International Journal of Fatigue xxx (2013) xxx–xxx

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International Journal of Fatigue

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In situ SEM studies of the low cycle fatigue behavior of DZ4 superalloyat elevated temperature: Effect of partial recrystallization

0142-1123/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.ijfatigue.2013.11.001

⇑ Corresponding author at: AML, School of Aerospace, Tsinghua University,Beijing 100084, China.

E-mail addresses: maxianfeng@gmail.com (X. Ma), shihj@mail.tsinghua.edu.cn(H.-J. Shi).

Please cite this article in press as: Ma X, Shi H-J. In situ SEM studies of the low cycle fatigue behavior of DZ4 superalloy at elevated temperature: Epartial recrystallization. Int J Fatigue (2013), http://dx.doi.org/10.1016/j.ijfatigue.2013.11.001

Xianfeng Ma a,b, Hui-Ji Shi a,⇑a AML, School of Aerospace, Tsinghua University, Beijing 100084, Chinab Department of Materials Science and Engineering, University of Michigan, Ann Arbor, MI 48109, USA

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

Article history:Received 23 August 2013Received in revised form 28 October 2013Accepted 1 November 2013Available online xxxx

Keywords:Directionally solidified superalloyRecrystallizationLow cycle fatigueShort fatigue crackin situ SEM

The influence of partial recrystallization (RX) on the low cycle fatigue (LCF) behavior of directionallysolidified DZ4 superalloy was studied using in situ scanning electron microscopy. Three typical recrystal-lization microstructures were prepared on the fatigue samples, via shot peening and subsequent heattreatment. The effect of RX on LCF life is evidently related to the specific RX microstructure. The RX layerdoes not necessarily reduce the LCF life. The fatigue crack initiation and short crack propagation behav-iors for both raw DZ4 and recrystallized samples were in situ examined, which rationalized the observedeffect of recrystallization on the LCF life variation.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

By using the modern directional solidification technology, themechanical properties of superalloys have been significantly im-proved due to elimination of the transverse grain boundaries[1,2]. Advanced superalloys have been developed and widely ap-plied in gas turbine components [1,3], e.g., directionally solidified(DS) or single crystal (SX) materials, because of their outstandingthermal, fatigue and creep properties under severe operation con-ditions. As a well-known issue in gas turbine industry, the superal-loys tend to be susceptible to environmental influence on thesurface, for instance, oxidation or recrystallization (RX) layersmay be formed during manufacturing process or high temperatureapplication. Specifically, the recrystallization grains were found onthe leading and trailing edges of as well as at the root or the coolingchannels of gas turbine blade. Zhang et al. [1] studied the fracturedturbine blades served in Chinese model fighters and observedintergranular cracks originated from the surface recrystallized areaon the leading and trailing edges of the DS blade. Studies in recentdecades [4–8] tend to suggest that the recrystallization, as an alienmicrostructure, is a safety threat to the DS/SX turbine components.

Generally one thinks that the surface recrystallized layer im-ports equiaxed crystal structure which is different from the trans-

versely isotropic DS superalloy with columnar crystal structure.Thus, the deformation incompatibility at the RX/matrix interfacewould induce stress concentration and may facilitate the cracknucleation under cyclic loading. Moreover, owing to the absenceof transverse grain boundary (GB), directionally solidified (DS)and single crystal (SX) alloys generally reduce the concentrationof grain boundary strengthening elements (e.g. B) which may de-crease the creep resistance and melting temperature of superalloy,hence they are able to maintain the strength and operate at highertemperature [1]. In this case, the transverse GB’s introduced byrecrystallization are relatively weak sites, especially for the turbineblades serving at elevated temperatures. It explains why intergran-ular cracking was observed in the recrystallization zone of turbinecomponents [1,9].

There has been some work on the recrystallization behaviors ofDS superalloy, such as the effects of deformation and heat treat-ment parameters on recrystallization [10,11]. In recent decades,more attention has been drawn to the influence of recrystallizationon the mechanical properties of DS alloys, e.g. the stress enduranceor creep property [6,7]. However, research work concerning the ef-fect of RX on low-cycle fatigue (LCF) is surprisingly limited in openliterature. And the studies on different alloy systems under differ-ent experimental conditions have not yet produced conclusive re-sults. For instance, the fracture mechanism of recrystallized DSblade is not yet clarified and conditions determining whether therecrystallized layer is fatal or inessential remain ambiguous. Zhanget al. [12] reported that DZ4 superalloy with continuous RX layerhad a fatigue life which was 10–50% of that of the original super-

ffect of

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alloy. Jia et al. [13] studied the same DZ4 alloy but at a higher tem-perature of 760 �C, which showed continuous decrease of LCF lifewith increasing RX thickness. Another study on DZ40 M DS alloyby Zhao et al. [8] showed a 60% decrease of fatigue life by a100 lm RX layer. Whereas LCF tests on several single crystal super-alloys with cellular recrystallization by Bürgel et al. [9] revealedthat the crack initiation life of surface recrystallized specimenhad no significant difference with the raw alloy, except for a muchhigher crack density.

On the other hand, due to lack of reliable experimental observa-tions, questions remain whether crack initiation/propagation inrecrystallized layer or the subsequent propagation in DS alloy isdominant in assessing the influence of RX on LCF life. Direct obser-vation of fatigue microcrack initiation and propagation is thereforein great need as the recrystallized specimen may exhibit quasi-brittleness and thus make the crack recording much more difficultand laborious by other methods such as plastic replica technique.With the development of observational techniques, in situ obser-vation with scanning electron microscope (SEM) can greatly en-hance the chance of identifying the surface crack at its very earlystage and quantitatively analyzing the effect of microstructure onthe small crack in low-cycle fatigue [14], as small as a few micronsor even smaller, enabling better understanding of the fracturemechanisms of materials based on the grain morphology, misori-entation and secondary phase etc. [15,16].

The present work attempts to add light to these issues. A seriesof RX microstructures by different shot peening (SP) pressures andheat treatments were specially prepared in directionally solidifiedDZ4 superalloy. The RX specimens, together with the original DZ4specimens, were subjected to LCF tests at 350 �C under in situ SEM.Smooth samples were used to study the influence of RX on the LCFlife. Also, a few notched samples were made to study the fatiguecrack growth behavior. Particular emphasis was placed on examin-ing the early stage of fatigue, including the crack initiation mecha-nisms in terms of RX microstructures and the short crack growthdata for a subsequent quantitative analysis.

2. Material and experimental procedure

The material used is a DS nickel-based superalloy DZ4 devel-oped for fabricating high-performance turbine blades in China.The nominal composition (in wt.%) is: 0.14C, 9.5Cr, 6.0Co, 5.3W,3.8Mo, 6.0Al, 1.8Ti, 0.02B and Ni balance. The microstructure iscomposed of paralleled columnar grains with their h001i axisalong the DS direction [17], as shown in Fig. 1a. Blocky carbideswere widely distributed in the alloy, which were identified to beMC type carbides (Fig. 1b). All the fatigue specimens were cut with

Fig. 1. Microstructure of DZ4 superalloy (a) o

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gauge length parallel to the solidification direction. The slab spec-imen has a dog-bone shape with a 2.5 mm by 0.4 mm gauge crosssection. For the notched specimens, a U-shaped notch of 50 lmdepth and about 30 lm radius was prepared manually [17] toinvestigate the fatigue short crack behaviors (Fig. 1a). To revealthe microstructure, the sample surfaces were polished and etchedin a solution of 4 g CuSO4 + 20 ml HCl + 20 ml H2O prior to fatiguetesting.

Partial recrystallization was specially prepared on the lateralsurface of fatigue specimens by shot peening and a subsequentheat treatment at 1220 �C for 4 h, followed by air cooling. To pre-vent oxidization, the specimens were enclosed in a quartz tubein which a positive pressure of argon gas was maintained. Shotpeening with three different pressures, i.e. 0.1 MPa, 0.3 MPa and0.5 MPa, were used to generate different RX at the specimen’s worksection. Hereafter the original and recrystallized DZ4 specimenswill be referred to as noRX, RX0.1, RX0.3 and RX0.5 respectively.

The fatigue tests were performed at 350 �C in the vacuumchamber of the SEM using a specially designed servo-hydraulictesting system. The elevated temperature was obtained by heatof resistance coils, accompanied with adjustable thermocouplesto get a stable control of temperature along the gauge length.The temperature profile remained within about ±1 �C of the nom-inal temperature throughout the test. The SEM was operated at anaccelerating voltage of about 15 kV. The fatigue tests were pausedat different loading cycles to capture the SEM images of fatiguecracks under the secondary-electron mode. More details aboutthe in situ SEM fatigue testing system can be found in Refs.[16,17]. The fatigue tests were load controlled under a maximalstress of 850 MPa and stress ratio of R = 0.1 at 2 Hz. After fatiguetesting, the samples were subjected to fractographic analysis usinga Scanning Electron Microscope of Hitachi S-4500.

3. Results and discussion

3.1. Microstructure of recrystallization

Fig. 2 shows the recrystallization formed on the specimen sur-face after shot peening and subsequent heat treatment. The surfacerecrystallization area was featured by the transverse grain bound-aries of equiaxed RX grains. No cellular recrystallization was ob-served in the samples. According to Wang’s study on DZ4 alloy[10], recrystallization occurs via fully recrystallized grains contain-ing comparable c0 size with that of original DZ4 when heat treatedat 1220 �C or above, as the c0 phases dissolve completely. Primarycarbides (MC) were widely distributed in RX grains and interden-dritic regions as reported in other DS superalloy [11], while the

verall view and (b) blocky MC carbides.

cycle fatigue behavior of DZ4 superalloy at elevated temperature: Effect ofe.2013.11.001

Fig. 2. Microstructure of recrystallized DZ4 sample obtained by shot peening of (a) 0.1 MPa, (b) 0.3 MPa and (c) 0.5 MPa, and heat treatment.

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secondary fine carbides (M23C6) precipitated preferentially on RXgrain boundaries acting with strengthening effect [15].

The shot peening pressure had an evident effect on the an-nealed RX microstructure, as shown in Fig. 2a–c. A continuouslayer of RX grains were generated for those at high shot-peeningpressures (0.3 MPa and 0.5 MPa), whereas discrete RX grains wereobtained for low shot-peening pressure of 0.1 MPa due to relativelysmall plastic deformation introduced [10]. The average depth of RXlayer is measured to be about 80 lm, 200 lm and 300 lm forRX0.1, RX0.3 and RX0.5 respectively. It is reasonable that the RXdepth increased with the increase of shot peening pressure, be-cause the plastic deformation induced by shot peening determinedthe amount of RX that can be generated in a full solution heattreatment (1220 �C/4 h in this study) [10]. Previous studies in sim-ilar directionally solidified superalloys have showed that the an-nealed recrystallization grains tend to have widely dispersedcrystal orientations [10,18].

3.2. Effect of recrystallization on LCF life

The fatigue lives of the noRX and recrystallized specimens wereplotted in terms of the shot peening pressure in Fig. 3, showingthat the recrystallization had a significant effect on the fatigue life.Especially for RX0.1 samples, they exhibited fatigue life of 5000–10,000 cycles in contrast to over 20,000 life cycles for noRX sam-ples, i.e. the fatigue life under rmax = 850 MPa was reduced bymore than 50% due to the surface RX. This degradation of fatiguelife by partial recrystallization is similar to that reported by Zhaoet al. [19]. Nevertheless, it is interesting to notice that the LCF lifeof recrystallized DZ4 gradually increased with SP pressures higher

Fig. 3. Effect of recrystallization on the fatigue life of DZ4 alloy.

Please cite this article in press as: Ma X, Shi H-J. In situ SEM studies of the lowpartial recrystallization. Int J Fatigue (2013), http://dx.doi.org/10.1016/j.ijfatigu

than 0.1 MPa, even exceeding that of noRX samples when the SPpressure was 0.5 MPa. The fatigue life of RX0.5 sample can be aslong as 43,654 cycles, which is about twice of that of noRX sample.This novel result has been rarely reported in the previous studiesregarding the effect of RX on the fatigue/creep properties of DSor SX superalloys. The recrystallization has been usually regardedas most detrimental to the turbine superalloys [1,5–8] due to thealien transverse boundaries, especially after the RX grains were re-lated to the crack initiation sites on some fractured DS blades [1].

3.3. Fatigue crack initiation

To understand the underlying mechanism for the effect of RX onfatigue life in Fig. 3, the fatigue crack initiation behaviors of bothnoRX and RX specimens were examined. The loading direction isalong vertical in Fig. 4a and b. For original DZ4 superalloy, inter-dendritic carbides vertical to the loading direction was prone tobe fractured in the early stage of cyclic loading. Multiple crackingof transverse carbides could be observed, among which the mostfavored crack further extended as the main crack, as shown inFig. 4a. The crack might grow into long crack by coalescence withother short cracks nearby. The transverse interdendritic bound-aries with inhomogeneities (e.g. blocky carbides) were reportedto be the crack initiation site in the low-cycle fatigued DZ4 [13].The in situ observation here indicated that it was the preferentialcracking mode for DZ4 under study.

For RX0.1, DZ4 with discrete RX grains were shown in Figs. 2aand 4b. The blocky carbides were commonly observed on the RXgrain boundaries, due to easier nucleation of recrystallizationaround these inhomogeneities. Moreover, the RX grains tend tohave orientations different from the h001i column grains. Hence,the difference in modulus and plastic properties would promotethe deformation incompatibility, leading to the cracking of thesebrittle carbides, which rationalized the crack initiation mode inRX0.1 specimens (arrowed in Fig. 4b). One of these small crackspropagated and evolved into the main crack. It is emphasized thatthe crack preferred transgranular mode in Fig. 4b, which is distinctfrom the reported intergranular cracking due to the weak trans-verse RX grain boundaries [6,8]. It is suggested that the differencein fracture mode is attributable to test condition, especially thetemperature [15], i.e. transgranular cracking dominates belowthe iso-strength temperature Ts (between 350 �C and 760 �C forDZ4).

For the RX0.3 and RX0.5 specimens, the main cracks were notdetected before the fracture of specimens, except for some non-propagating cracks along carbides. A post-examination of the frac-tured specimen was done under SEM. Fractographic analysis re-vealed that the crack in RX0.3 specimen was originated fromcarbides in RX layer (Fig. 4c), whereas the RX0.5 specimen withthe longest fatigue life showed to be originated from near-surface

cycle fatigue behavior of DZ4 superalloy at elevated temperature: Effect ofe.2013.11.001

Fig. 4. Fatigue crack initiation for different specimens (a) original DZ4: from interdendritic carbides, (b) RX0.1: from large carbides near RX boundary, (c) RX0.3: from carbidesin RX layer and (d) RX0.5: from near surface casting pores or carbides in RX layer.

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casting pore or carbide in RX layer (Fig. 4d). The river patterns wereevident, indicating crystallographic fracture along active slipplanes. These planes in Fig. 4d were identified to be {111}, as theoctahedral slip was preferentially operated at lower temperatureand higher frequency in FCC superalloys [2]. Crystallographic frac-ture was commonly observed in all examined specimens. Regard-ing the carbides that initiated cracks during fatigue loading, astatistical analysis of the carbide size (denoted by square root ofthe area) and distribution was shown in Fig. 5. In practical case, lar-ger blocky carbides would facilitate the fatigue crack initiation inthis alloy. The observations in this section indicated that the differ-

Fig. 5. Statistic of the fractured carbides during the low cycle fatigue of DZ4 alloy.

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ence in crack initiation may be responsible for the observed lifevariation with different RX.

3.4. Fatigue crack propagation behaviors

The fatigue crack propagation features were also studied byin situ SEM for both noRX and RX specimens. Edge notch was spe-cially prepared on the samples of noRX, RX0.3 and RX0.5 to localizethe crack initiation site. For RX0.1 specimens which did not haveuniform RX layer as in RX0.3 and RX0.5, smooth samples wereused, in which the fatigue cracks could be found around the car-bides on RX grain boundaries (see Fig. 4b). Fig. 6 shows the fatigueshort crack propagation in the noRX specimen. Early in the fatigueloading process (less than 10% of the total life), crack initiated fromthe notch root (Fig. 6a). Notch depth is included in the crack lengthin the caption of Fig. 6, also adopted in Figs. 7–9. In the followingthousands of cycles, the small crack had a slow development(Fig. 6b). The cracking mechanism was crystallographic fractureof the dendrite of DZ4 column grains, which was driven by re-peated cycling on the plastic zone ahead of crack tip. This is similarto the fatigue cracking of DZ4 at room temperature in vacuum [12].After 13,810 cycles in total, a slantwise crack emerged ahead of theformer crack tip (arrowed in Fig. 6b). Subsequently, the crack prop-agated along this inclined plane, accompanied with evident sliptraces especially on the upper side of crack, indicating the crackis driven by shearing mode (Fig. 6c). In Fig. 6d, duplex slip traceswas featured on both sides of the crack tip and the crack propaga-tion is about 45� inclined to the loading axis.

For RX0.1 sample in Fig. 7a, cracking of carbides was commonlyobserved in early period (indicated by arrows), in addition to thefractured carbides on RX grain boundary from which the fatigue

cycle fatigue behavior of DZ4 superalloy at elevated temperature: Effect ofe.2013.11.001

Fig. 6. Fatigue crack growth of original DZ4 specimen (a) 1992 cycles, 59.9 lm, (b) 13,810 cycles, 93.4 lm, (c) 15,858 cycles, 110.9 lm and (d) 16,092 cycles, 135.1 lm.

Fig. 7. Fatigue crack growth of RX0.1 specimen: crack I (a) and (c) and crack II (b) and (d). (a) 11,569 cycles, 47.1 lm, (b) 11,569 cycles, 78.5 lm, (c) 14,005 cycles, 105.8 lmand (d) 14,005 cycles, 90.0 lm.

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Fig. 8. Fatigue crack growth of RX0.3 specimen (a) 63 cycles, 69.2 lm, (b) 1915 cycles, 106.8 lm, (c) 2101 cycles, 115.7 lm and (d) 2272 cycles, 148.4 lm.

Fig. 9. Fatigue crack growth of RX0.5 specimen (a) 7120 cycles, 64.3 lm, (b) 11,069 cycles, 154.4 lm, (c) 11,189 cycles, 157.7 lm and (d) 11,403 cycles, 229.7 lm.

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Please cite this article in press as: Ma X, Shi H-J. In situ SEM studies of the low cycle fatigue behavior of DZ4 superalloy at elevated temperature: Effect ofpartial recrystallization. Int J Fatigue (2013), http://dx.doi.org/10.1016/j.ijfatigue.2013.11.001

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crack initiated. The crack propagated in a transgranular mannerevidently, similar to that in Fig. 6. The crack path appeared curving,indicating the interaction of small crack growth with the localmicrostructure. Upper slip traces were dominant after exceedingcrack length of 50 lm, indicating this RX grain is probably sin-gle-slip-oriented. Another macroscopic crack in this specimen ap-peared in duplex slip manner (see Fig. 7b). Of interest is theinteraction between fatigue crack and carbides. In Fig. 7b, carbide‘A’ served as fatigue crack origin, whereas carbide ‘B’ in the wayof crack was passed by, acting as obstacle to impede crack growth.Carbides ‘C’ and ‘D’ were by-passed, which is responsible for thecrack path fluctuation and variation in growth rate. After 14,005cycles, the crack in Fig. 7a and b evolved into that in Fig. 7c andd respectively. After the crack reached the interface between RXgrain and matrix DZ4, it took thousands of loading cycles for thecrack to penetrate it (Fig. 7d), implying the blocking effect of theinterface.

For RX0.3 specimen in Fig. 8, the predominated transgranularcracking mode was similar to that of RX0.1. The crack initiatedfrom carbide near the notch root (Fig. 8a). Subsequently, the inter-action of fatigue crack and carbides, i.e. by-pass or pass through, aswell as the inhibition of crack growth (Fig. 8b) was observed [15]. Asharp increase in fatigue crack growth rate (FCGR) was observedafter passing through the carbide in Fig. 8b. Slip bands were ob-served in front of the crack tip in Fig. 8c. The zig-zag cracking man-ner was evident in Fig. 8d, in which the other slip band wasalternatively operated when one slip band was blocked by disloca-tion piling-up. The specimen was fractured before the crack tipcompletely passed through the RX/DZ4 interface.

For RX0.5 specimen, interaction between small crack and car-bides was observed in Fig. 9a, similar to that in Fig. 8b. Slip tracesterminated at the carbides (Fig. 9a). The fatigue crack deflectedfrom the horizontal direction when the crack was about74.63 lm (Fig. 9b). Subsequently the crack propagated inclinedto the loading direction, accompanied with evident slip traces.The interaction between persistent slip band (PSB) and twinboundary (TB) (Fig. 9c) was discussed elsewhere [15]. A remark-able increase in crack opening displacement (COD) was noticedin Fig. 9c compared to Fig. 9b, with no evident increase in cracklength. This was also observed when the crack tip was blockedby carbides in Fig. 8b. The COD method is one of the most effectivemeasures for the FCGR in low cycle fatigue [16]. However, theabove observation provided direct evidence of deviation from therelationship between COD and FCGR, i.e. when crack growth inter-acted with local microstructure. Crack coalescence can be anotherchallenge in quantitative evaluation of small crack growth rate. Forinstance, measuring COD would be difficult in Fig. 9d when theblunted old crack tip was replaced by the sharp new crack tip.The ligament in Fig. 9d endured plastic slip and its fracture dramat-ically increased the crack length and the stress intensity corre-spondingly, leading to the catastrophic rupture soon.

Fig. 10. Fatigue crack growth rate versus the SIF range for different DZ4 specimens.

3.5. Evaluation of fatigue crack growth rate

The fatigue crack length was measured from the high resolutionSEM images captured at different loading cycles during the in situobservations (e.g. Figs. 6–9). The fatigue crack growth can be esti-mated by the stress intensity range as first proposed by Paris andErdogan [20], i.e.

da=dN ¼ CðDKÞm ð1Þ

where a is the fatigue crack length defined as the projected lengthvertical to loading axis, DK is the range of stress intensity factor(SIF), C and m are known as Paris parameters.

Please cite this article in press as: Ma X, Shi H-J. In situ SEM studies of the lowpartial recrystallization. Int J Fatigue (2013), http://dx.doi.org/10.1016/j.ijfatigu

For single edge notched plate of finite width in this study, theSIF range is given by [21]:

DK ¼ Drffiffiffiffiffiffi

pap

Fða=wÞ ð2Þ

where Dr is the stress range, a is the total crack length, and w is thewidth of specimen. F(a/w) is a function of a/w to account for thegeometry effect [21].

Fig. 10 shows the fatigue crack growth rate as a function of theSIF range for each type of specimen studied in this experimentalextent. The data points corresponding to the primary crack of thespecimens in Figs. 6–9 were used. Another specimen of noRXwas also displayed here (marked as red open circle). It is noticedthat the fatigue crack growth rates followed the same trend asthe fatigue lives regarding the effect of RX. For recrystallized spec-imens, RX0.1 and RX0.3 showed overall higher FCGR, indicating thedeterioration of fatigue crack growth resistance due to recrystalli-zation. Especially, the RX0.1 specimen showed the highest FCGR,which could be used to explain its lowest LCF life. For RX0.5 spec-imens with compact RX grains, although the RX layer was thickest,enhanced fatigue crack resistance was observed, which partlyrationalized the improved fatigue life performance in Fig. 3.

For RX0.1 sample, it was most featured by the isolated RXgrains, which might possess orientations other than the [001] ori-entation of as-cast DZ4 column grains. Hence, the preferentiallyoriented RX grains would be subjected to higher local stress dueto difference in modulus and plasticity, which would promotethe subsequent fatigue crack growth in these grains. For RX0.3and RX0.5 samples, thicker continuous RX layers were formed.Especially, the multiple refined RX grains in RX0.5 sample weremost beneficial to the fatigue resistance of DS superalloy (seeFig. 2c). Firstly, it increased the yield stress according to theHall–Petch relationship [22], that is, a stronger surface layer wouldprohibit the crack initiation due to smaller slip length for cyclicstrain localization; Secondly, the higher density of grain boundarywas favorable to hinder the fatigue crack growth [17,23]. In addi-tion, annealing twins were observed in RX layers after solutionheat treatment (Fig. 9). TB is a special kind of coherent high-angleGB with the lowest interfacial energy [24], which has been exper-imentally supported to provide an intrinsically higher resistance tofatigue cracking than the high-energy GB [24]. The impeding effectof TB was manifested by a deceleration of crack growth and possi-ble arrest, as observed in RX0.5 sample (Fig. 9c).

3.6. Indication from the influence of recrystallization

The results in this study showed that the effect of recrystalliza-tion on LCF behavior was closely related to the specific RX micro-

cycle fatigue behavior of DZ4 superalloy at elevated temperature: Effect ofe.2013.11.001

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structure, which was different from the previous understandingthat thicker recrystallization layer would lead to larger drop inmechanical performance [4,6,8,13]. It was found that the discreteRX grains were most detrimental to the fatigue performance ofDS superalloy, which led to not only multiple crack initiation atearlier stage but also a promoted FCGR. While on the other hand,the RX0.5 sample with compact RX layer showed beneficial effectof RX on crack initiation and a lower crack growth rate due tothe refined grains and annealing twin boundaries. For RX0.3 sam-ple, there was a competition between the two mechanisms men-tioned above.

Regarding the enhanced fatigue life in RX0.5 samples, althoughthis phenomenon has been rarely reported in DS or SX superalloys,there is a previous study on a coarse-grained K417 nickel-basedpolycrystalline superalloy [25]. The authors succeeded in improv-ing the fatigue life of K417 alloy by introducing a 150 lm-thickRX layer (grain size of 80 lm) on the surface. And they attributedthe fatigue life enhancement to smaller RX grain size, grain bound-ary effect and better surface quality after recrystallization [25]. Shiet al. [26] also observed a comparable fatigue life of recrystallizedDZ4 sample by high shot peening pressure as compared to that oforiginal DZ4. Note that in the above cases, the recrystallizationmicrostructure, i.e. the morphology/size of c0 is an important as-pect in affecting the strength of the alien RX grains. With lowerannealing temperatures, cellular recrystallization would be ex-pected, as in Bürgel’s [9] and Okazaki’s [5] work. As the cellularRX microstructure had a much coarser distribution of c/c0, the sur-face RX layer was relatively weak according to the c0 size depen-dence of strength in nickel-based superalloy. While in this studyafter recrystallization at solution temperature of 1220 �C, the c0

in RX grains has comparable microstructural feature as the as-castc0 [10]. Hence it allows the possibility of surface strengthening ef-fect and barrier effect of recrystallized GB and TB.

Another benefit of this strengthening RX layer should be consid-ered. As the dislocation and dynamic recovery become easier athigh temperature [27], it may decrease the contribution ofstrengthening techniques associated with high dislocation density,e.g. shot peening and surface mechanical attrition treatment(SMAT) technique. Other techniques by introducing surface nano-structure are also disabled at high temperature due to recrystalli-zation and/or growth of the nanograins. However, the RXmicrostructure of grains and twins obtained in this study (e.g.RX0.5 samples) are thermally stable [10,27]. This point is quiteimportant and attractive for turbine superalloys used at hightemperature.

Also it should be noted that the above discussion is under thepresent experimental condition, i.e. the temperature. At high tem-peratures (e.g. at 760 �C), intergranular cracking along the GBs andTBs instead of transgranular cracking is expected to be the domi-nant fracture mode of RX grains. The effect of RX layer on LCFshould be reevaluated in that case. The transgranular fracture ofRX layer observed in this study implies that intergranular crackingis not always predominant or responsible for fatigue failure ofrecrystallized DS superalloy. The fatigue crack initiation modeand crack growth resistance showed dependence on the micro-structural characteristics of RX layer. As a result, it had a non-monotonic influence on the fatigue life of DZ4 superalloy. Furtherinvestigation is underway in quantitative prediction of the LCF lifeof recrystallized DS superalloy.

4. Conclusions

By in situ SEM fatigue testing, the effect of partial recrystalliza-tion on the LCF behavior of DZ4 superalloy was studied. The obser-vations have resulted in the following conclusions:

Please cite this article in press as: Ma X, Shi H-J. In situ SEM studies of the lowpartial recrystallization. Int J Fatigue (2013), http://dx.doi.org/10.1016/j.ijfatigu

(1) Three typical recrystallization microstructures wereobtained by varying shot peening pressure: discrete RXgrains of 80 lm thick at 0.1 MPa, 200 lm-thick RX layersat 0.3 MPa, and 300 lm-thick RX layer at 0.5 MPa.

(2) The LCF life of DZ4 decreased with the recrystallizationwhen the shot-peening pressure was below 0.3 MPa, butincreased dramatically for RX sample with shot-peeningpressure of 0.5 MPa. The partial recrystallization generatedunder 0.1 MPa shot peening is most detrimental to the LCFproperty of DZ4 alloy.

(3) The interdendritic carbides were the preferential crack initi-ation sites for original DZ4 alloy under study. For RX0.1 sam-ples, multiple crack initiation from the blocky carbides onRX GB was predominant. For RX0.3 and RX0.5 samples, fati-gue crack primarily initiated from the carbides or castingdefects in the RX layer.

(4) Transgranular cracking of RX grains was dominant for allrecrystallized DZ4 specimens in this study, indicating theimportance of predicting the fracture mode according tothe exposed temperature. Short crack growth in RX layerand DZ4 showed interaction with local microstructure, e.g.the carbide, GB and TB. The fatigue crack growth rates fol-lowed the same trend as the fatigue lives in terms of recrys-tallization influence, which can be used to rationalize theobserved variation of fatigue life.

(5) The inferior LCF performance of RX0.1 sample was attributedto multiple crack initiation from carbides on discrete RX GBand favored crack growth in RX grain. The enhanced fatiguelife of RX0.5 was suggested to be due to compact recrystal-lized grains on the specimen surface and better crack growthresistance assisted by GB and TB etc. This RX layer has apotential benefit of thermal stability over other surfacestrengthening techniques. The results here indicate a morecomprehensive viewpoint in evaluating the effect of recrys-tallization on the mechanical performance of turbinesuperalloy.

Acknowledgements

The authors would like to thank Ms. Manqiong Xu in TsinghuaUniversity for her help in setting up the test. Financial supportsfrom the National Natural Science Foundation of China (Nos.11372151 and 51071094) are highly appreciated.

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