Low cycle fatigue behaviour of a low interstitial Ni-base...

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Acta Materialia 57 (2009) 3450–3459

Low cycle fatigue behaviour of a low interstitial Ni-base superalloy

K. Gopinath a,b,*, A.K. Gogia a, S.V. Kamat a, R. Balamuralikrishnan a, U. Ramamurty b

a Defence Metallurgical Research Laboratory, DRDO, Hyderabad 500 058, Indiab Department of Materials Engineering, Indian Institute of Science, Bangalore 560 012, India

Received 9 December 2008; received in revised form 26 March 2009; accepted 28 March 2009Available online 7 May 2009

Abstract

The low cycle fatigue behaviour of precipitation strengthened nickel-base superalloy 720Li containing a low concentration of inter-stitial carbon and boron was studied at 25, 400 and 650 �C. Cyclic stress response at all temperatures was stable under fully reversedconstant total strain amplitude (De/2) when De/2 6 0.6%. At De/2 > 0.6%, cyclic hardening was followed by softening, until fractureat 25 and 650 �C. At 400 �C, however, cyclic stress plateaued after initial hardening. Dislocation–dislocation interactions and precipitateshearing were the micromechanisms responsible for the cyclic hardening and softening, respectively. The number of reversals to failurevs. plastic strain amplitude plot exhibits a bilinear Coffin–Manson relation. Transmission electron microscopy substructures revealedthat planar slip was the major deformation mode under the conditions examined. However, differences in its distribution were observedto be the cause for the bilinearity in fatigue lives. The presence of fine deformation twins at low De/2 at 650 �C suggests the role of twin-ning in homogenization of cyclic deformation.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: 720Li; LCF; Planar slip

1. Introduction

Precipitation strengthened superalloys retain excellentmonotonic strength levels up to high fractions of theirmelting point (�0.6Tm). However, in gas turbine applica-tions, the lives of nickel-base superalloys are often limitedby fatigue considerations. It is noted that the endurancelimit of nickel-base superalloys (for 107 cycles) normalizedwith yield strength is only �0.25 compared with 0.5–1.0observed for many engineering alloys [1]. While degrada-tion of the strengthening precipitates through repeatedshearing during deformation has been a contributing factor[2], more often, cracks that are initiated from the structuralheterogeneities present in the microstructure are consideredresponsible for the relatively poor performance of manysuperalloys under conditions of cyclic deformation [3].Such microstructural defects include microporosity in cast

1359-6454/$36.00 � 2009 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2009.03.046

* Corresponding author. Address: Defence Metallurgical ResearchLaboratory, DRDO, Hyderabad 500 058, India.

E-mail address: [email protected] (K. Gopinath).

alloys, pores or prior particle boundaries in powder metal-lurgy products and brittle phases such as carbides andnitrides in cast and wrought products [4]. Though initialefforts to achieve better fatigue properties in wroughtnickel-base disc alloys by reducing carbon content andimproving the distribution of carbides did meet with somesuccess, carbon levels could not be reduced below a certainlevel without adverse effects on stress–rupture properties[5]. Subsequently, this problem could be overcome in mod-ern superalloys by controlling the content of boron and zir-conium [6] as microalloying elements. Alloy 720Li (‘lowinterstitial’ abbreviated as ‘Li’), developed for high-integ-rity rotating components such as turbine discs in gas tur-bines [7], belongs to this class of advanced wroughtsuperalloys containing reduced levels of carbon and boroncompared with other contemporary wrought disc alloys [8].In view of the high speed of rotation and associated centrif-ugal stresses experienced by discs, disc alloys should pos-sess good elevated temperature tensile strength. Stressesthat are associated with changes in the rotational speedof the turbine and the stresses that are of thermal origin

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K. Gopinath et al. / Acta Materialia 57 (2009) 3450–3459 3451

are cyclic in nature, emphasizing the need for disc alloys topossess good low cycle fatigue (LCF) resistance at operat-ing temperatures.

The chemistry of alloy 720Li evolved from that of alloy720, originally developed for blade application in land-based gas turbines [9]. However, to achieve the strengthand LCF resistance desired for disc application, the process-ing and heat treatment of alloy 720Li is different from thatadopted for alloy 720 for creep-resistant blade application[9,10]. The resultant changes in grain size, precipitate shape,size and distribution are closely linked to the macroscopicproperties of the alloy. The tensile behaviour of alloy720Li was addressed by the authors recently [11]. Thosestudies revealed that 0.2% yield strength and ultimate tensilestrength of alloy 720Li are insensitive to temperatures up to600 and 500 �C, respectively. Above 650 �C, both yield andultimate strength decreased markedly with temperature.The alloy exhibited positive strain rate sensitivity at 25and 750 �C. At 400 �C, the alloy exhibited serrated flowand negative strain rate sensitivity due to dynamic strainageing (DSA) [12]. Heterogeneous planar slip, involvingshearing of precipitates, was found to be the predominantdeformation mechanism at all strain rates at 25 �C whereas,at 750 �C and low strain rate (10�5 s�1), homogeneous slipwas observed. While some aspects of fatigue crack growthrates, which are the essential inputs for modern design pro-cesses, have been studied in alloy 720Li [13,14], the results ofstudies on LCF behaviour [13,15,16] and deformation pro-cesses are relatively limited in open literature. Studies onwrought 720Li with low interstitial content thus providesan ideal opportunity to assess the substructures and damagemechanisms associated with cyclic deformation, develop-ment of which are often curtailed in many superalloys byfailures initiated from defects discussed earlier. Therefore,this paper examines the LCF behaviour of alloy 720Li at25, 400 and 650 �C. While 400 and 650 �C are typical tem-peratures at the bore/web and rim of turbine discs, respec-tively, studies at 25 �C formed the reference to assess theinfluence of thermal effects at elevated temperatures. Inaddition, tests at 400 �C would also help to understand man-ifestations of DSA on LCF behaviour. The cyclic stressresponse, plastic strain dependence of fatigue lives andmicromechanisms involved are discussed in light of thedeformation substructures obtained.

2. Materials and experiments

Alloy 720Li used in this study (composition in Table 1)was supplied by M/s Aubert and Duval (Les Ancizes,France) in the form of forged blocks of dimensions150 � 150 � 100 mm, processed and heat-treated to

Table 1Chemical composition of alloy 720Li.

Element Cr Co Ti Mo Al

Wt.% 16.03 14.42 4.93 2.97 2.4

achieve fine grains and high strength meant for turbine discapplication. The processing involved triple melting sequen-tially through vacuum induction melting, electro slag refin-ing and vacuum arc refining and subsequent processingthrough the conventional ingot metallurgy route. The alloywas solutionized and aged. Solutionizing was carried out ata sub-solvus temperature of 1090 �C for 4 h, followed byoil quenching. First-stage ageing was carried out at650 �C for 24 h and air cooled. The second-stage ageingwas at 760 �C for 16 h followed by air cooling.

The specimens for LCF testing were machined fromcylindrical blanks that were electric discharge machinedfrom the fully heat-treated forged blocks. All tests wereconducted on threaded round specimens with a gaugediameter of 4.5 mm and a gauge length of 14 mm. Circum-ferential machining marks in the gauge length portion ofthe specimens were fully removed by polishing parallel tothe stress axis to an average surface roughness(Ra) < 0.2 lm. Tests were carried out at ambient tempera-ture (25 �C), 400 and 650 �C in fully reversed mode(R = �1) at total strain amplitudes De/2 of 0.4%, 0.5%,0.6%, 0.7%, 0.8%, 0.9% and 1.0%, using extensometers ofgauge lengths 10 and 12 mm. While tests at De/2 of 1.2%were also carried out at 25 and 400 �C, they were not car-ried out at 650 �C, as only extremely low lives (<50 cycles)would have resulted. Tests at different strain amplitudeswere carried out at appropriate frequencies which wouldresult in a constant strain rate of 5 � 10�3 s�1. Load andextension data were recorded using an on-line computer.Tests were conducted as per ASTM E 606-04 [17] in aclosed-loop servo-hydraulic machine (M1000RK, Dartec,UK) of 100-kN load capacity equipped with a resistance-heated three-zone split furnace. All tests were started aftersoaking for 30 min at the test temperature. Temperature T

during the test was monitored using a K-type thermocou-ple and recorded with a strip chart recorder.

Samples for optical metallography were polished usingstandard metallographic procedures and electrolyticallyetched with 10% orthophosphoric acid in water. Scanningelectron microscopy (SEM) studies were carried out in aFEI Quanta 400 microscope equipped with an energy dis-persive X-ray spectroscopy (EDS) facility. Deformationstructures were studied through transmission electronmicroscopy (TEM) of thin foils sliced from the gauge sec-tions of fatigue tested specimens at 45� to the loading axisusing a low speed saw. Discs 3 mm in diameter werepunched out after the slices were mechanically thinneddown to 100 lm. They were twin-jet electropolished to per-foration in a solution of mixed acids (78 vol.% methanol,10 vol.% lactic acid, 7 vol.% sulphuric acid, 3 vol.% nitricacid and 2 vol.% hydrofluoric acid) at �30 �C and 12 V

W C B Zr Ni

9 1.21 0.011 0.015 0.030 Balance

Fig. 1. Microstructure of alloy 720Li in sub-solvus solution treated andtwo-stage aged condition. Optical micrographs showing (a) primary c0

precipitates (bright) at grain boundaries (electrolytic orthophosphoricetch) and (b) presence of annealing twins (modified Kalling’s etch). (c)Dark-field TEM image showing secondary and tertiary c0 precipitates.

3452 K. Gopinath et al. / Acta Materialia 57 (2009) 3450–3459

DC. The foils were examined using a FEI Tecnai 20Tmicroscope operating at 200 kV.

3. Results

3.1. Microstructure

The average grain size of the material used for the study isestimated to be ASTM 10 (�11 lm) with occasional grainsup to ASTM 8 (�22 lm). Optical metallography of electro-lytically etched specimen reveals primary c0 precipitates ingood contrast as bright phase (size 1–10 lm and volume frac-tion �18%) at the grain boundaries of the darker c matrix(Fig. 1a). Wide annealing twins were confirmed to be inher-ent in the standard microstructure of the material in as-received condition by examination of multiple specimensfrom locations far apart (Fig. 1b). The presence of irregularlyshaped particles was also noted, and they were confirmed byEDS analysis to be mainly TiC and Ti(C, N) particles. TEMexamination of the as-received material reveals (Fig. 1c)intragranular secondary c0 precipitates (�120 nm) and finetertiary c0 precipitates (�50 nm).

3.2. Cyclic stress response

Cyclic stress response of alloy 720Li (Fig. 2) is influencedby both De/2 and T. The alloy did not show appreciable hard-ening or softening at any T for De/2 6 0.6%. However, at 25and 650 �C and at De/2 > 0.6%, the alloy exhibited initialhardening, which was followed by softening (Fig. 2). Thecyclic stress responses at 400 �C were distinct because ofthe absence of softening following the initial hardening atDe/2 between 0.6% and 1%. However, mild softening couldbe noticed at 400 �C at De/2 = 1.2% (Fig. 2d). The degreeof hardening D is computed using the following equation:

D ð%Þ ¼ ðDr=2Þmax � ðDr=2Þ1ÞðDr=2Þ1

� �� 100 ð1Þ

where (Dr/2)max is the maximum stress amplitude exhibitedduring cycling, and (Dr/2)1 is the stress amplitude from thefirst cycle. Fig. 3 shows the variation in D as a function ofDe/2. It is seen that D depends on both T and De/2. It is evi-dent that, at all T, D is low at De/2 < 0.6% and startsincreasing when De/2 P 0.6%, reaching a maxima at De/2 = 1.0%, as clearly evident at 25 and 400 �C (Fig. 3). At650 �C, D is comparatively lower than at 25 and 400 �C,and the drop in D after maxima could not be observed(Fig. 3) as tests were not conducted at De/2 of 1.2% at650 �C. Though the trend in the variation in D was similarat the three T, the magnitude of D at 400 �C was clearly thehighest among the three T studied.

3.3. Cyclic stress–strain response

From the tests conducted to determine fatigue lives as afunction of De/2, the cyclic stress–strain (CSS) responseswere extracted (i.e. companion specimen method [18] by

plotting the Dr/2 values recorded at mid-life against theimposed De/2, as done conventionally (Fig. 4)). These werecompared with the monotonic stress–strain responsesobtained from the first cycle. Fig. 4 indicates that alloy720Li exhibits mixed behaviour at 25 and 650 �C, i.e. cyclicsoftening at low De/2 and cyclic hardening at high De/2. At400 �C, while no cyclic softening could be observed at lowDe/2, cyclic hardening was clearly evident at De/2 > 0.6%.

Monotonic and CSS responses were analysed using thefollowing power law relationships, respectively, [19]:

Fig. 2. Cyclic stress response of alloy 720Li.

Fig. 3. Degree of hardening D computed from cyclic stress response data.

Fig. 4. Monotonic and CSS curves: (a) 25 �C; (b) 400 �C; (c) 650 �C.Stresses (r) are from the tensile segments of the first cycle, and stressamplitudes (Dr/2) are from mid-life hysteresis loops.


r ¼ Kenp ð2aÞ

Dr2¼ K 0

Dep

2

� �n0

ð2bÞ

where n and n0 are monotonic and cyclic strain hardeningexponents, K and K0 are monotonic and cyclic strengthcoefficients, r is the true stress, and Dr/2 and Dep/2 arestress amplitude and plastic strain amplitude at half life.The values of the exponents and coefficients determinedas in Eqs. (2a) and (2b) at 25, 400 and 650 �C are tabulatedin Table 2. Increased work hardening in fatigue is evidentin the values of n0 which are larger than n.

3.4. Low cycle fatigue lives

The fracture life Nf, obtained as a function of De/2 at thethree T studied, is plotted in Fig. 5. Dest-fit curves in Fig. 5show that fatigue lives at 650 �C are clearly lower thanthose at 25 and 400 �C at all De/2. Nf at 400 �C is only

Table 2Monotonic and cyclic strength coefficients and strain hardening exponentsobtained by fitting Eq. (2) to the respective stress–strain curves.

Temp. (�C) Monotonic Cyclic

K (MPa) n K0 (MPa) n0

25 1412 0.031 1306 0.072400 1410 0.042 1325 0.062650 1285 0.037 1272 0.118

Fig. 5. Dependence of fatigue life (Nf) on imposed strain amplitude (De/2).


marginally lower than that at 25 �C for De/2 P 0.6%.However, at the low De/2 levels of 0.4% and 0.5%, Nf

at 400 �C are higher than that at 25 �C, suggesting acrossover.

In view of the dominating influence of Dep/2 on fatiguedamage, fatigue lives were analysed using the Coffin–Man-son (CM) equation [19] below:

Dep

2¼ e0f ð2Nf Þc ð3Þ

where e0f is the fatigue ductility coefficient, c is the fatigueductility exponent, and 2Nf is the number of reversals tofailure. According to Eq. (3), a plot of Dep/2 vs. 2Nf onlog–log axes would result in a straight line with slope c.At all three T studied, 2Nf data plotted against Dep/2 atmid-life (Fig. 6) show that the data are better fitted bytwo straight-line segments, one of shallow slope at highDep/2 and another of steep slope at low Dep/2. Such bilin-earity in slope is of practical significance, as extrapolationbased on any one segment alone would lead to overestima-

Fig. 6. CM plots of plastic strain amplitude Dep/2 vs. number of strainreversals 2Nf, exhibiting bilinearity in slope (Eq. (3)).

tion of fatigue lives. Values of e0f and c pertaining to thelow Dep/2 and high Dep/2 regimes are tabulated in Table 3.

3.5. Deformation substructures

TEM investigations were conducted to assess whetherthere were any changes in deformation substructures thatcould explain the observed bilinearity in the CM plots(Fig. 6). Deformation substructures pertaining to speci-mens tested to fracture at 25 and 650 �C at both low Dep/2 (steep slope regime) and high Dep/2 (shallow sloperegime) were studied. In specimens subjected to De/2 = 0.4% (mid-life Dep/2 = 0.006%) at 25 �C, grain interiorswere relatively free of dislocations. Minimal dislocationactivity could be observed at grain boundaries which weredecorated with primary c0 precipitates (Fig. 7a). Isolatedslip bands with relatively low dislocation density could alsobe observed occasionally. In specimens subjected to De/2 = 1.0% (mid-life Dep/2 = 0.4%) at 25 �C, most grainsshowed the presence of many intense slip bands alongone family of {1 1 1} planes. High-angle tilting experimentssuggested that dislocation activity between the slip bandswas minimal (Fig. 7b). Activation of secondary slip alonganother set of {1 1 1} planes could also be noticed, but onlyin a few grains. Deformation is thus heterogeneous andconfined mainly to slip bands, but numerous bands spreadthroughout the grain were noted.

Substructure of specimens subjected to De/2 = 0.4%(mid-life Dep/2 = 0.01%) at 650 �C revealed presence offew slip bands and homogeneously distributed dislocationsof low density in the inter-band regions (Fig. 8a). In addi-tion, fine twins (50–100 nm thick) were observed in numer-ous grains, many of them with multiple parallel twins(Fig. 8b). While coarse annealing twins are regular micro-structural features of the as-received material (Fig. 1b),the fine twins were observed only rarely in as-receivedmaterial. The high density of fine twins in fatigued speci-mens suggests that they are deformation induced. Speci-mens subjected to De/2 = 1.0% (mid-life Dep/2 = 0.38%)at 650 �C contained intense slip bands (Fig. 9a–d) as wellas significant dislocation activity between the bands.

4. Discussion

It is known that low stacking fault energy (SFE) materi-als exhibit high degree of planarity in slip, especially at

Table 3Values of fatigue ductility coefficient e0f and fatigue ductility exponent c

pertaining to low and high Dep/2 regimes at three test temperatures.

T (�C) Regime e0f (%) c

25 Low Dep/2 4.8 � 105 �1.52High Dep/2 10.7 �0.45

400 Low Dep/2 4 � 105 �1.71High Dep/2 2.3 �0.28

650 Low Dep/2 2. � 103 �1.26High Dep/2 3.7 �0.44

Fig. 7. TEM micrographs of deformation substructure in specimensfatigued at 25 �C. (a) De/2 = 0.4% (Dep/2 = 0.006%), Nf = 60047 cycles;general bright-field image of low dislocation activity limited to grainboundaries and primary c0 boundaries. (b) De/2 = 1.0% (Dep/2 = 0.4%),Nf = 467 cycles; bright-field image showing intense slip bands along (1 1 1)planes.

Fig. 8. General bright-field TEM micrographs of deformation substruc-ture in specimens fatigued at 650 �C, De/2 = 0.4% (Dep/2 = 0.01%),Nf = 11132 cycles: (a) slip bands and low dislocation activity in theinter-band regions; (b) multiple twins within a grain. SAD pattern in theinset taken in h1 1 0i zone axis orientation clearly shows twin diffractionspots.


T < 0.4Tm [20]. From the alloy chemistry (Table 1), it isreasonable to assume low SFE for alloy 720Li as it con-tains Cr, Co and W, all of which are known to reducethe SFE of Ni significantly [21–23]. As a consequence, pla-narity of slip is expected in alloy 720Li. The presence ofcoherent precipitates, which in the case of alloy 720Li con-stitute �30 vol.%, also promotes slip planarity [24]. In linewith expectation, a high degree of planarity in slip anddeformation localized on slip bands was observed in alloy720Li at 25 �C in unidirectional tensile deformation [11].Homogeneous slip was prevalent only at high T (750 �C)and low strain rates (10�5 s�1). It is established that paral-lelism between unidirectional and cyclic deformation existin polycrystalline face centred cubic (fcc) metals [25] andalloys containing particles penetrable by dislocations [26].In particular, hardening behaviour at low and high plasticstrain amplitudes in fatigue are considered distinct [25] andcorrelated with different stages of unidirectional deforma-tion. The results obtained and deformation substructuresobserved as part of this study are interpreted in light ofthese findings, which suggest that cyclic straining willaccentuate the dislocation processes prevalent in unidirec-

tional deformation. As cyclic deformation behaviour andthe substructures developed dictate the fatigue lives, theseaspects are considered prior to the discussion on fatiguelives.

4.1. Characteristics of cyclic deformation

4.1.1. Cyclic stress response

At low De/2 levels at all T studied, the stress response isstable from the initial cycles (Fig. 2a). The magnitude ofDep/2 at mid-life developed for the imposed De/2 is shownin Fig. 10. At De/2 = 0.6%, Dep/2 values are about an orderof magnitude smaller than the imposed De/2 at all T. Thisdifference increases with decreasing De/2. Deformation infatigue at low Dep/2 is considered similar to stage I harden-ing in unidirectional deformation [25]. At low De/2, the dis-location density and average travel distances ofdislocations are small, hence only a few complex disloca-tion interactions or tangles occur. These were confirmedduring TEM studies of specimens cycled at De/2 = 0.4%,both at 25 �C and 650 �C, wherein only minimal disloca-tion activity was observed within the grains. However,

Fig. 9. General bright-field TEM micrographs of deformation substructure in specimens fatigued at 650 �C and De/2 = 1.0% (Dep/2 = 0.38%), Nf = 196cycles: (a) intense slip bands and homogeneous dislocation activity in the inter-band regions; (b) intense slip bands spread through out the grain; (c)isolated slip band and homogeneous dislocation distribution; (d) homogeneous dislocation activity, same grain as (c), at higher magnification.

Fig. 10. Dep/2 developed at mid-life as a function of imposed De/2 at threetemperatures studied.


some dislocation activity could be observed near grainboundaries (Figs. 7a and 8a) where relatively more strainis accommodated so as to maintain the continuity betweengrains. Such low levels of Dep/2 and resultant minimal dis-

location activity are considered the reason behind the rela-tively low D values at low De/2 levels (Fig. 3).

At De/2 > 0.6%, the alloy exhibited initial hardening,which was followed by softening at 25 and 650 �C(Fig. 2b–d). It is known that crack initiation can lead tosoftening, but in such a scenario the peak stress in tensionwould show a much higher reduction compared with thepeak stress in compression. However, in the present case,the peak stress in tension and compression were found toreduce equally, confirming that softening is intrinsic tothe material and not due to crack initiation. At a givenDe/2, the increase in Dr/2 during the initial cycles is gener-ally attributed increasing dislocation-dislocation interac-tions and associated increase in dislocation density. Inthe case of planar slip materials, these interactions leadto dislocation tangling within the slip bands [25] makingthem denser (Figs. 7b and 9b). The reduction in Dr/2beyond the maxima in Fig. 2b–d suggests that some soften-ing process attains significance over the hardening phe-nomenon prominent until then. Shearing of precipitates[26–28] and dissolution of sheared precipitates leading tothe formation of precipitate-free deformation bands[29,30] have been observed during the softening phase insuperalloys. Shearing of secondary c0 precipitates could

Fig. 11. Dark-field TEM micrograph showing shearing of secondary c0

precipitates: 25 �C, De/2 = 1.0% (Dep/2 = 0.4%), Nf = 467 cycles.


be observed (Fig. 11) in TEM foils from fractured speci-mens subjected to De/2 = 1% both at 25 and 650 �C. It istherefore apparent that, in alloy 720Li, precipitate shearingand work hardening of the c matrix happen simulta-neously, with the more dominating effect dictating the netcyclic stress response. Work hardening of the c matrix willtend to dominate during the initial cycles. The extent towhich c matrix hardens (i.e. (Dr/2)max � (Dr/2)1), beforethe softening due to precipitate shearing starts to dominate,depends on the stress levels generated to achieve theimposed De/2. Therefore, it is interpreted that the stress lev-els pertaining to De/2 values beyond the maxima in Fig. 3are high enough right from the first cycle, leading to earlyprecipitate shearing and lesser extent of hardening, beforesoftening effects dominate. From the application point ofview, cyclic softening could redistribute the stresses, allevi-ating stress-concentration effects at geometric features incomponents such as fir-tree roots and bolt holes that aretypical in turbine discs [31].

The tendency to soften after initial hardening, observedfrom intermediate levels of De/2 at 25 and 650 �C (Fig. 2band 2c), is observed only at the highest applied De/2 of1.2% at 400 �C (Fig. 2d). While the softening at high De/2 could be due to reasons discussed above, the absence ofsoftening at De/2 6 1% suggests the prevalence of an addi-tional strengthening mechanism not operative at 25 and650 �C, which offsets softening due to precipitate shearing.It is considered that this is a manifestation of DSA prevail-ing at this T [12]. During DSA, diffusing solute atoms lockmobile dislocations, resulting in an increase in flow stress[32]. As diffusion of solute atoms is critical for DSA,increased vacancy concentration during cyclic deformation[33] compared with that in monotonic deformation tends tomagnify the effects of DSA. In addition, unloading seg-ments during fatigue cycling reduces the driving force fordislocation mobility and provides additional time for soluteatoms to diffuse and lock mobile dislocations morestrongly. Hence, both the absence of softening (Fig. 2b)and the higher degree of cyclic hardening at 400 �C than

at 25 and 650 �C (Fig. 3) are considered manifestationsof DSA. Detailed studies on the effect of DSA on theLCF behaviour of this alloy are currently under way andare therefore not reported here.

4.1.2. Monotonic and CSS behaviourCSS curves reflect the influence of cycling on the stress–

strain relationship in materials. In many materials, micro-structural changes that happen during fatigue tend to stabi-lize after an initial shakedown period, during which thematerial may harden or soften [34]. Evolution of stable sub-structures in pure metals and simpler alloys is often com-pleted by mid-life and is, generally, the unstatedpresumption when Dr/2 values at mid-life are used for plot-ting the CSS curves. In alloy 720Li, the position of the CSScurves can vary, depending on the fraction of life chosenfor plotting them, as continuous softening from a peak stressis observed at all intermediate and high De/2 at 25 and650 �C. However, following the convention, CSS curvesplotted using Dr/2 values at mid-life (Fig. 4) and the coeffi-cients and exponents determined from them (Table 2) canbe useful for comparison between similar alloys.

As expected from the monotonic and CSS curves(Fig. 4), the values of cyclic strain hardening exponents n0

at the T examined are higher than the correspondingmonotonic strain hardening exponents in tension n (Table2), signifying an increased work hardening rate in fatigue.It is pertinent to note that the observed values of n (0.03–0.04) are very low compared with those generally reportedfor metals (0.1–0.5). Monotonic tensile studies on alloy720Li [11] had revealed that the plastic regime on thelog–log plot of r vs. ep is better represented by two linearstages, one with a shallow slope n1, at lower strains, andthe other with a steeper slope n2, at higher strains, ratherthan by a single linear fit of slope n expected as in Eq.(2a). The magnitudes of n2 were found to conform to typ-ical n values for metals, and the deviations at low strainsresulting in lower slope n1 are considered to correspondto a regime of planar glide typical of low SFE materials[35]. Thus, only low n values (Table 2) would result atthe low plastic strains used for determining them.

4.2. Fatigue lives and deformation substructures

From the plot of De/2 vs. Nf (Fig. 5), it is evident thatboth De/2 and T influence LCF lives. Lives at 650 �C aresignificantly lower than those at 25 and 400 �C. Lives at400 �C are only marginally lower than that at 25 �C, exceptat De/2 = 0.4% and 0.5%, where the lives are higher thanthat at 25 �C. The observed higher lives at 400 �C than at25 �C at low De/2 levels can be comprehended if lives areassessed as a function of Dep/2. As Dep/2 = De/2 � [(Dr/2)/E], Dep/2 developed for any imposed De/2 depends onthe flow stress r required to attain De/2 and Young’s mod-ulus E. E of alloy 720Li gradually reduces from 227 GPa at25 �C to 193 GPa at 400 �C [10]. The monotonic tensileyield strength and ultimate strength of alloy 720Li remain


unaffected by T until �500 �C [11], typical of manywrought nickel-base superalloys strengthened by c0 precip-itates of 40–50% volume fraction [36]. Therefore, in alloy720Li, as T increase between 25 and 400 �C, E decreases,and r remains unaffected, and therefore Dep/2 at 400 �Cis lower than that at 25 �C, as indicated in Fig. 10. FromFig. 10, it can also be noted that the magnitude of Dep/2developed at low De/2 (0.4% and 0.5%) at 400 �C is lowerthan that at 25 �C by at least an order, and is consideredresponsible for higher LCF lives at 400 �C than at 25 �C.This is also reflected in plots of Dep/2 vs. 2Nf (Fig. 6).

Lower Dep/2 at 400 �C compared with 25 �C should haveresulted in comparatively higher LCF lives at all De/2.

However, the marginally lower lives at 400 �C suggest theinfluence of T and other associated effects on cyclic defor-mation mechanisms. Although the monotonic tensilestrength is essentially unaffected by an increase in T from25 to 400 �C [11], this increase could introduce an elementof homogeneity to slip and increase the irreversibility ofslip compared with 25 �C. As DSA is known to disperseplanar slip [20], it would increase fatigue lives at 400 �C,as monotonic tensile deformation studies had indicated aprevalence of DSA over a T range of 250–475 �C [12]. Thisessentially means that the LCF lives observed at 400 �C arethe net result of the effects considered above rather than afunction of Dep/2 alone. Homogeneity of slip increases withincreased thermal activation at elevated T, resulting inlower lives at 650 �C compared with 25 �C (Fig. 5) eventhough Dep/2 levels are not significantly different (Fig. 10).

Considering LCF principally as a crack propagationprocess and the representation of Dep/2 vs. 2Nf as that per-taining to crack propagation, it was suggested that bilinear-ity in the CM plot (similar to that at Fig. 6) is due to achange in fracture mode [37]. If changes in fracture modeand consequent reduction in life were arising out of envi-ronmental effects such as oxidation, it should have beenobvious only in the high life regime at elevated tempera-tures. However, in the case of alloy 720Li, the presenceof bilinearity even at 25 �C (Fig. 6), where the role of envi-ronment is expected to be least, suggests that the environ-mental effect may not be the reason for the bilinearity ofCM plots. Change in deformation process with Dep/2 isanother possibility which can be manifest in bilinearity[38]. A third possibility is associated with the changes inhomogeneity of deformation [39].

Results of the TEM studies carried out to understand thedeformation processes are in Figs. 7–9. The substructures inFigs. 7a and 8a and b represent a steep slope regime (lowDep/2), and those in Figs. 7b and 9 are typical of a shallowslope regime (high Dep/2) in the CM plot (Fig. 6). While pla-nar slip appears to be the dominant deformation mecha-nism under the different conditions examined, there arenoticeable differences in the homogeneity of slip band distri-bution depending on Dep/2 and T. At 25 �C and low Dep/2,slip is confined primarily to a few isolated slip bands and istherefore highly heterogeneous. This favours strain locali-zation, thereby leading to early crack initiation and easy

crack propagation resulting in lower fatigue lives. However,at high Dep/2, the imposed higher strain has resulted in anincreased density of slip bands which, in a sense, representsa move towards more homogeneous deformation. At650 �C and low Dep/2, slip bands are isolated, and a few dis-locations homogeneously distributed in the inter-bandregions are also seen (Fig. 8a). In addition, fine deformationtwins can also be observed in many of the grains (Fig. 8b).At 650 �C and high Dep/2, the deformation is much morehomogeneous. Not only are many slip bands present, butthe density of dislocations in the inter-band regions is alsoconsiderably higher (Fig. 9a). Occasional grains with denseslip bands spread homogeneously throughout the grains arealso observed (Fig. 9b). Smaller grains have only relativelyfew slip bands but exhibit extensive homogeneous deforma-tion in the inter-band regions (Fig. 9c and d). The role ofthermal activation in increasing the homogeneity of slip isclearly evident from the deformation substructures at650 �C (Fig. 9a–d).

At 25 �C, the bilinearity in slope of the CM plot (Fig. 6)can be attributed clearly to the difference in the density ofslip bands observed within the grains. At lower Dep/2(fewer slip bands), the stronger localization of deformationis responsible for the lower lives observed in the steep sloperegime (Fig. 6). At 650 �C, the magnitudes of the slopes inboth the low and high Dep/2 regimes are similar to those at25 �C (Table 3). As the slope in low Dep/2 is not affected bytwinning, it is clear that bilinearity at 650 �C is also due tothe differences in homogeneity of slip between low and highDep/2. Instances of deformation twins being confined tostrengthening precipitates [40] and spreading through thematrix and precipitates [41] during tensile deformationhave been reported earlier. Twinning associated with fati-gue deformation has also been reported for Ni–Fe-basesuperalloy 718 [42–44]. Deformation-induced microtwin-ning during creep has recently been reported [45,46] onnickel-base superalloys of very similar microstructures.While these observations confirm the possibility of twin-ning in fcc structures of low SFE where cross-slip is diffi-cult, detailed studies are required to understand theirorigin and precise role in fatigue at high T and low Dep/2, as observed in this study. However, the presence of twinsin specimens fatigued at elevated temperatures where theoverall deformation is homogeneous suggests a possiblerole of twinning in homogenizing the deformation, albeitindirectly. As twinning can place new slip systems in an ori-entation favourable for subsequent slip, it is possible thatthe formation of more microtwins aids homogenizationof deformation. Newer microtwins have been shown toemerge to accommodate cyclic deformation in superalloy718 [44]. In comparison with alloys where twinning isrestricted to precipitates [40], homogeneity of deformationin alloy 720Li is expected to be higher, as the twins havepropagated through both the c matrix and c0 precipitates,as confirmed by the presence of twin spots in the diffractionpattern for both regular and superlattice reflections(Fig. 8b).


5. Conclusions

LCF studies on precipitation strengthened nickel-basesuperalloy 720Li at 25, 400 and 650 �C under fully reversed(R = �1) constant strain amplitude (De/2) cycling revealedthe following:

(a) At all three temperatures studied, at low De/2(60.6%), cyclic stress response was stable from theinitial cycles. This is linked to the low level of disloca-tion activity that prevailed.

(b) At 25 and 650 �C, at intermediate and higher De/2(P0.7%), the alloy initially exhibited cyclic harden-ing, which was followed by softening until fracture.Interactions of dislocations with other dislocationsand precipitates are considered to account for the ini-tial hardening, and shearing of precipitates is consid-ered key to subsequent softening.

(c) Fatigue lives exhibit a bilinear relationship with Dep/2at all T studied. The deformation substructuresrevealed planar slip as the dominant deformationmode. However, the extent of localization of defor-mation which depends on the imposed De/2 is consid-ered responsible for bilinearity in Dep/2 vs. 2Nf plots.

Acknowledgements

The authors would like to thank Defence Research andDevelopment Organization for permitting and supportingthe above research. Motivation, encouragement and the al-loy for the study provided by the Gas Turbine ResearchEstablishment, Bangalore, is gratefully acknowledged. Spe-cial thanks are due to Mr. D.G. Sathe and Mr. ChennaReddy of AMTL for their help in specimen preparationand testing.

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