International Journal of Fatigue 63 (2014) 137–144

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International Journal of Fatigue

journal homepage: www.elsevier .com/locate / i j fa t igue

Comparison of low-cycle fatigue behaviors between two nickel-basedsingle-crystal superalloys

http://dx.doi.org/10.1016/j.ijfatigue.2014.01.0180142-1123/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding authors. Tel.: +86 24 83978226 (P. Li). Tel.: +86 24 23971757(T. Jin).

E-mail addresses: pli@imr.ac.cn (P. Li), tjin@imr.ac.cn (T. Jin).

P. Li a,⇑, Q.Q. Li a,b, T. Jin b,⇑, Y.Z. Zhou b, J.G. Li b, X.F. Sun b, Z.F. Zhang a

a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, 110016 Shenyang, PR Chinab Superalloy Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, 110016 Shenyang, PR China

a r t i c l e i n f o a b s t r a c t

Article history:Received 4 November 2013Received in revised form 12 January 2014Accepted 14 January 2014Available online 24 January 2014

Keywords:Nickel-based superalloyHigh temperature low-cycle fatigue (LCF)Slip characterCracking mode

The deformation mechanisms of nickel-base single-crystal superalloys containing 3 wt.% Re (3Re) andwithout Re (0Re) were systematically investigated in the temperature range from room temperature to900 �C. At all testing temperatures, the alloy 3Re showed higher stress amplitudes and longer fatiguelives. With increasing temperature, slip mode gradually changes from planar to wavy slip. Dislocationclimb and cross-slip, even by-passing became dominant deformation mechanisms of the two alloysand the addition of Re further reinforced this trend. In this process, the deformation homogenizationbecame dominant, which resulted in the transition of cracking mode from shear fracture to normal frac-ture. At low temperature the specimen eventually cracks along the slip bands by shear, but at high tem-perature the specimen cracks by normal and shows mode I cracking.

� 2014 Elsevier Ltd. All rights reserved.

0

1. Introduction

Nickel-based single-crystal (SC) superalloys are one of the mostwidely used high-temperature structural materials because theyhave an excellent balance of physical and mechanical properties,including good fatigue and fracture resistance and outstanding oxi-dation, corrosion and creep resistance at elevated temperatures[1,2]. The single-crystal (SC) superalloys are usually classified intofirst-, second- and third-generations. The second and third genera-tions contain about 3 and 6 wt.% of Re [3–5], respectively. Re is avery expensive element, the addition of which can lead to animprovement in the creep strength. It is suggested that the en-hanced resistance to creep comes from the promotion of raftingby Re, which partitions into the c phase and brings the lattice mis-fit to be more negative [6–8]. Furthermore, Janotti et al. [9] and Fuet al. [10] considered that the Re/vacancy exchanging energy in theNi-rich matrix was remarkably high, which may lead to retardmost of diffusion-driven processes in these materials.

For low-cycle fatigue (LCF) behaviors of nickel-base superalloys,the temperature always plays a more important role. Previousstudies have shown [11–14] that planar slip in superalloy is dom-inant at low temperature. The motion of dislocations is restrictedto thin slip bands parallel to {111} slip planes [15]. The appearanceof slip bands is responsible for the initial hardening and shearing of

c precipitates. This deformation mode led to softening [16]. On theother hand, slip band density increases with increasing total strainamplitude. Petrenec et al. [17] found that highly inhomogeneousplanar dislocation arrangements in the form of bands were ob-served after cyclic straining at all temperatures from room temper-ature up to 800 �C. However, Pineau and Antolovich [18] also foundthat as the temperature was increased, the dislocation structurebecame much homogenous and the dislocation density withinthe bands was much lower than those at low temperatures. Inaddition, as temperature increased, both average slip line spacingand slip offset measured by atomic force microscopy (AFM) de-creased [19]. Above 800 �C, the slip bands were not as frequentlyobserved and homogenous non-planar slip became dominant[20–22]. Chu et al. [23] compared the fatigue behaviors of nickel-base superalloy from 700 �C to 900 �C and found that at 700 �Cthe deformation was concentrated in slip bands running across c0

particles but at 900 �C no slip bands appeared anywhere. It is usu-ally considered that the LCF deformation is controlled by the Oro-wan mechanism and that the dislocation moves by cross-slip in thehorizontal c channel to form dislocation loops at 700 �C [24]. Fur-thermore, dislocation climb, slip and interaction between disloca-tions and carbides are the main LCF deformation mechanisms at900 �C [25]. It is likewise found that at high temperature interfacialnetworks of near-edge dislocations are the most common in the c0

precipitates with a few stacking faults [26,27]. Zhou et al. [28] fur-ther pointed out that stacking faults (SFs) also appeared in the ma-trix and more completed rafts and smaller interfacial dislocationspacing at 1100 �C were observed.

Table 1Nominal composition of the two nickel-based SC superalloys.

Specimen Element (wt.%)

Cr Co Mo W Ta Re Al Ti Ni

Alloy 0Re 6.0 10.0 2.0 6 8.0 – 4.0 2.0 BalAlloy 3Re 6.0 10.0 2.0 3 8.0 3 4.0 2.0 Bal

Table 2Fatigue conditions and data of two nickel-based SC superalloys at differenttemperatures.

Temperature (�C) Specimen cpl (%) ES (GPa) Cyclic No.

25 0Re 0.82 123 ± 1 10003Re 0.83 127 ± 1 2000

250 0Re 0.85 112 ± 2 900

500 0Re 0.78 108 ± 3 17003Re 0.89 112 ± 3 1800

750 0Re 0.79 101 ± 6 13703Re 0.77 105 ± 5 4500

900 0Re 0.80 95 ± 7 11003Re 0.84 101 ± 9 1900

138 P. Li et al. / International Journal of Fatigue 63 (2014) 137–144

It seems that when the temperature exceeds 800 �C, the strainlocalization phenomenon will disappear and slip band cannot existbecause of the homogenous deformation. But the actual mecha-nism is much more complicated. It is well known that althoughthe stacking fault energy (SFE) value of nickel itself is high, withthe addition of various elements, the SFE value of nickel-basesuperalloy becomes very low at room temperature [2,29], whichcan be used as a prerequisite for planar slip. On the other hand,the formation of stacking faults is probably due to the ordered c0

precipitates. Shearing of the c0 precipitates requires the passageof dislocation pairs to maintain the ordered condition of the parti-cle. Such dislocation arrangement can be regarded as two-dimensional structure since dislocation movement cannot escapefrom {111} plane. As the temperature increases, the SFE value willreturn to higher values [30]. By rights, the superalloy should exhi-bit its wavy slip feature and form more three dimensional disloca-tion arrangements just like copper, nickel and aluminum [31]. Butclearly, it is not entirely true because although at higher tempera-ture the dislocations may be easier to cross slip, the interactionbetween dislocations and c0 phase will change the original accu-mulation and evolution process of dislocations. Therefore moreprecisely, nickel-base superalloys experience a transition fromthe strain localization to the homogenous deformation withincreasing temperature. The question is how the slip mode changesand what this change means for the fatigue damage mechanisms ofsuperalloy.

It is well known that the development of a life prediction schemerequires a proper understanding of LCF damage mechanisms in thetemperature range in service. At room temperature, fatigue crackshave been found to initiate from slip bands [32–35], twin bound-aries [36], grain boundaries [37], carbides [38], or casting porosities[39]. On the other hand, no matter where the fatigue crack initiatesalong, at low and intermediate temperatures (T < 760 �C) a fatiguecrack propagates predominantly along {111} crystallographicplanes [40]. At elevated temperatures, plastic deformation is muchmore homogeneous and is not confined to slip bands due to ther-mally activated processes and fracture surfaces appear to be non-crystallographic [41], although fracture features are obscured byoxidation. Ott and Mughrabi [42] revealed that at high temperaturethe fatigue cracks propagated along the c channels and avoided cut-ting the c0 phase, thus fatigue cracking was perpendicular to thestress axis. But for specimen with c/c0 rafts parallel to the stressaxis, the crack propagated obliquely to stress axis.

In fact, single-crystal (SX) superalloys have different deforma-tion mechanisms, slip morphologies and cracking modes at differ-ent temperatures. The addition of the alloy elements also affectsthe dislocation movement and distribution. Therefore, both tem-perature and the alloy elements should affect the LCF behaviorsof nickel-base SX superalloy together with influencing the SFE va-lue, the c/c0 relative strength as well as the interface mismatch.However, the systematic research and comprehensive analysisconsidering both the effects of temperature and alloy elementsare rather scarce. In view of this, the present study selects twonickel-base SX superalloys with and without the element Re asmodel materials and concentrates on discussing the roles of theabove factors in the LCF processes and damage mechanisms ofnickel-base SX superalloy.

2. Experimental procedures

Two sets of nickel-base SX superalloys containing 3 wt.% Re (al-loy 3Re) and without Re (alloy 0Re) were prepared by means ofcrystal selection method in a directional solidification vacuum fur-nace under a high thermal gradient. The alloy 0Re contains 6 wt.%W, while the alloy 3Re contains 3 wt.% W and 3 wt.% Re. The nom-

inal compositions of the two alloys are listed in Table 1 in detail.Solution treatment by 1300 �C/1 h + 1305 �C/3 h + 1315 �C/4 h(AC) and two-step aging treatment by 1080 �C/6 h (AC) + 870 �C /24 h (AC) were carried out on the two alloys. Following the heattreatment, samples with a diameter of 6 mm and a gauge lengthof 15 mm were made using an electro-spark cutting machine.LCF specimens with [001] orientation were machined parallel tothe longitudinal direction of the heat-treated bars. The LCF testswere all conducted with an Instron8862 servo hydraulic testingmachine, and were carried out under total strain (Det = 1.6%) con-trol in air at 25 �C, 250 �C, 500 �C, 750 �C and 900 �C, respectively.A triangle waveform with a constant strain rate of 5 � 10�3 s�1

and strain ratio R = �1 was used. Table 2 shows the variation ofYoung’s modulus and plastic strain amplitude with temperature.It can be found that the Young’s modulus gradually decreases withincreasing temperatures in both alloys.

After fatigue test, the surface morphologies, crack initiation andpropagation of all the specimens were observed by a LEO supra35field emission scanning electron microscope (SEM). Some speci-mens after high temperature LCF tests were polished to producea mirror-like surface for crack tip observations after etching with7% HClO4 + 90% Alcohol. The deformation microstructures afterLCF failure were observed with the help of transmission electronmicroscope (TEM). TEM observations were generally performedon {001} thin foils in order to examine interface dislocations andclearly exhibit the dislocation activity within the channels. Thesamples were initially thinned mechanically to 50 lm, then elec-tropolished at 15 V in a solution of 10% perchloric acid and 90%ethanol using the double-jet polishing technique at room temper-ature. TEM observations were performed on a Tecnai F20 FEI TEM.Eventually, the reproducibility of all the deformation microstruc-tures, or configurations shown in what follows has been carefullychecked.

3. Results and discussion

3.1. Cyclic stress response of SX superalloys at different temperatures

Fig. 1(a) and (b) shows the cyclic stress response curves of thetwo alloys tested at different temperatures. At all testing tempera-tures, the alloys 0Re and 3Re both displayed fairly constant stress

Fig. 1. Comparison of LCF behaviors between the two SX superalloys at different temperatures: (a) cyclic stress response curves of alloy 0Re; (b) cyclic stress response curvesof alloy 3Re; (c) saturation stress amplitude-temperature curves; and (d) Dep–Nf curves of the two alloys.

P. Li et al. / International Journal of Fatigue 63 (2014) 137–144 139

amplitudes from the start of the test. By comparing the two sets ofcurves, it can be found that with increasing the testing tempera-tures, the average stress amplitudes of the two alloys gradually de-creased. For the alloy 0Re, the longest fatigue life appeared at500 �C. However, the longest fatigue life of the alloy 3Re occurredat 750 �C. It seemed that the addition of Re improved the temper-ature at which the longest fatigue life appeared. Meanwhile, at thesame temperature, the stress amplitude of the alloy 3Re is slightlyhigher than that of the alloy 0Re, but the fatigue life of the alloy3Re is obviously higher than that of the alloy 0Re. At room temper-ature, the two alloys showed continuous cyclic softening, whichcould be caused by the shear-induced specimen damage. From250 �C to 750 �C, the cyclic stress response curves of the two alloysexhibited a distinct plateau behavior which meant that the stresslevel almost kept constant before the final fracture. At 900 �C, thetwo alloys displayed the slow and sustained cyclic softening duringthe whole cyclic deformation.

Besides the effect of temperature, the effect of Re on the LCFbehaviors of SX superalloys should not be negligible. Fig. 1(c) dis-play a comparison diagram of the two alloys concerning the stressamplitude at a certain temperature. At all testing temperatures, thestress amplitudes of the alloy 3Re are generally 20–40 MPa higherthan those of the alloy 0Re. Fig. 1(d) describes the relationshipbetween plastic strain amplitude and the fatigue life of the twoalloys. With decreasing the plastic strain amplitude, the corre-sponding fatigue lives are increasing by degrees. Based on the cyc-lic deformation behaviors of the two alloys, it can be understoodthat the deformation mechanisms of SX superalloys at differenttemperatures are distinctly different. The effect of temperature

on the deformation mode of SX superalloys mainly focused onthe surface slip morphologies, dislocation movement and crackingmodes, which will be analyzed and discussed in detail in the fol-lowing sections.

3.2. Surface morphologies and cracking mode at differenttemperatures

From room temperature to high temperature, the two alloysshowed a similar behavior. Fig. 2 summarizes the behavior andpoints out the effect of Re on the surface morphologies, especiallyslip mode. At room temperature, numerous slip band (SBs) appearon the specimen surface, the severe intrusion and extrusion repre-sent the strain localization and eventually the crack propagatesalong multiple sets of SBs simultaneously with cutting c0 phase.At 500 �C, the degree of intrusion and extrusion becomes weakand the number of activated slip systems is lower. Around900 �C, the SBs disappear and the deformation becomes homoge-neous, which results in the cracking mode changing from shearfracture to normal fracture. Meanwhile, the oxidation aggravatesthe fatigue damage with increasing temperature. Compared withthe alloy 0Re, the addition of Re at both room temperature andmedium temperature stimulates the activation of more slip sys-tems and thus the crack propagates along more slip planes,although the intrusion and extrusion are not so much severe asthat of the alloy 0Re, which will conduct to the extension of fatiguelife. Although the crack propagation in two alloys at 900 �C is per-pendicular to the loading axis, the addition of Re makes the phasesize smaller and phase distribution more homogeneous, which will

Fig. 2. Surface slip morphologies and fatigue crack propagations in the two alloys at different temperatures: (a), (b) and (c) alloy 0Re; (d), (e) and (f) alloy 3Re.

140 P. Li et al. / International Journal of Fatigue 63 (2014) 137–144

slow down the degree of oxidation and improve the fatigueproperties.

Compared with the alloy 0Re, the surface morphologies of thealloy 3Re become more complex because of the addition of Re.Thornton et al. [43] studied the temperature dependence of theyield stress based upon Ni3Al. They found that with increasingtemperature, the micro-plastic yield stress approaches a peak va-lue at approximately 800 �C, where the maximum stress is closelyrelated to the alloying element and strain rate. The addition of Tiand Cr in Ni3Al can significantly increase the maximum stressand the temperature up to the peak. The authors also consideredthat the operation of the additional {100} slip systems is arisinglargely as the deformation temperature is increased and the alloy-ing addition affected the temperature dependence probably byaltering the tendency for {100} slip. Then, along with temperature,the addition of Re is expected to affect the LCF behaviors of nickel-base SX superalloys.

In summary, the cracking mode is closely related with surfaceslip morphologies. Different slip morphologies reflect the variationof dislocation movement with slip character. Therefore, the furtherdiscussion will focus on the regularity of dislocation movement intwo nickel-base SX superalloys.

3.3. Dislocation movement and deformation mechanisms at differenttemperatures

Fig. 3 shows the fatigued dislocation structures of the two alloysat room temperature. A large number of planar slip bands corre-sponding to the macroscopic SBs form in the alloy 0Re, as shownin Fig. 3(a). These bands mainly consist of high density of tangleddislocations bowing out the matrix and continuously cut throughc0 phase, which is also found in the alloy 3Re. Petrenec et al. [17]found that the observed band parallel to the {111} slip planesand the structure of this band is similar to the ladder structure ofthe persistent slip band observed in cyclically deformed fcc metals[44–47]. Only a small number of SFs appear in the bands of the al-loy 0Re. Unlike the alloy 0Re, two sets of dislocation bands appearin the alloy 3Re (see Fig. 3(b)) and plenty of SFs form within thesebands (see Fig. 3(c)). Together with dislocation tangles and pileup,they are composed of the dislocation bands. Antolovich et al.

[48,49] found that the appearance of dislocation substructure in2D can be very misleading. They used one 3D technique in a studyof the LCF of the nickel-based superalloy René 95, which resulted ina completely new interpretation from traditional 2D micrograph.In addition, the addition of Re reduced the SFE of the alloy at roomtemperature and planar slip characteristic becomes more signifi-cant. On the other hand, more anti-phase boundary (APB) existedin the alloy 3Re [50], which made the activation of new slip systemeasier due to the stronger resistance, as shown in Fig. 3(d) and (e).In any case, the dislocation movement is mainly based on the pla-nar slip at room temperature. The dislocations cut through c0 phaseto form plenty of planar SBs by planar slip, which displayed theclassical LCF mechanism of the two alloys, no matter whether0Re or 3Re, as shown in Fig. 3(f) schematically.

With increasing temperature, the slip mode gradually changesfrom planar slip to wavy slip. Meantime, the hardness of the twoalloys also begins to change. Daymond et al. [51] compared thetensile behaviors of polycrystalline nickel-base superalloy at differ-ent temperatures using neutron diffraction and found that withincreasing temperature the relative hardness of the two phasespresent in a c0 superalloy gradually changed. Between roomtemperature and 500 �C, the c0 phase is the softer phase. At highertemperatures, the c matrix is softer. Therefore, when the test tem-perature rises to 750 �C, the dislocation movement in fatigued SXsuperalloys includes both planar slip and wavy slip. As shown inFig. 4, no matter whether 0Re or 3Re alloys, the dislocation bandsdisappeared and instead numerous dislocations bowed out aroundc0 phase. Fig. 4(a) and (b) shows the dislocation structure of the al-loy 0Re. The dislocation lines are parallel to the (001) plane andthe expanding dislocation loops display very long and narrowshapes with zigzagged features. However, the dislocations of thealloy 3Re in Fig. 4(c) are of remarkably different morphologies. Incontrast, by cross-slip in the horizontal c channels, in some regionsc/c0 interfacial dislocation networks have formed as shown inFig. 4(d). Zhang et al. [52] pointed out that these dislocations havecompleted the reorientation from the deposition direction h110i tothe mismatch direction h100i in the (001) c/c0 interfacial plane. Infact, the addition of Re increases the lattice mismatch which thenchanges the distribution of the geometrically necessary disloca-tions (GNDs) at the interface. It is well known that the GNDs

Fig. 3. Dislocation configurations in alloy 0Re and 3Re after LCF failure at room temperature: (a) 0Re; (b), (c), (d) and (e) 3Re; (f) dislocation movement at room temperature.

Fig. 4. Dislocation configurations in alloy 0Re and 3Re after LCF failure at 750 �C: (a) and (b) 0Re; (c), (d) and (e) 3Re; (f) dislocation movement at 750 �C.

P. Li et al. / International Journal of Fatigue 63 (2014) 137–144 141

actually lower the interfacial energy by strain compensation, butany change would increase the energy and provide resistance. Thusthe addition of Re will contribute to the interface strengthening ofthe superalloy. Meanwhile, those zigzagged characteristics fromvery long, narrow shaped expanding dislocation loops indicatedthat these dislocations are moving by cross-slip in the horizontalc channels (see Fig. 4(e) and (f)). Eventually the larger lattice misfitresults in denser c/c0 interfacial dislocation networks, which is thekey factor controlling the LCF behaviors of the two alloys at 750 �C.

Fig. 5(a) shows the dislocation structure of the fatigued alloy0Re at 900 �C. In this figure, the deformed zone consisting of thedislocation by-passing with high dislocation density was separatedby low dislocation density areas. Furthermore, matrix dislocation

and interface dislocations can be seen in Fig. 5(b). These resultsare in agreement with other findings on the cyclic deformationin Ref. [25]. Furthermore, Fig. 5(c) demonstrates that the disloca-tion loops are arrested at the c/c0 interface, the arrested interfacialdislocations with Burgers vectors oblique to the c/c0 interface canmove towards the edge of the c0 particles by a combined glide-climb process, including gliding in {111} plane in the matrix andclimbing around the cuboidal c/c0 interface. Similar to the alloy0Re, the fatigued alloy 3Re also demonstrated the dislocation by-passing by glide-climb motion along the c/c0 interfaces inFig. 5(d). Different from the alloy 0Re, the dislocation distributionis much homogeneous in the alloy 3Re and the difference betweendeformed zone (D-zone) and un-deformed zone (U-zone) becomes

Fig. 5. Dislocation configurations in alloys 0Re and 3Re after LCF failure at 900 �C: (a), (b) and (c) 0Re; (d) and (e) 3Re; (f) dislocation movement at 900 �C.

Fig. 6. Effect of Re on deformation mechanisms at 25 �C and 900 �C. (a) Alloy 0Re at 25 �C; (b) alloy 3Re at 25 �C; (c) alloy 0Re at 900 �C; (d) alloy 3Re at 900 �C.

142 P. Li et al. / International Journal of Fatigue 63 (2014) 137–144

very little (see Fig. 5(e)). The formation of homogeneous deformedmicrostructures is probably due to the APB energy effect. In fact, nomatter whether the alloys 0Re or 3Re is, dislocation climb and glidein the c channels are the main deformation mechanisms at 900 �C,as shown in Fig. 5(f) schematically.

The above results showed that the effect of temperature on LCFbehaviors of the alloys 0Re and 3Re is the same. However, thestrengthening mechanisms due to the addition of Re are differentat different temperatures. Therefore, the effect of Re can be dividedinto two cases to be discussed. As shown in Fig. 6(a), the disloca-tions move by planar slip mode at room temperature and formnumerous planar slip bands cutting through the c0 phases. Theaddition of the alloy element Re further decreased the SFE valueof the superalloy so that its planar slip characteristic is more obvi-ous [53]. In order to better coordinate deformation, more slip sys-tems operate as a result of the decrease of the SFE, as shown in

Fig. 6(b). At high temperature, such as more than 900 �C, the dislo-cation movement is composed of the cross slip of screw dislocationand the climb of the edge dislocation along the c channels. Asshown in Fig. 6(c), the deformation tends to homogenize by Oro-wan by-passing at this temperature. It is well known that the ele-ment Re retards all diffusion-driven process and further inhibitsthe growth of c0 phase to reduce the width of the c channel [53],which will induce the activation of more dislocation movementby a combined glide-climb process so that the deformation homog-enization becomes higher, as shown in Fig. 6(d).

Concerning the effect of temperature on the deformation mech-anism of the SX superalloy, Lall et al. [54] proposed a model calledLCP-model. They considered that the increasing strength of c0

phase with increasing temperature is a thermally activatedprocess, where the cross-slip of dislocations from the octahedralplanes {111} to the cube planes {100} governs the plastic

Fig. 7. Variation of deformation mechanisms in the two nickel-base SX superalloys fatigued at different temperatures.

P. Li et al. / International Journal of Fatigue 63 (2014) 137–144 143

deformation of the L12 crystals, including c0 phase. At this point theleading pair of Shockley partials acts as a pinning point to hinderthe glide in the primary (111) plane. Further research work [55]showed that this so-called non-Schmid behavior is a characteristicfeature not only for the L12 crystals, but also for all superalloyscontaining a high, medium or even low volume fraction of the c0

phase, but not for a disordered solid solution of approximatelythe same chemical composition.

According to the above discussion, it should be recognized thatas the temperature changes, the SFE and relative strength of SXsuperalloy are changed to promote the occurrence of differentdislocation movement modes and the formation of various disloca-tion structures. For [001] SX superalloys, the short-distance cross-slip of screw dislocation gradually converts from {111} to {001}planes with increasing temperature. From room temperature toabout 750 �C, the dislocation movement generates a transitionfrom planar to wavy slip. The relative strength of the c0 phasehas gradually increased. Part of the screw dislocations bowed outaround the c0 phase and the more edge dislocations piled up atthe c/c0 interface, as shown in Fig. 7. Eventually the piling-up dis-locations still cut through the c0 phase and bring about the strainlocalization. At 900 �C, the screw dislocation continues to moveby cross-slip and the edge dislocation begins to climb withfreedom. At this temperature, the relative strength of c0 phasehas become enough high. Together with the increase of SFE, it ismore difficult to cut through the c0 phase. The dislocations bypassthe c0 phase by glide-climb motion, which is contributive to thehomogeneous distribution of the dislocations. In addition, withincreasing temperature, the mismatch stresses at the c/c0 interfaceincrease, leading to plenty of misfit dislocations nucleated at the

interface. These misfit dislocations strongly hindered the move-ment of the follow-up dislocations, which also contributes to thedeformation homogenization.

Apart from this, the addition of Re further amplifies the effect oftemperature. On one hand, the SFE value is further reduced by theaddition of Re at the same temperature [53]. On the other hand,with increasing temperature the addition of Re induces the segre-gation of other element which causes the increase of interface mis-match. In short, the SFE value, the c0/c relative strength and theinterface mismatch play a key role in the effect of temperatureon the LCF behaviors of nickel-base SX superalloy. They dominatethe deformation mechanism of the superalloy together by adjust-ing the mode of dislocation movement. More useful discussionswill be made after further research on the LCF behaviors of nick-el-base SX superalloy.

4. Conclusions

The effect of temperature on the LCF behaviors of two nickel-base SX superalloys containing 3 wt.% Re and without Re was stud-ied. Some conclusions can be drawn:

(1) At room temperature and 250 �C, the two alloys show thecontinuous cyclic softening and are dominated by plasticdeformation in the initial stage of cyclic deformation. Bycomparison, the fatigue life of the alloy 3Re is about twotimes higher than that of the alloy 0Re. The SBs appear infatigued two alloys with severe intrusion and extrusionand more slip systems operate in the alloy 3Re. The crackpropagates along multiple sets of SBs. These planar SBs

144 P. Li et al. / International Journal of Fatigue 63 (2014) 137–144

consist of dislocation tangles and SFs mixed in them. Theaddition of Re promotes the activation of more slip systemsand thereby improves the fatigue properties of the SXsuperalloy.

(2) At 500 �C and 750 �C, the cyclic stress response curves of thetwo alloys exhibit a distinct plateau behavior which meansthat the stress level almost keeps constant before final frac-ture. The SBs are still main slip features, but the degree ofintrusion and extrusion becomes weak. Although the addi-tion of Re is likewise beneficial to the activation of slip sys-tems, the total number is still reducing with increasingtemperature. Some cracks initiate along the oxide layerand still propagate along the SBs. The dislocation movementis gradually changed from planar to wavy slip. The formationof dislocation network caused by cross-slip becomes the keyfactor controlling the LCF behaviors of the two alloys.

(3) At 900 �C, the two alloys display similar slow and sustainedcyclic softening, but the addition of element Re improvedthe ability of the resistance to the plastic deformation andextended the LCF lives of SX superalloys. At this tempera-ture, the SBs disappear and the deformation homogenizationbecomes dominant, resulting in the normal cracking of spec-imen and showing mode I crack. Owing to the APB energyeffect, the dislocations display homogeneous distribution.In the c channels, the dislocations bypass the c0 phase by acombined glide-climb process.

Acknowledgements

The authors are grateful to Prof. Z.G. Wang and Prof. S.X. Li fortheir good suggestions and advices. Thanks are also due to Y.J.Xu, K. Du, J.L. Wen and Q.Q. Duan. This work is supported by theNational Basic Research Program of China under Grant No.2010CB631206, the National Natural Science Foundation of China(NSFC) under Grant Nos. 51001104, 50931004 and U1037601.

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