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Temperature dependence on tensile deformation mechanisms in a novel Nickel-based single crystal superalloy
Materials Science & Engineering A 776 (2020) 138997
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Temperature dependence on tensile deformation mechanisms in a novel Nickel-based single crystal superalloy Z.H. Tan a, b, X.G. Wang a, *, Y.L. Du a, T.F. Duan c, Y.H. Yang a, J.L. Liu a, J.D. Liu a, L. Yang b, J. G. Li a, Y.Z. Zhou a, **, X.F. Sun a, *** a b c
Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China School of Materials Science and Engineering, Shenyang University of Technology, 111 Shenliao Road, Shenyang, 110870, China Liaoning Research Institute of Light Industry, 3 Chongshan Road, Shenyang, 110030, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Single crystal superalloy Tensile property Fracture characteristic Stacking fault Deformation mechanism
The affordability has become a key element in the development of the modern aero-engines thus the design and research of low-cost single crystal superalloys are in great demand. A kind of novel Nickel-based single crystal superalloy with cost reduction was designed in this work and the temperature dependence on the microstructure modification as well as corresponding deformation mechanisms during tensile tests were systematically inves tigated. The experimental alloy exhibited a remarkable yield strength of 912 MPa but relatively poor ductility at 760 � C. At higher temperatures, an overt strain softening occurred before the tensile rupture and the fracture features were identified as dimples induced by the accumulated micro-pores. The stacking faults shearing mechanism prevailed at room temperature and there presented two types of stacking faults in the γ0 precipitates. Both decomposition and cross-slip of the a/2 <101> superdislocation were observed at 760 � C while the deformation mechanism was controlled by APB-coupled dislocation pairs shearing the γ0 phase at 980 � C. With temperature increasing to 1100 � C and 1120 � C, the amount of shearing dislocation pairs decreased dramatically, besides, the interfacial dislocation networks and rafted γ/γ0 structures were formed. The degradation of me chanical properties was considerably slight from 1100 � C to 1120 � C, however, three primary microstructure modifications were emphasized.
1. Introduction Superalloys, especially nickel-based single crystal superalloys have been widely used to manufacture high-pressure turbine blades in aeroengines due to their outstanding mechanical properties and corrosion resistance. The remarkable tensile strength and creep resistance under elevated temperatures of these single crystal alloys originate from the inherent γ/γ0 two-phase microstructure, i.e., the cubic γ0 phase with L12 structure coherently being embedded in the γ matrix [1–4]. During recent decades, quantities of refractory elements like Mo, W, Cr have been added to single crystal alloys to further enhance their mechanical properties. Particularly, Re has become the representative element of second and third generation single crystal superalloys for the consider able strengthening effects it can offer [5–7]. Hence alloying has become the primary method to enhance the comprehensive properties of the
alloys. However, the increasing addition of alloying elements and especially, strategic rare element Re has definitely brought about the rising costs and higher density of the single crystal alloys which conse quently limits their extensive applications. Therefore, the design and development of novel affordable single crystal superalloys have attrac ted great attentions. According to the strategy of rhenium reduction, Fink et al. [8] had developed a kind of second generation single crystal alloy with lower Re addition, which could maintain the creep property and oxidation resistance of Ren� e N5 (3 wt% Re). Han et al. [9] sys tematically studied the high cycle fatigue behavior of a rhenium free second generation single crystal alloy, the results showed that the fa tigue life decreased with the increase of stress under the condition of 900 � C. Nevertheless, the design and research work about low-cost third generation single crystal alloys remains rare. Recently, Institute of Metal Research, Chinese Academy of Science
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (X.G. Wang), [email protected] (Y.Z. Zhou), [email protected] (X.F. Sun). https://doi.org/10.1016/j.msea.2020.138997 Received 30 October 2019; Received in revised form 19 January 2020; Accepted 21 January 2020 Available online 24 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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had designed a novel Nickel-based single crystal alloy with only 3 wt% Re addition. This novel alloy was developed by the combination of composition optimization and microstructure modification. The addi tion of W and Mo were reasonably increased compared to Ren�e N6 and CMSX-10 alloys and served as the substitute of Re [10,11]. The Co content was chosen at a middle level of 8 wt% to balance the solution effect and microstructural stability. Moreover, the contents of Al and Ta which contribute to the formation of γ0 phase were chosen at the average level. On the other hand, the microstructure modification was empha sized in the present work. A similar γ/γ0 microstructure to that of Ren�e N6 alloy was obtained by the optimization of heat treatment regimes. In addition, the smelting technology was improved to minimize the defects as possible. The alloy exhibited similar creep properties with Ren�e N6 (5.4 wt% Re) under various creep conditions, while the total cost of the alloy had come down by nearly half [11–13]. Tensile property is considered as a significant indicator to charac terize the comprehensive property of single crystal superalloys and can provide further guidance for creep and thermal mechanical fatigue. The tensile behaviors of first and second generation single crystal alloys have been investigated in detail [14–17]. It was generally acknowledged that the a/2 <101> dislocations cutting into the γ0 precipitations was the dominant mechanism at room temperature and the dislocations began to decompose at middle temperature range [14,15]. Thus the a/3 <112> superpartials shearing mechanism appeared and the stacking faults could be observed. At elevated temperatures, the dislocation gliding along the γ/γ0 interface and the Orowan by-pass mechanism became common. In Wang et al.’ work [18] about fourth generation single crystal alloy, however, the stacking faults defects rather than parallel slip lines were observed at room temperature. Additionally, although the dislocation gliding and by-passing mechanisms are common at elevated temperatures, the mechanism of paired dislocations shearing the γ0 phase cannot be ignored [19]. In general, with the continuous optimization of alloy composition, new features have emerged during tensile deformation of advanced single crystal alloys at different tem peratures. It has the necessity to clarify the basic tensile deformation mechanisms of these alloys to make them function better during the service. In the present work, tensile tests from room temperature to elevated temperatures were carried out on the above novel single crystal alloy to elucidate the temperature dependence of tensile deformation mecha nism. Particularly, 1120 � C tensile test was conducted in this study to figure out the mechanical degradation of the alloy under ultimate ser vice condition.
After fully heat treated, standard cylindrical specimens with diam eter of 5 mm were cut from single crystal bars and machined for me chanical testing. The tensile tests were carried out at the temperatures of RT, 760 � C, 980 � C, 1100 � C and 1120 � C along the [001] orientation to rupture by using AG-250KNE machines. The constant strain rate was kept at 2.4 � 10 4 s 1 during the tests. The experimental metallographic specimens were grinded and polished mechanically and then etched with 4 g CuSO4þ10 ml HClþ20 ml H2O. The microstructure and fracture surfaces were examined by INSPECT F50 and TESCAN MIRA3 scanning electron microscope (SEM). After tensile tests, thin foils with the thickness of 600 μm for trans mission electron microscopy (TEM) observation were cut perpendicular to the [001] orientation and approximately 5–7 mm away from the fracture surfaces. These foils were mechanically ground to 40 μm and then electrochemically thinned in a solution of 10% perchloric acid and 90% ethanol at 20 � C by using a twin jet polisher. Ion milling process was eventually carried out at the angle of 8� and 4� successively for better observation. A JEM 2100 transmission electron microscope was used in the present work to examine the deformed structure of the experimental specimens. 3. Results 3.1. γ/γ0 microstructure Fig. 1 shows the γ/γ0 microstructure of the experimental alloy after standard heat treatments, the element segregation between dendritic cores and interdendritic regions has been eliminated to a great degree after high temperature solution treatments. The cubic γ0 phase with the average size of 340 nm was clearly observed while the volume fraction of the precipitates was approximately 68.5%. The distribution of the γ0 phase was quite uniform and no distinct difference was found between dendritic and interdendritic areas. Moreover, the average lattice pa rameters of the γ0 phase and γ matrix were acquired via X-ray diffraction method [20]. The lattice misfit δ between the γ0 phase and γ matrix at room temperature can be obtained through Eq. (1) as following:
where aγ0 and aγ represents the lattice parameters of γ phase and γ matrix respectively, and the calculated results illustrate that the γ/γ0 lattice misfit in the experimental alloy is 0.06%. 0
3.2. Tensile behavior
2. Experimental procedures
The tensile engineering strain-stress curves of the experimental sin gle crystal alloy at various typical temperatures are displayed in Fig. 2, the inset magnified pictures are presented to determine the existence of serrated flow at each temperature. Obviously, the experimental alloy demonstrated various tensile behavior at different temperatures and especially, after the occurrence of yield, the flow stress exhibited distinctive characteristic at each typical temperature. A well-defined upper yield point was observed at room temperature and it was fol lowed by the drop of flow stress to a steady platform which indicated that the work hardening effect was relatively weak. Note that a slight serrated flow was then seen in the strain-stress curve (magnified im ages). A strong work hardening took place after yield at 760 � C, addi tionally, the plastic instability which was related to dynamic strain aging became rather serious. Furthermore, it was clear that double yield phenomenon appeared at this temperature which implied the successive initiation of two kinds of deformation mechanisms. As the test temper ature increased to 980 � C, there presented only mild work hardening and the PLC effect disappeared. The flow stresses demonstrated similar characteristic after yield behavior at 1100 � C and 1120 � C, the overt softening behavior was observed before the rupture of the specimens. Similarly, no distinct PLC effect was found in the magnified tensile
Table 1 lists the nominal and measured chemical compositions of the experimental alloy. The pure raw materials were smelt in a vacuum induction furnace to obtain the master alloys and the single crystal bars with the diameter of 16 mm were prepared by means of selecting crystal method in the directional solidification furnace ZGD -15. Only single crystal rods with orientations within 8� deviating from [001] were selected in this work. The incipient melting point of the experimental alloy was deter mined via metallurgical analysis method and the standard heat treat ment regimes were as follows: 1330 � C � 1 h þ 1335 � C � 3 h þ 1345 � C � 4 h, air cooling (AC) →1100 � C � 6 h, AC→870 � C � 24 h, AC. Table 1 Chemical composition of the experimental single crystal superalloy (wt.%). Element
Al
Co
Mo
W þ Cr þ Ta
Re
Ni
Nominal Measured
5.8 5.74
8.0 8.03
2.0 2.04
18.0 17.94
3.0 3.02
Bal Bal
(1)
δ ¼ 2(aγ0 – aγ) / (aγ0 þ aγ)
2
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Fig. 1. γ/γ0 microstructure after standard heat treatments: (a) SEM micrograph of the experimental alloy, (b) size and distribution of the γ0 phase.
Fig. 2. Tensile engineering strain-stress curves of the experimental alloy at various temperatures. Inset-magnified images of each tested temperature.
curves above 1100 � C. It was interesting to found that these serrated flows appeared and vanished at different temperatures, which was closely related to the distinctive microstructure and deformation char acteristic at each temperature. The detailed formation mechanism would be discussed in the later chapter. Fig. 3 illustrates the variation of the tensile strength, elongation and reduction of area in the experimental alloy at various temperatures. The yield strength of the alloy increased gradually to a maximum of 912 MPa
from room temperature to 760 � C and then it dropped dramatically with further increase of test temperatures. It was noteworthy that the degradation of the yield strength from 1100 � C to 1120 � C was quite modest which guaranteed the mechanical properties of the alloy under ultimate temperatures. The ultimate tensile strength also rocketed to the peak at 760 � C and after that, it demonstrated similar trend to the yield strength as the temperature increased. Nevertheless, the percentage of elongation in the experimental alloy presented the completely opposite 3
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Fig. 3. Tensile properties of the experimental alloy at different temperatures. (a) Yield strengthen and ultimate tensile strength, (b) Elongation and reduction of area.
Fig. 4. Fracture surface of the experimental alloy at various temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. Inset-fracture features at higher magnification. 4
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trend. The elongation nearly remained the same from room temperature to 760 � C and subsequently, it increased to a great extent with tem perature increasing. The reduction of area reached its minimum at 760 � C, and what was interesting was that another valley value appeared between 980 � C and 1120 � C. However, the reduction of area and elongation displayed similar tendency at different temperatures in general.
cracks were clearly observed and several micro-pores were also found in the fracture surface. A macroscopically sharp and wedge-shaped frac ture was seen at 760 � C, and the fracture mechanism was defined to be pure shear fracture. As the test temperature increased, the plastic deformation capacity of the γ matrix was promoted and consequently, the dislocation gliding in the (111) plane became rather frequent compared to room temperature. The considerable plastic deformation had led to the separation of slip planes, thereby causing the fracture of specimen ultimately. Note that the micro-pores still existed at 760 � C. The fracture surface demonstrated mixed cleavage and dimples char acteristic at 980 � C. These dimples were formed by shrinkage porosity while the cleavage features presented at the edge of the surface. A distinct boundary between dimples and cleavage could be observed on the fracture surface (seen in higher magnification image). As the tem perature increased to 1100 � C, the spongy fracture surface which composed of larger dimples and tear ridges was observed, hence the fracture mechanism turned into ductile fracture. The fracture surface at
3.3. Fracture characteristic The temperature dependence on the fracture surface characteristic was shown in Fig. 4. At room temperature, the fracture surface exhibited elliptical morphology macroscopically and there presented some irreg ular fluctuation in the central of the surface. Plenty of river patterns could be observed at the edge of the specimen, therefore, the fracture mechanism was identified as cleavage. At higher magnification, cleav age steps with high parallelism which derived from propagation of the
Fig. 5. Longitudinal morphologies of the experimental alloy at various temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. 5
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1120 � C displayed similar features to that at 1100 � C, however, more micro-pores were found in the surface. These micro-holes nucleated and grew during the tensile process, and they would further aggregate and lead to the rupture. Fig. 5 shows the longitudinal morphologies of the experimental al loys after tensile fracture at various temperatures. The γ0 phase with cubic shape was observed at room temperature, and there presented a typical slip line at the angle of 45� along the (001) orientation. The gliding of a/2<110> dislocation on the (111) plane was considered as the dominant deformation mechanism of single crystal alloys at room temperature [21,22], thus the above slip line was inferred to be gener ated from (111) plane. No evident changes in the size and shape of γ0 phase were observed at 760 � C. However, more slip lines were found to cross γ/γ0 structures and most of them were in the same directions. The dislocation slipping on the (111) plane is still believed as the main deformation mode, however, it is hard to define whether the other slip system was activated under this condition, further TEM observation will be discussed. At 980 � C, the γ/γ0 microstructure was subjected to a serious distortion in the crossed slip lines, which manifested that more than one set of slip systems had been activated. With temperature increasing to 1100 � C, the slip bands disappeared and the rafted γ0 phase
and γ channel were seen in the specimen. The rafted structures were also prevalent in the specimen fractured at 1120 � C, however, the lamellar structures were found to be increased in thickness, which could be rationalized by the accelerated element diffusions under higher tem perature condition. Further discussion is required to explain the tem perature dependence of microstructure evolution in the experimental alloy after tensile rupture. 3.4. Deformation microstructure Bright TEM images of the deformation microstructure after tensile fracture at various temperatures are displayed in Fig. 6. At first glance, the deformed microstructure at room temperature was occupied by the miniative stacking faults while the dislocation pileup was considerably serious in the γ channels. Note that the stacking faults were generally located in the same [011] directions, indicating that only one primary slip system was launched. The above inference conforms to the SEM observation. After detailed inspection, however, several short disloca tions were observed, as shown by red triangles. These dislocations were originated from the complex dislocation interactions in tangled γ matrix and then sheared into the γ0 phase. Moreover, no dislocation cross-
Fig. 6. TEM micrographs of the experimental alloy after tensile fracture at different temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. 6
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slipping or climbing was observed at room temperature, the dislocations mainly glided on the typical (111) plane and were not hindered by other slip systems, which could rationalize the steady trend of strength vari ation in tensile curve. For specimen at 760 � C, larger stacking faults were observed in the γ0 phase and the occurrence of dislocations became more frequent compared to room temperature. The stacking faults presented two extension directions in the γ0 phase thus it could be inferred that the other slip system had been activated. Both straight and bend dislocations were observed at 760 � C, which provided evidence for dislocation crossslipping. When the temperature reached 980 � C, the most striking characteristic was the diminishing of the stacking faults. Moreover, plenty of curved dislocations with high energy were clearly seen to cut through the whole γ matrix into γ0 phase. Hence, the corresponding tensile deformation mechanism was believed to change compared with lower temperature conditions. With temperature increasing to 1100 � C, the γ0 phase began to lost its cubic shape to some extent, and the other significant feature was the presence of γ/γ0 interfacial dislocation networks. These dislocation networks approximately lay in a direction vertical to [010] and were similar to the configurations after creep rupture under elevated tem peratures [23,24]. Furthermore, long superdislocation pairs containing anti-phase boundary (APB) were seen to totally truncate through the γ0 precipitates, as marked by red triangles in Fig. 6 (d). Note that the initial superdislocation pairs appeared at lower temperatures in this work, nevertheless, they were generally considered to prevail under elevated temperature conditions. At 1120 � C, the interfacial dislocation networks became more regular and compact compared to that at 1100 � C, while the cubic γ0 precipitates turned into distinct spherical shape. Moreover, considering the higher thermal activation at 1120 � C, the coarsening behavior of γ0 phase in the horizontal direction became more serious and some of the adjacent particles might be linked together, which was consistent with the previous SEM observations.
rationalized by the addition of Re and Co. According to previous work [26–28], these planar faults were obtained via different dislocation re actions, and the stacking faults in γ channel were formed by the following way: a/2 <101> → a/6 <112> þ a/6 <121> þ SF in γ matrix
(2)
During tensile deformation, the interfacial dislocation a/2 <101> is generally considered to decompose at γ/γ0 interface, and the two a/6 <112> Shockley dislocations can be separated constantly driving by the flow stress and γ/γ0 misfit stress, thus producing a stacking fault in γ matrix. For stacking faults in the γ0 phase, however, they were basically generated in the process of a/3 <121> superpartials shearing into the γ0 phase, after the decomposition of a/2 <101> dislocations. Furthermore, the a/6 <112> Shockley dislocations were left behind at γ/γ0 interface. Detailed reaction equation is as following: a/2 <101> → a/6 <112> þ a/3 <121> þ SF in γ0 phase
(3)
From equation (3), it could be deduced that the a/6 <112> Shockley superpartials would currently not cut into the γ0 phase, or the atomic misarrangement would be replaced by APB with higher energy. What was interesting was that the second type of stacking fault was observed in the specimen, as marked by circle in Fig. 7. These configurations were also known as stacking fault loop in Refs. [19,25], which were formed in the following way: a/2 <101> → a/3 <112> þ a/3 <121> þ SF in γ0 phase
(4)
Nevertheless, in view of the relatively low thermal activation and misfit stress at room temperature, the amount of these small stacking faults was considerably less than that of the previous type. In addition, the straight dislocations after tensile fracture at room temperature were identified as a/3 <121> superpartials, only a few a/2 <101> superdislocation pairs plus APB could be observed [21]. Actu ally, there exists a competitive relationship between SFE and APB energy during tensile deformation, and the deformation mechanism will be controlled by the one with lower energy. Based on the above observa tions, it is reasonable to deduce that the APB energy is much higher than SFE at room temperature, and the dominant deformation mechanism in the experimental alloy is stacking faults shearing the γ0 precipitates.
4. Discussion 4.1. Deformation mechanism at RT As mentioned above, the stacking faults were identified as the dominant characteristic of the specimens fractured at room temperature and 760 � C, especially room temperature. Fig. 7 illustrates the magnified configurations of the stacking faults, it was clear that they were pre sented in both γ0 phase and γ channel by the comparison of bright and dark field images. Essentially, the presence of stacking faults in γ matrix was previously seen in ruthenium (Ru)-containing single crystal super alloys during tensile tests at room temperature [25], as a consequence of their inherent lower stacking fault energy (SFE) caused by Ru addition. In the present work, however, the experimental alloy without Ru is also inferred to have relatively low SFE at room temperature, which might be
4.2. Deformation mechanism at intermediate temperature At 760 � C, the stacking faults still prevailed in the microstructure and the detailed features were shown in Fig. 8. Different from the charac teristic under room temperature condition, only large stacking faults with superpartials (marked by rectangle) which developed from reac tion (3) were observed. Note that the intersections of these planar faults (marked by circle) were found which further confirmed the activation of multiple slip systems. Besides, as mentioned above, the dislocations
Fig. 7. Stacking faults in the γ0 phase and γ matrix at room temperature: (a) bright field image, (b) dark field image. 7
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4.3. Deformation mechanism at elevated temperatures As the temperature increased to 980 � C, the order degree in the γ0 precipitates would decrease to considerable extent and the APB energy was inferred to decline extremely. Hence the decomposition of a/2 <101> dislocation at γ/γ0 interface is not considered as the optimal option and the stacking faults can be completely replaced by APBcoupled dislocation pairs, as marked in Fig. 9 (a). Moreover, the other two significant changes related to the microstructure should be emphasized. Firstly, we all know the lattice misfit between the γ0 phase and matrix will become more negative with the increasing temperature in single crystal alloys [19,33,34]. Thus the initial irregular dislocation networks were formed at γ/γ0 interface to help release the increasing coherency strains (marked by quadrangle). Secondly, the strength of γ0 precipitate was believed to decrease when the temperature exceeded approximately 900 � C [3,35]. Very limited hindrance to dislocation shearing behaviors could be expected by these irregular dislocation networks and softened γ0 phase. Consequently, with the onset of dislo cation pairs crossing interfacial dislocation networks and shearing the γ0 precipitates, a strain softening can be found (after the point of UTS) in tensile curves. According to the above experimental results, a critical transformation temperature of tensile deformation mechanism regarding stacking fault shearing and APB shearing mechanisms is inferred to exist between 760 � C and 980 � C, further thorough tests are required in the future works. At 1100 � C, the dislocation shearing the γ0 phase became less prev alent, and only several long superdislocation pairs with APB were observed, which were mainly defined as a/2 <101> pairs. Note that the evolution of dislocation networks was promoted by the much reduced γ/γ0 lattice misfit while the edge of the γ0 precipitates was found to dissolve into the matrix. The rafted structures had been formed and the deformation mechanism was initially controlled by Orowan by-passing. Subsequently, as the flow stress continued to increase, the paired dis locations containing APB were deduced to cut through the interfacial dislocation networks and γ0 phase. The tensile deformation mechanism and microstructure character istic of the fractured specimen at 1120 � C are basically similar to that at 1100 � C, however, three primary modifications need to be pointed out. Firstly, driven by the promoted thermal activation and high stress as well as the further reduction in γ/γ0 lattice misfit, the perfection process of interfacial dislocation networks had been completed within a short tensile period. Secondly, to reach a thermodynamic equilibrium at higher temperature, the rafting and dissolving behaviors of the γ0 pre cipitates were accelerated compared to 1100 � C. What is the most noteworthy are the variations in the amount and configuration of the cutting superdislocation pairs. Unlike the long straight dislocation pairs at 1100 � C, the shearing dislocations had significantly declined in quantity and they were short in length. Obviously, the dislocation pairs have been impeded from truncating the γ0 phase by the denser interfacial
Fig. 8. Configurations of typical stacking faults and dislocations after tensile fracture at 760 � C.
amount increased to a certain degree while the cross-slip motion was seen. The marked dislocations in Fig. 8 implied that the a/2 <101> superdislocation had cross-slipped from {111} planes to {001} planes, thereby hindering the further slipping of the rest dislocations in {111} planes. The above dislocation configuration is known as Kear-Wilsdorf lock [29], which greatly contributes to the work hardening during tensile deformation at 760 � C. An increasing number of dislocation pairs was also observed in the experimental alloy, thus the deformation mechanism was turned to the mix of stacking fault and APB shearing the γ0 phase, which also helped explain the presence of double yield plateaus in tensile curves. The other striking characteristic at 760 � C is the occurrence of PLC phenomenon after yield behavior in the specimen. Essentially, the serrated plastic flow is commonly found in many Ni-based superalloys within a certain temperature range, and it is believed to be influenced by multiple factors like diffusion rates of solute atoms, activation energy, SFE, etc. [30–32]. In the present work, the discontinuous plastic defor mation began to appear at room temperature and became serious at 760 � C, however, it disappeared above 980 � C. Therefore, it is rational to speculate that the above phenomenon is closely correlated with the presence of stacking faults in the microstructure. A dynamic pinning and unpinning processes existed between the solute atoms and the stacking faults with mobile dislocations, which brought about the typical mode of plastic flow in strain-stress curves. Moreover, affected by the higher activation energy at 760 � C, the specimen exhibited more intensive PLC effect compared to room temperature.
Fig. 9. Typical dislocation configurations after tensile fracture at elevated temperatures: (a) 980 � C, (b) 1120 � C. 8
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dislocation networks while the rafted structures have also exerted ob stacles to the dislocation climbing. Interestingly, a kind of curved dislocation pairs (region A) was observed at elevated temperatures, additionally, it was much more common at 1120 � C, as shown in Fig. 9 (b). It could be identified as a <010> superdislocation, which was generated by the reaction of a/2 <011> dislocations with different Burgers vectors in the matrix [36]. The above dislocation configuration was previously seen in some single crystal alloys after creep rupture at elevated temperatures, and it was inferred to move in a combined way of sliding and climbing [23,37]. Besides, researchers found that the minimum creep rates of the alloys could be lowered down by these dislocations with inferior mobility [38]. Nevertheless, no significant improvement related to the alloy mechan ical properties is expected in this study, considering the much higher local stress during tensile deformation. When the flow stress reaches a threshold value, lots of paired dislocations with APB can easily cut into the γ0 phase, hence the softening stage in stain-stress curves has become definitely evident under elevated temperature conditions.
curation, Supervision. J.G. Li: Resources. Y.Z. Zhou: Project adminis tration, Funding acquisition. X.F. Sun: Funding acquisition, Supervision.
5. Conclusions
References
The tensile behaviors of a novel third generation single crystal su peralloy at various temperatures were studied in this work. The tem perature dependence on the fracture surfaces, microstructure and deformation mechanisms were studied in detail via scanning electron microscope and transmission electron microscope. The main conclu sions could be drawn as follows:
[1] C. Sims, N. Stoloff, W. Hagel, Superalloy ІІ, John Wiley & Sons, New York, 1987. [2] Y.B. Hu, T.S. Cao, C.Q. Cheng, L. Zhang, J. Zhao, Oxidation behavior of a singlecrystal Ni-based superalloy over the temperature range of 850 � C-950 � C in air, Appl. Surf. Sci. 484 (2019) 209–218. [3] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, 2006. [4] R.H. Wu, Z.F. Yue, M. Wang, Effect of initial γ/γ0 microstructure on creep of single crystal nickel-based superalloys: a phase-field simulation incorporating dislocation dynamics, J. Alloys Compd. 779 (2019) 326–334. [5] Q.Q. Ding, S.Z. Li, L.Q. Chen, X.D. Han, Z. Zhang, Q. Yu, J.X. Li, Re segregation at interfacial dislocation network in a nickel-based superalloy, Acta Mater. 154 (2018) 137–146. [6] F.H. Latief, K. Kakehi, Effects of Re content and crystallographic orientation on creep behavior of aluminized Ni-base single crystal superalloys, Mater. Des. 49 (2013) 485–492. [7] L. Liu, J. Meng, J.L. Liu, M.K. Zou, H.F. Zhang, X.D. Sun, Y.Z. Zhou, Influences of Re on low-cycle fatigue behaviors of single crystal superalloys at intermediate temperature, J. Mater. Sci. Technol. 35 (2019) 1917–1924. [8] P.J. Fink, J.L. Miller, D.G. Konitzer, Rhenium reduction-alloy design using an economically strategic element, JOM 62 (2010) 55–57. [9] G.M. Han, Z.X. Zhang, J.G. Ji, T. Jin, X.F. Sun, Z.Q. Hu, High cycle fatigue behavior of a nickel-based single crystal superalloy DD98M at 900 � C, Acta Metall. Sin. 48 (2012) 170–175. [10] G.L. Erickson, The Development and Application of CMSX-10, TMS, Warrendale, PA, 1996, pp. 35–44. [11] W.S. Walston, K.S. O0 Hara, E.W. Rose, Ren�e N6: Third Generation Single Crystal Superalloy, TMS, Warrendale, PA, 1996, pp. 27–34. [12] T. Hino, T. Kobayashi, Y. Koizumi, H. Harada, T. Yamagata, Development of a New Single Crystal Superalloy for Industrial Gas Turbines, TMS, Warrendale, PA, 2000, pp. 729–736. [13] P. Caron, High γ0 Solvus New Generation Nickel-Based Superalloys for Single Crystal Turbine Blade Applications, TMS, Warrendale, PA, 2000, pp. 737–746. [14] W.W. Milligan, S.D. Antolovich, The mechanisms and temperature dependence of superlattice stacking fault formation in the single-crystal superalloy PWA 1480, Metal. Trans. A. 22 (1991) 2309–2318. [15] Z.X. Shi, S.Z. Liu, J. Yu, J.R. Li, Tensile behavior of the second generation single crystal superalloy DD6, J. Iron Steel Res. Int. 22 (2015) 738–742. [16] W.P. Yang, J.R. Li, S.Z. Liu, Z.X. Shi, J.Q. Zhao, X.Q. Wang, Transverse tensile properties of a Nickel-based third generation single crystal superalloy, Rare Metal, Mater. Eng. 47 (2018) 2964–2969. [17] X.H. Xiong, D.M. Quan, P.D. Dai, Z.P. Wang, Q.X. Zhang, Z.F. Yue, Tensile behavior of nickel-base single crystal superalloy DD6, Mater. Sci. Eng., A 636 (2015) 608–612. [18] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, The effects of ruthenium additions on tensile deformation mechanisms of single crystal superalloys at different temperatures, Mater. Des. 63 (2014) 286–293. [19] G.L. Wang, J.L. Liu, J.D. Liu, X.G. Wang, Y.Z. Zhou, X.D. Sun, H.F. Zhang, T. Jin, Temperature dependence of tensile behavior and deformation microstructure of a Re-containing Ni-base single crystal superalloy, Mater. Des. 130 (2017) 131–139. [20] X.P. Tan, J.L. Liu, X.P. Song, T. Jin, X.F. Sun, Z.Q. Hu, Measurements of γ/γ0 lattice misfit and γ0 volume fraction for a Ru-containing nickel-based single crystal superalloy, J. Mater. Sci. Technol. 27 (2011) 899–905. [21] N. Tian, S.G. Tian, H.J. Yan, D.L. Shu, S.K. Zhang, G.Q. Zhao, Deformation mechanisms and analysis of a single crystal nickel-based superalloy during tensile at room temperature, Mater. Sci. Eng., A 744 (2019) 154–162. [22] P. Zhang, Y. Yuan, B. Li, G. Yang, X. Song, Investigation of tensile deformation mechanisms at room temperature in a new Ni-based single crystal superalloy, Phil. Mag. Lett. 96 (2016) 238–245.
Acknowledgements The financial supports provided by National Key R&D Program of China under Grant No. 2017YFA0700704, National Science and Tech nology Major Project under Grant No. 2017-VI-0002-0072, Natural Science Foundation of Liaoning Province under Grant No. 20170520038, National Natural Science Foundation of China (NSFC) under Grant Nos. 51601192 and 51671188, State Key Lab of Advanced Metals and Materials Open Fund under Grant No.2018-Z07 and Youth Innovation Promotion Association, Chinese Academy of Sciences for carrying out this work are gratefully acknowledged. The authors are grateful to J.M. Liu, X.L. Wang (Institute of Metal Research) and S.N. Nie (Shenyang University of Technology) for the indispensable assistance in TEM observation.
(1) The flow stress of the experimental alloy demonstrated distinc tive characteristic after yield. A strong work hardening was observed at 760 � C while the softening behavior presented at elevated temperatures. (2) The fracture mode was identified as cleavage and pure shearing at room temperature and 760 � C respectively. The fracture sur face exhibited a mix of dimples and cleavage at 980 � C, and the fracture mechanism turned into ductile features with tempera ture increasing to 1100 � C and 1120 � C. (3) Stacking faults-related microstructure was clearly observed at room temperature and 760 � C, however, it disappeared at 980 � C. The interfacial dislocation networks were formed at elevated temperatures to help release the increasing γ/γ0 lattice strain. In addition, several modifications with regard to the microstructure were illustrated when the temperature was raised from 1100 � C to 1120 � C. (4) As the temperature increased from room temperature to 980 � C, the deformation mechanism changed from stacking faults shearing the γ0 precipitates to APB shearing the strengthening phase, additionally, these two modes were believed to coexist at 760 � C. Above 980 � C, the dislocation climbing and Orowan bypassing became common and the APB shearing mechanism was considered to predominate when the applied stress reached a critical value. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Z.H. Tan: Writing - original draft, Methodology, Investigation. X.G. Wang: Conceptualization, Project administration, Funding acquisition. Y.L. Du: Methodology, Writing - review & editing. T.F. Duan: Funding acquisition, Writing - review & editing. Y.H. Yang: Software, Visuali zation. J.L. Liu: Formal analysis. J.D. Liu: Validation. L. Yang: Data 9
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Materials Science & Engineering A 776 (2020) 138997 [31] P. Fernandez-Zelaia, B.S. Adair, V.M. Barker, S.D. Antolovich, The portevin-le chatelier effect in the Ni-based superalloy IN100, Metall. Mater. Trans. 46 (2015) 5596–5609. [32] C.Y. Cui, C.G. Tian, Y.Z. Zhou, T. Jin, X.F. Sun, Dynamic Strain Aging in Ni Base Alloys with Different Stacking Fault Energy, TMS, Warrendale, PA, 2012, pp. 715–722. [33] S. Ma, V. Seetharaman, B.S. Majumdar, CRSS of γ/γ’ phases from in situ neutron diffraction of a directionally solidified superalloy tension tested at 900 � C, Acta Mater. 56 (2008) 4102–4113. [34] G. Bruno, S. Schumacher, H.C. Pinto, C. Schulze, Measurement of the lattice misfit of the nickel-base superalloy SC16 by high-energy synchrotron radiation, Metall. Mater. Trans. 34 (2003) 193–197. [35] P. Beadmore, R.G. Davies, T.L. Johnston, On the temperature dependence of the flow stress of Nickel-base alloys, Trans. AIME. 245 (1969) 1537–1545. [36] G. Eggeler, A. Dlouhy, On the formation of <010>-dislocations in the γ0 -phase of superalloy single crystals during high temperature low stress creep, Acta Mater. 45 (1997) 4251–4262. [37] P.M. Sarosi, R. Srinivasan, G.F. Eggeler, M.V. Nathal, M.J. Mills, Observations of a <010> dislocations during the high-temperature creep of Ni-based superalloy single crystals deformed along the [001] orientation, Acta Mater. 55 (2007) 2509–2518. [38] R. Srinivasan, G.F. Eggeler, M.J. Mills, γ0 -cutting as rate-controlling recovery process during high-temperature and low-stress creep of superalloy single crystals, Acta Mater. 48 (2000) 4867–4878.
[23] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Z.Q. Hu, J.H. Do, B.G. Choi, I.S. Kim, C.Y. Jo, Dislocation motion during high-temperature low-stress creep in Ru-free and Rucontaining single-crystal superalloys, Mater. Des. 67 (2015) 543–551. [24] J.X. Zhang, T. Murakumo, Y. Koizumi, T. Kobayashi, H. Harada, Slip geometry of dislocations related to cutting of the γ0 phase in a new generation single-crystal superalloy, Acta Mater. 51 (2003) 5073–5081. [25] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Tensile behaviors and deformation mechanisms of a nickel-base single crystal superalloy at different temperatures, Mater. Sci. Eng., A 598 (2014) 154–161. [26] S.G. Tian, Z. Zheng, F.S. Liang, Z. Chao, L. Chen, Creep behavior of a 4.5%-Re single crystal nickel-based superalloy at intermediate temperatures, Mater. Sci. Eng., A 543 (2012) 104–109. [27] J.L. Liu, J.J. Yu, T. Jin, X.F. Sun, H.R. Guan, Z.Q. Hu, Influence of temperature on tensile behavior and deformation mechanism of Re-containing single crystal superalloy, Trans. Nonferrous Metals Soc. China 21 (2011) 1518–1523. [28] L. Kovarik, R.R. Unocic, J. Li, P. Sarosi, C. Shen, Y. Wang, M.J. Mills, Microtwinning and other shearing mechanisms at intermediate temperatures in Nibased superalloys, Prog. Mater. Sci. 54 (2009) 839–873. [29] K.J. Hemker, M.J. Mills, W.D. Nix, An investigation of the mechanisms that control intermediate temperature creep of Ni3Al, Acta Metall. Mater. 39 (1991) 1901–1913. [30] Y.L. Cai, C.G. Tian, S.H. Fu, G.M. Han, C.Y. Cui, Q.C. Zhang, Influence of γʹ precipitates on portevin–le chatelier effect of Ni-based superalloys, Mater. Sci. Eng., A 638 (2015) 314–321.
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Temperature dependence on tensile deformation mechanisms in a novel Nickel-based single crystal superalloy Z.H. Tan a, b, X.G. Wang a, *, Y.L. Du a, T.F. Duan c, Y.H. Yang a, J.L. Liu a, J.D. Liu a, L. Yang b, J. G. Li a, Y.Z. Zhou a, **, X.F. Sun a, *** a b c
Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China School of Materials Science and Engineering, Shenyang University of Technology, 111 Shenliao Road, Shenyang, 110870, China Liaoning Research Institute of Light Industry, 3 Chongshan Road, Shenyang, 110030, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Single crystal superalloy Tensile property Fracture characteristic Stacking fault Deformation mechanism
The affordability has become a key element in the development of the modern aero-engines thus the design and research of low-cost single crystal superalloys are in great demand. A kind of novel Nickel-based single crystal superalloy with cost reduction was designed in this work and the temperature dependence on the microstructure modification as well as corresponding deformation mechanisms during tensile tests were systematically inves tigated. The experimental alloy exhibited a remarkable yield strength of 912 MPa but relatively poor ductility at 760 � C. At higher temperatures, an overt strain softening occurred before the tensile rupture and the fracture features were identified as dimples induced by the accumulated micro-pores. The stacking faults shearing mechanism prevailed at room temperature and there presented two types of stacking faults in the γ0 precipitates. Both decomposition and cross-slip of the a/2 <101> superdislocation were observed at 760 � C while the deformation mechanism was controlled by APB-coupled dislocation pairs shearing the γ0 phase at 980 � C. With temperature increasing to 1100 � C and 1120 � C, the amount of shearing dislocation pairs decreased dramatically, besides, the interfacial dislocation networks and rafted γ/γ0 structures were formed. The degradation of me chanical properties was considerably slight from 1100 � C to 1120 � C, however, three primary microstructure modifications were emphasized.
1. Introduction Superalloys, especially nickel-based single crystal superalloys have been widely used to manufacture high-pressure turbine blades in aeroengines due to their outstanding mechanical properties and corrosion resistance. The remarkable tensile strength and creep resistance under elevated temperatures of these single crystal alloys originate from the inherent γ/γ0 two-phase microstructure, i.e., the cubic γ0 phase with L12 structure coherently being embedded in the γ matrix [1–4]. During recent decades, quantities of refractory elements like Mo, W, Cr have been added to single crystal alloys to further enhance their mechanical properties. Particularly, Re has become the representative element of second and third generation single crystal superalloys for the consider able strengthening effects it can offer [5–7]. Hence alloying has become the primary method to enhance the comprehensive properties of the
alloys. However, the increasing addition of alloying elements and especially, strategic rare element Re has definitely brought about the rising costs and higher density of the single crystal alloys which conse quently limits their extensive applications. Therefore, the design and development of novel affordable single crystal superalloys have attrac ted great attentions. According to the strategy of rhenium reduction, Fink et al. [8] had developed a kind of second generation single crystal alloy with lower Re addition, which could maintain the creep property and oxidation resistance of Ren� e N5 (3 wt% Re). Han et al. [9] sys tematically studied the high cycle fatigue behavior of a rhenium free second generation single crystal alloy, the results showed that the fa tigue life decreased with the increase of stress under the condition of 900 � C. Nevertheless, the design and research work about low-cost third generation single crystal alloys remains rare. Recently, Institute of Metal Research, Chinese Academy of Science
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (X.G. Wang), [email protected] (Y.Z. Zhou), [email protected] (X.F. Sun). https://doi.org/10.1016/j.msea.2020.138997 Received 30 October 2019; Received in revised form 19 January 2020; Accepted 21 January 2020 Available online 24 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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had designed a novel Nickel-based single crystal alloy with only 3 wt% Re addition. This novel alloy was developed by the combination of composition optimization and microstructure modification. The addi tion of W and Mo were reasonably increased compared to Ren�e N6 and CMSX-10 alloys and served as the substitute of Re [10,11]. The Co content was chosen at a middle level of 8 wt% to balance the solution effect and microstructural stability. Moreover, the contents of Al and Ta which contribute to the formation of γ0 phase were chosen at the average level. On the other hand, the microstructure modification was empha sized in the present work. A similar γ/γ0 microstructure to that of Ren�e N6 alloy was obtained by the optimization of heat treatment regimes. In addition, the smelting technology was improved to minimize the defects as possible. The alloy exhibited similar creep properties with Ren�e N6 (5.4 wt% Re) under various creep conditions, while the total cost of the alloy had come down by nearly half [11–13]. Tensile property is considered as a significant indicator to charac terize the comprehensive property of single crystal superalloys and can provide further guidance for creep and thermal mechanical fatigue. The tensile behaviors of first and second generation single crystal alloys have been investigated in detail [14–17]. It was generally acknowledged that the a/2 <101> dislocations cutting into the γ0 precipitations was the dominant mechanism at room temperature and the dislocations began to decompose at middle temperature range [14,15]. Thus the a/3 <112> superpartials shearing mechanism appeared and the stacking faults could be observed. At elevated temperatures, the dislocation gliding along the γ/γ0 interface and the Orowan by-pass mechanism became common. In Wang et al.’ work [18] about fourth generation single crystal alloy, however, the stacking faults defects rather than parallel slip lines were observed at room temperature. Additionally, although the dislocation gliding and by-passing mechanisms are common at elevated temperatures, the mechanism of paired dislocations shearing the γ0 phase cannot be ignored [19]. In general, with the continuous optimization of alloy composition, new features have emerged during tensile deformation of advanced single crystal alloys at different tem peratures. It has the necessity to clarify the basic tensile deformation mechanisms of these alloys to make them function better during the service. In the present work, tensile tests from room temperature to elevated temperatures were carried out on the above novel single crystal alloy to elucidate the temperature dependence of tensile deformation mecha nism. Particularly, 1120 � C tensile test was conducted in this study to figure out the mechanical degradation of the alloy under ultimate ser vice condition.
After fully heat treated, standard cylindrical specimens with diam eter of 5 mm were cut from single crystal bars and machined for me chanical testing. The tensile tests were carried out at the temperatures of RT, 760 � C, 980 � C, 1100 � C and 1120 � C along the [001] orientation to rupture by using AG-250KNE machines. The constant strain rate was kept at 2.4 � 10 4 s 1 during the tests. The experimental metallographic specimens were grinded and polished mechanically and then etched with 4 g CuSO4þ10 ml HClþ20 ml H2O. The microstructure and fracture surfaces were examined by INSPECT F50 and TESCAN MIRA3 scanning electron microscope (SEM). After tensile tests, thin foils with the thickness of 600 μm for trans mission electron microscopy (TEM) observation were cut perpendicular to the [001] orientation and approximately 5–7 mm away from the fracture surfaces. These foils were mechanically ground to 40 μm and then electrochemically thinned in a solution of 10% perchloric acid and 90% ethanol at 20 � C by using a twin jet polisher. Ion milling process was eventually carried out at the angle of 8� and 4� successively for better observation. A JEM 2100 transmission electron microscope was used in the present work to examine the deformed structure of the experimental specimens. 3. Results 3.1. γ/γ0 microstructure Fig. 1 shows the γ/γ0 microstructure of the experimental alloy after standard heat treatments, the element segregation between dendritic cores and interdendritic regions has been eliminated to a great degree after high temperature solution treatments. The cubic γ0 phase with the average size of 340 nm was clearly observed while the volume fraction of the precipitates was approximately 68.5%. The distribution of the γ0 phase was quite uniform and no distinct difference was found between dendritic and interdendritic areas. Moreover, the average lattice pa rameters of the γ0 phase and γ matrix were acquired via X-ray diffraction method [20]. The lattice misfit δ between the γ0 phase and γ matrix at room temperature can be obtained through Eq. (1) as following:
where aγ0 and aγ represents the lattice parameters of γ phase and γ matrix respectively, and the calculated results illustrate that the γ/γ0 lattice misfit in the experimental alloy is 0.06%. 0
3.2. Tensile behavior
2. Experimental procedures
The tensile engineering strain-stress curves of the experimental sin gle crystal alloy at various typical temperatures are displayed in Fig. 2, the inset magnified pictures are presented to determine the existence of serrated flow at each temperature. Obviously, the experimental alloy demonstrated various tensile behavior at different temperatures and especially, after the occurrence of yield, the flow stress exhibited distinctive characteristic at each typical temperature. A well-defined upper yield point was observed at room temperature and it was fol lowed by the drop of flow stress to a steady platform which indicated that the work hardening effect was relatively weak. Note that a slight serrated flow was then seen in the strain-stress curve (magnified im ages). A strong work hardening took place after yield at 760 � C, addi tionally, the plastic instability which was related to dynamic strain aging became rather serious. Furthermore, it was clear that double yield phenomenon appeared at this temperature which implied the successive initiation of two kinds of deformation mechanisms. As the test temper ature increased to 980 � C, there presented only mild work hardening and the PLC effect disappeared. The flow stresses demonstrated similar characteristic after yield behavior at 1100 � C and 1120 � C, the overt softening behavior was observed before the rupture of the specimens. Similarly, no distinct PLC effect was found in the magnified tensile
Table 1 lists the nominal and measured chemical compositions of the experimental alloy. The pure raw materials were smelt in a vacuum induction furnace to obtain the master alloys and the single crystal bars with the diameter of 16 mm were prepared by means of selecting crystal method in the directional solidification furnace ZGD -15. Only single crystal rods with orientations within 8� deviating from [001] were selected in this work. The incipient melting point of the experimental alloy was deter mined via metallurgical analysis method and the standard heat treat ment regimes were as follows: 1330 � C � 1 h þ 1335 � C � 3 h þ 1345 � C � 4 h, air cooling (AC) →1100 � C � 6 h, AC→870 � C � 24 h, AC. Table 1 Chemical composition of the experimental single crystal superalloy (wt.%). Element
Al
Co
Mo
W þ Cr þ Ta
Re
Ni
Nominal Measured
5.8 5.74
8.0 8.03
2.0 2.04
18.0 17.94
3.0 3.02
Bal Bal
(1)
δ ¼ 2(aγ0 – aγ) / (aγ0 þ aγ)
2
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Fig. 1. γ/γ0 microstructure after standard heat treatments: (a) SEM micrograph of the experimental alloy, (b) size and distribution of the γ0 phase.
Fig. 2. Tensile engineering strain-stress curves of the experimental alloy at various temperatures. Inset-magnified images of each tested temperature.
curves above 1100 � C. It was interesting to found that these serrated flows appeared and vanished at different temperatures, which was closely related to the distinctive microstructure and deformation char acteristic at each temperature. The detailed formation mechanism would be discussed in the later chapter. Fig. 3 illustrates the variation of the tensile strength, elongation and reduction of area in the experimental alloy at various temperatures. The yield strength of the alloy increased gradually to a maximum of 912 MPa
from room temperature to 760 � C and then it dropped dramatically with further increase of test temperatures. It was noteworthy that the degradation of the yield strength from 1100 � C to 1120 � C was quite modest which guaranteed the mechanical properties of the alloy under ultimate temperatures. The ultimate tensile strength also rocketed to the peak at 760 � C and after that, it demonstrated similar trend to the yield strength as the temperature increased. Nevertheless, the percentage of elongation in the experimental alloy presented the completely opposite 3
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Fig. 3. Tensile properties of the experimental alloy at different temperatures. (a) Yield strengthen and ultimate tensile strength, (b) Elongation and reduction of area.
Fig. 4. Fracture surface of the experimental alloy at various temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. Inset-fracture features at higher magnification. 4
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trend. The elongation nearly remained the same from room temperature to 760 � C and subsequently, it increased to a great extent with tem perature increasing. The reduction of area reached its minimum at 760 � C, and what was interesting was that another valley value appeared between 980 � C and 1120 � C. However, the reduction of area and elongation displayed similar tendency at different temperatures in general.
cracks were clearly observed and several micro-pores were also found in the fracture surface. A macroscopically sharp and wedge-shaped frac ture was seen at 760 � C, and the fracture mechanism was defined to be pure shear fracture. As the test temperature increased, the plastic deformation capacity of the γ matrix was promoted and consequently, the dislocation gliding in the (111) plane became rather frequent compared to room temperature. The considerable plastic deformation had led to the separation of slip planes, thereby causing the fracture of specimen ultimately. Note that the micro-pores still existed at 760 � C. The fracture surface demonstrated mixed cleavage and dimples char acteristic at 980 � C. These dimples were formed by shrinkage porosity while the cleavage features presented at the edge of the surface. A distinct boundary between dimples and cleavage could be observed on the fracture surface (seen in higher magnification image). As the tem perature increased to 1100 � C, the spongy fracture surface which composed of larger dimples and tear ridges was observed, hence the fracture mechanism turned into ductile fracture. The fracture surface at
3.3. Fracture characteristic The temperature dependence on the fracture surface characteristic was shown in Fig. 4. At room temperature, the fracture surface exhibited elliptical morphology macroscopically and there presented some irreg ular fluctuation in the central of the surface. Plenty of river patterns could be observed at the edge of the specimen, therefore, the fracture mechanism was identified as cleavage. At higher magnification, cleav age steps with high parallelism which derived from propagation of the
Fig. 5. Longitudinal morphologies of the experimental alloy at various temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. 5
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1120 � C displayed similar features to that at 1100 � C, however, more micro-pores were found in the surface. These micro-holes nucleated and grew during the tensile process, and they would further aggregate and lead to the rupture. Fig. 5 shows the longitudinal morphologies of the experimental al loys after tensile fracture at various temperatures. The γ0 phase with cubic shape was observed at room temperature, and there presented a typical slip line at the angle of 45� along the (001) orientation. The gliding of a/2<110> dislocation on the (111) plane was considered as the dominant deformation mechanism of single crystal alloys at room temperature [21,22], thus the above slip line was inferred to be gener ated from (111) plane. No evident changes in the size and shape of γ0 phase were observed at 760 � C. However, more slip lines were found to cross γ/γ0 structures and most of them were in the same directions. The dislocation slipping on the (111) plane is still believed as the main deformation mode, however, it is hard to define whether the other slip system was activated under this condition, further TEM observation will be discussed. At 980 � C, the γ/γ0 microstructure was subjected to a serious distortion in the crossed slip lines, which manifested that more than one set of slip systems had been activated. With temperature increasing to 1100 � C, the slip bands disappeared and the rafted γ0 phase
and γ channel were seen in the specimen. The rafted structures were also prevalent in the specimen fractured at 1120 � C, however, the lamellar structures were found to be increased in thickness, which could be rationalized by the accelerated element diffusions under higher tem perature condition. Further discussion is required to explain the tem perature dependence of microstructure evolution in the experimental alloy after tensile rupture. 3.4. Deformation microstructure Bright TEM images of the deformation microstructure after tensile fracture at various temperatures are displayed in Fig. 6. At first glance, the deformed microstructure at room temperature was occupied by the miniative stacking faults while the dislocation pileup was considerably serious in the γ channels. Note that the stacking faults were generally located in the same [011] directions, indicating that only one primary slip system was launched. The above inference conforms to the SEM observation. After detailed inspection, however, several short disloca tions were observed, as shown by red triangles. These dislocations were originated from the complex dislocation interactions in tangled γ matrix and then sheared into the γ0 phase. Moreover, no dislocation cross-
Fig. 6. TEM micrographs of the experimental alloy after tensile fracture at different temperatures: (a) RT, (b) 760 � C, (c) 980 � C, (d) 1100 � C, (e) 1120 � C. 6
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slipping or climbing was observed at room temperature, the dislocations mainly glided on the typical (111) plane and were not hindered by other slip systems, which could rationalize the steady trend of strength vari ation in tensile curve. For specimen at 760 � C, larger stacking faults were observed in the γ0 phase and the occurrence of dislocations became more frequent compared to room temperature. The stacking faults presented two extension directions in the γ0 phase thus it could be inferred that the other slip system had been activated. Both straight and bend dislocations were observed at 760 � C, which provided evidence for dislocation crossslipping. When the temperature reached 980 � C, the most striking characteristic was the diminishing of the stacking faults. Moreover, plenty of curved dislocations with high energy were clearly seen to cut through the whole γ matrix into γ0 phase. Hence, the corresponding tensile deformation mechanism was believed to change compared with lower temperature conditions. With temperature increasing to 1100 � C, the γ0 phase began to lost its cubic shape to some extent, and the other significant feature was the presence of γ/γ0 interfacial dislocation networks. These dislocation networks approximately lay in a direction vertical to [010] and were similar to the configurations after creep rupture under elevated tem peratures [23,24]. Furthermore, long superdislocation pairs containing anti-phase boundary (APB) were seen to totally truncate through the γ0 precipitates, as marked by red triangles in Fig. 6 (d). Note that the initial superdislocation pairs appeared at lower temperatures in this work, nevertheless, they were generally considered to prevail under elevated temperature conditions. At 1120 � C, the interfacial dislocation networks became more regular and compact compared to that at 1100 � C, while the cubic γ0 precipitates turned into distinct spherical shape. Moreover, considering the higher thermal activation at 1120 � C, the coarsening behavior of γ0 phase in the horizontal direction became more serious and some of the adjacent particles might be linked together, which was consistent with the previous SEM observations.
rationalized by the addition of Re and Co. According to previous work [26–28], these planar faults were obtained via different dislocation re actions, and the stacking faults in γ channel were formed by the following way: a/2 <101> → a/6 <112> þ a/6 <121> þ SF in γ matrix
(2)
During tensile deformation, the interfacial dislocation a/2 <101> is generally considered to decompose at γ/γ0 interface, and the two a/6 <112> Shockley dislocations can be separated constantly driving by the flow stress and γ/γ0 misfit stress, thus producing a stacking fault in γ matrix. For stacking faults in the γ0 phase, however, they were basically generated in the process of a/3 <121> superpartials shearing into the γ0 phase, after the decomposition of a/2 <101> dislocations. Furthermore, the a/6 <112> Shockley dislocations were left behind at γ/γ0 interface. Detailed reaction equation is as following: a/2 <101> → a/6 <112> þ a/3 <121> þ SF in γ0 phase
(3)
From equation (3), it could be deduced that the a/6 <112> Shockley superpartials would currently not cut into the γ0 phase, or the atomic misarrangement would be replaced by APB with higher energy. What was interesting was that the second type of stacking fault was observed in the specimen, as marked by circle in Fig. 7. These configurations were also known as stacking fault loop in Refs. [19,25], which were formed in the following way: a/2 <101> → a/3 <112> þ a/3 <121> þ SF in γ0 phase
(4)
Nevertheless, in view of the relatively low thermal activation and misfit stress at room temperature, the amount of these small stacking faults was considerably less than that of the previous type. In addition, the straight dislocations after tensile fracture at room temperature were identified as a/3 <121> superpartials, only a few a/2 <101> superdislocation pairs plus APB could be observed [21]. Actu ally, there exists a competitive relationship between SFE and APB energy during tensile deformation, and the deformation mechanism will be controlled by the one with lower energy. Based on the above observa tions, it is reasonable to deduce that the APB energy is much higher than SFE at room temperature, and the dominant deformation mechanism in the experimental alloy is stacking faults shearing the γ0 precipitates.
4. Discussion 4.1. Deformation mechanism at RT As mentioned above, the stacking faults were identified as the dominant characteristic of the specimens fractured at room temperature and 760 � C, especially room temperature. Fig. 7 illustrates the magnified configurations of the stacking faults, it was clear that they were pre sented in both γ0 phase and γ channel by the comparison of bright and dark field images. Essentially, the presence of stacking faults in γ matrix was previously seen in ruthenium (Ru)-containing single crystal super alloys during tensile tests at room temperature [25], as a consequence of their inherent lower stacking fault energy (SFE) caused by Ru addition. In the present work, however, the experimental alloy without Ru is also inferred to have relatively low SFE at room temperature, which might be
4.2. Deformation mechanism at intermediate temperature At 760 � C, the stacking faults still prevailed in the microstructure and the detailed features were shown in Fig. 8. Different from the charac teristic under room temperature condition, only large stacking faults with superpartials (marked by rectangle) which developed from reac tion (3) were observed. Note that the intersections of these planar faults (marked by circle) were found which further confirmed the activation of multiple slip systems. Besides, as mentioned above, the dislocations
Fig. 7. Stacking faults in the γ0 phase and γ matrix at room temperature: (a) bright field image, (b) dark field image. 7
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4.3. Deformation mechanism at elevated temperatures As the temperature increased to 980 � C, the order degree in the γ0 precipitates would decrease to considerable extent and the APB energy was inferred to decline extremely. Hence the decomposition of a/2 <101> dislocation at γ/γ0 interface is not considered as the optimal option and the stacking faults can be completely replaced by APBcoupled dislocation pairs, as marked in Fig. 9 (a). Moreover, the other two significant changes related to the microstructure should be emphasized. Firstly, we all know the lattice misfit between the γ0 phase and matrix will become more negative with the increasing temperature in single crystal alloys [19,33,34]. Thus the initial irregular dislocation networks were formed at γ/γ0 interface to help release the increasing coherency strains (marked by quadrangle). Secondly, the strength of γ0 precipitate was believed to decrease when the temperature exceeded approximately 900 � C [3,35]. Very limited hindrance to dislocation shearing behaviors could be expected by these irregular dislocation networks and softened γ0 phase. Consequently, with the onset of dislo cation pairs crossing interfacial dislocation networks and shearing the γ0 precipitates, a strain softening can be found (after the point of UTS) in tensile curves. According to the above experimental results, a critical transformation temperature of tensile deformation mechanism regarding stacking fault shearing and APB shearing mechanisms is inferred to exist between 760 � C and 980 � C, further thorough tests are required in the future works. At 1100 � C, the dislocation shearing the γ0 phase became less prev alent, and only several long superdislocation pairs with APB were observed, which were mainly defined as a/2 <101> pairs. Note that the evolution of dislocation networks was promoted by the much reduced γ/γ0 lattice misfit while the edge of the γ0 precipitates was found to dissolve into the matrix. The rafted structures had been formed and the deformation mechanism was initially controlled by Orowan by-passing. Subsequently, as the flow stress continued to increase, the paired dis locations containing APB were deduced to cut through the interfacial dislocation networks and γ0 phase. The tensile deformation mechanism and microstructure character istic of the fractured specimen at 1120 � C are basically similar to that at 1100 � C, however, three primary modifications need to be pointed out. Firstly, driven by the promoted thermal activation and high stress as well as the further reduction in γ/γ0 lattice misfit, the perfection process of interfacial dislocation networks had been completed within a short tensile period. Secondly, to reach a thermodynamic equilibrium at higher temperature, the rafting and dissolving behaviors of the γ0 pre cipitates were accelerated compared to 1100 � C. What is the most noteworthy are the variations in the amount and configuration of the cutting superdislocation pairs. Unlike the long straight dislocation pairs at 1100 � C, the shearing dislocations had significantly declined in quantity and they were short in length. Obviously, the dislocation pairs have been impeded from truncating the γ0 phase by the denser interfacial
Fig. 8. Configurations of typical stacking faults and dislocations after tensile fracture at 760 � C.
amount increased to a certain degree while the cross-slip motion was seen. The marked dislocations in Fig. 8 implied that the a/2 <101> superdislocation had cross-slipped from {111} planes to {001} planes, thereby hindering the further slipping of the rest dislocations in {111} planes. The above dislocation configuration is known as Kear-Wilsdorf lock [29], which greatly contributes to the work hardening during tensile deformation at 760 � C. An increasing number of dislocation pairs was also observed in the experimental alloy, thus the deformation mechanism was turned to the mix of stacking fault and APB shearing the γ0 phase, which also helped explain the presence of double yield plateaus in tensile curves. The other striking characteristic at 760 � C is the occurrence of PLC phenomenon after yield behavior in the specimen. Essentially, the serrated plastic flow is commonly found in many Ni-based superalloys within a certain temperature range, and it is believed to be influenced by multiple factors like diffusion rates of solute atoms, activation energy, SFE, etc. [30–32]. In the present work, the discontinuous plastic defor mation began to appear at room temperature and became serious at 760 � C, however, it disappeared above 980 � C. Therefore, it is rational to speculate that the above phenomenon is closely correlated with the presence of stacking faults in the microstructure. A dynamic pinning and unpinning processes existed between the solute atoms and the stacking faults with mobile dislocations, which brought about the typical mode of plastic flow in strain-stress curves. Moreover, affected by the higher activation energy at 760 � C, the specimen exhibited more intensive PLC effect compared to room temperature.
Fig. 9. Typical dislocation configurations after tensile fracture at elevated temperatures: (a) 980 � C, (b) 1120 � C. 8
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dislocation networks while the rafted structures have also exerted ob stacles to the dislocation climbing. Interestingly, a kind of curved dislocation pairs (region A) was observed at elevated temperatures, additionally, it was much more common at 1120 � C, as shown in Fig. 9 (b). It could be identified as a <010> superdislocation, which was generated by the reaction of a/2 <011> dislocations with different Burgers vectors in the matrix [36]. The above dislocation configuration was previously seen in some single crystal alloys after creep rupture at elevated temperatures, and it was inferred to move in a combined way of sliding and climbing [23,37]. Besides, researchers found that the minimum creep rates of the alloys could be lowered down by these dislocations with inferior mobility [38]. Nevertheless, no significant improvement related to the alloy mechan ical properties is expected in this study, considering the much higher local stress during tensile deformation. When the flow stress reaches a threshold value, lots of paired dislocations with APB can easily cut into the γ0 phase, hence the softening stage in stain-stress curves has become definitely evident under elevated temperature conditions.
curation, Supervision. J.G. Li: Resources. Y.Z. Zhou: Project adminis tration, Funding acquisition. X.F. Sun: Funding acquisition, Supervision.
5. Conclusions
References
The tensile behaviors of a novel third generation single crystal su peralloy at various temperatures were studied in this work. The tem perature dependence on the fracture surfaces, microstructure and deformation mechanisms were studied in detail via scanning electron microscope and transmission electron microscope. The main conclu sions could be drawn as follows:
[1] C. Sims, N. Stoloff, W. Hagel, Superalloy ІІ, John Wiley & Sons, New York, 1987. [2] Y.B. Hu, T.S. Cao, C.Q. Cheng, L. Zhang, J. Zhao, Oxidation behavior of a singlecrystal Ni-based superalloy over the temperature range of 850 � C-950 � C in air, Appl. Surf. Sci. 484 (2019) 209–218. [3] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, 2006. [4] R.H. Wu, Z.F. Yue, M. Wang, Effect of initial γ/γ0 microstructure on creep of single crystal nickel-based superalloys: a phase-field simulation incorporating dislocation dynamics, J. Alloys Compd. 779 (2019) 326–334. [5] Q.Q. Ding, S.Z. Li, L.Q. Chen, X.D. Han, Z. Zhang, Q. Yu, J.X. Li, Re segregation at interfacial dislocation network in a nickel-based superalloy, Acta Mater. 154 (2018) 137–146. [6] F.H. Latief, K. Kakehi, Effects of Re content and crystallographic orientation on creep behavior of aluminized Ni-base single crystal superalloys, Mater. Des. 49 (2013) 485–492. [7] L. Liu, J. Meng, J.L. Liu, M.K. Zou, H.F. Zhang, X.D. Sun, Y.Z. Zhou, Influences of Re on low-cycle fatigue behaviors of single crystal superalloys at intermediate temperature, J. Mater. Sci. Technol. 35 (2019) 1917–1924. [8] P.J. Fink, J.L. Miller, D.G. Konitzer, Rhenium reduction-alloy design using an economically strategic element, JOM 62 (2010) 55–57. [9] G.M. Han, Z.X. Zhang, J.G. Ji, T. Jin, X.F. Sun, Z.Q. Hu, High cycle fatigue behavior of a nickel-based single crystal superalloy DD98M at 900 � C, Acta Metall. Sin. 48 (2012) 170–175. [10] G.L. Erickson, The Development and Application of CMSX-10, TMS, Warrendale, PA, 1996, pp. 35–44. [11] W.S. Walston, K.S. O0 Hara, E.W. Rose, Ren�e N6: Third Generation Single Crystal Superalloy, TMS, Warrendale, PA, 1996, pp. 27–34. [12] T. Hino, T. Kobayashi, Y. Koizumi, H. Harada, T. Yamagata, Development of a New Single Crystal Superalloy for Industrial Gas Turbines, TMS, Warrendale, PA, 2000, pp. 729–736. [13] P. Caron, High γ0 Solvus New Generation Nickel-Based Superalloys for Single Crystal Turbine Blade Applications, TMS, Warrendale, PA, 2000, pp. 737–746. [14] W.W. Milligan, S.D. Antolovich, The mechanisms and temperature dependence of superlattice stacking fault formation in the single-crystal superalloy PWA 1480, Metal. Trans. A. 22 (1991) 2309–2318. [15] Z.X. Shi, S.Z. Liu, J. Yu, J.R. Li, Tensile behavior of the second generation single crystal superalloy DD6, J. Iron Steel Res. Int. 22 (2015) 738–742. [16] W.P. Yang, J.R. Li, S.Z. Liu, Z.X. Shi, J.Q. Zhao, X.Q. Wang, Transverse tensile properties of a Nickel-based third generation single crystal superalloy, Rare Metal, Mater. Eng. 47 (2018) 2964–2969. [17] X.H. Xiong, D.M. Quan, P.D. Dai, Z.P. Wang, Q.X. Zhang, Z.F. Yue, Tensile behavior of nickel-base single crystal superalloy DD6, Mater. Sci. Eng., A 636 (2015) 608–612. [18] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, The effects of ruthenium additions on tensile deformation mechanisms of single crystal superalloys at different temperatures, Mater. Des. 63 (2014) 286–293. [19] G.L. Wang, J.L. Liu, J.D. Liu, X.G. Wang, Y.Z. Zhou, X.D. Sun, H.F. Zhang, T. Jin, Temperature dependence of tensile behavior and deformation microstructure of a Re-containing Ni-base single crystal superalloy, Mater. Des. 130 (2017) 131–139. [20] X.P. Tan, J.L. Liu, X.P. Song, T. Jin, X.F. Sun, Z.Q. Hu, Measurements of γ/γ0 lattice misfit and γ0 volume fraction for a Ru-containing nickel-based single crystal superalloy, J. Mater. Sci. Technol. 27 (2011) 899–905. [21] N. Tian, S.G. Tian, H.J. Yan, D.L. Shu, S.K. Zhang, G.Q. Zhao, Deformation mechanisms and analysis of a single crystal nickel-based superalloy during tensile at room temperature, Mater. Sci. Eng., A 744 (2019) 154–162. [22] P. Zhang, Y. Yuan, B. Li, G. Yang, X. Song, Investigation of tensile deformation mechanisms at room temperature in a new Ni-based single crystal superalloy, Phil. Mag. Lett. 96 (2016) 238–245.
Acknowledgements The financial supports provided by National Key R&D Program of China under Grant No. 2017YFA0700704, National Science and Tech nology Major Project under Grant No. 2017-VI-0002-0072, Natural Science Foundation of Liaoning Province under Grant No. 20170520038, National Natural Science Foundation of China (NSFC) under Grant Nos. 51601192 and 51671188, State Key Lab of Advanced Metals and Materials Open Fund under Grant No.2018-Z07 and Youth Innovation Promotion Association, Chinese Academy of Sciences for carrying out this work are gratefully acknowledged. The authors are grateful to J.M. Liu, X.L. Wang (Institute of Metal Research) and S.N. Nie (Shenyang University of Technology) for the indispensable assistance in TEM observation.
(1) The flow stress of the experimental alloy demonstrated distinc tive characteristic after yield. A strong work hardening was observed at 760 � C while the softening behavior presented at elevated temperatures. (2) The fracture mode was identified as cleavage and pure shearing at room temperature and 760 � C respectively. The fracture sur face exhibited a mix of dimples and cleavage at 980 � C, and the fracture mechanism turned into ductile features with tempera ture increasing to 1100 � C and 1120 � C. (3) Stacking faults-related microstructure was clearly observed at room temperature and 760 � C, however, it disappeared at 980 � C. The interfacial dislocation networks were formed at elevated temperatures to help release the increasing γ/γ0 lattice strain. In addition, several modifications with regard to the microstructure were illustrated when the temperature was raised from 1100 � C to 1120 � C. (4) As the temperature increased from room temperature to 980 � C, the deformation mechanism changed from stacking faults shearing the γ0 precipitates to APB shearing the strengthening phase, additionally, these two modes were believed to coexist at 760 � C. Above 980 � C, the dislocation climbing and Orowan bypassing became common and the APB shearing mechanism was considered to predominate when the applied stress reached a critical value. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Z.H. Tan: Writing - original draft, Methodology, Investigation. X.G. Wang: Conceptualization, Project administration, Funding acquisition. Y.L. Du: Methodology, Writing - review & editing. T.F. Duan: Funding acquisition, Writing - review & editing. Y.H. Yang: Software, Visuali zation. J.L. Liu: Formal analysis. J.D. Liu: Validation. L. Yang: Data 9
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Materials Science & Engineering A 776 (2020) 138997 [31] P. Fernandez-Zelaia, B.S. Adair, V.M. Barker, S.D. Antolovich, The portevin-le chatelier effect in the Ni-based superalloy IN100, Metall. Mater. Trans. 46 (2015) 5596–5609. [32] C.Y. Cui, C.G. Tian, Y.Z. Zhou, T. Jin, X.F. Sun, Dynamic Strain Aging in Ni Base Alloys with Different Stacking Fault Energy, TMS, Warrendale, PA, 2012, pp. 715–722. [33] S. Ma, V. Seetharaman, B.S. Majumdar, CRSS of γ/γ’ phases from in situ neutron diffraction of a directionally solidified superalloy tension tested at 900 � C, Acta Mater. 56 (2008) 4102–4113. [34] G. Bruno, S. Schumacher, H.C. Pinto, C. Schulze, Measurement of the lattice misfit of the nickel-base superalloy SC16 by high-energy synchrotron radiation, Metall. Mater. Trans. 34 (2003) 193–197. [35] P. Beadmore, R.G. Davies, T.L. Johnston, On the temperature dependence of the flow stress of Nickel-base alloys, Trans. AIME. 245 (1969) 1537–1545. [36] G. Eggeler, A. Dlouhy, On the formation of <010>-dislocations in the γ0 -phase of superalloy single crystals during high temperature low stress creep, Acta Mater. 45 (1997) 4251–4262. [37] P.M. Sarosi, R. Srinivasan, G.F. Eggeler, M.V. Nathal, M.J. Mills, Observations of a <010> dislocations during the high-temperature creep of Ni-based superalloy single crystals deformed along the [001] orientation, Acta Mater. 55 (2007) 2509–2518. [38] R. Srinivasan, G.F. Eggeler, M.J. Mills, γ0 -cutting as rate-controlling recovery process during high-temperature and low-stress creep of superalloy single crystals, Acta Mater. 48 (2000) 4867–4878.
[23] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Z.Q. Hu, J.H. Do, B.G. Choi, I.S. Kim, C.Y. Jo, Dislocation motion during high-temperature low-stress creep in Ru-free and Rucontaining single-crystal superalloys, Mater. Des. 67 (2015) 543–551. [24] J.X. Zhang, T. Murakumo, Y. Koizumi, T. Kobayashi, H. Harada, Slip geometry of dislocations related to cutting of the γ0 phase in a new generation single-crystal superalloy, Acta Mater. 51 (2003) 5073–5081. [25] X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Tensile behaviors and deformation mechanisms of a nickel-base single crystal superalloy at different temperatures, Mater. Sci. Eng., A 598 (2014) 154–161. [26] S.G. Tian, Z. Zheng, F.S. Liang, Z. Chao, L. Chen, Creep behavior of a 4.5%-Re single crystal nickel-based superalloy at intermediate temperatures, Mater. Sci. Eng., A 543 (2012) 104–109. [27] J.L. Liu, J.J. Yu, T. Jin, X.F. Sun, H.R. Guan, Z.Q. Hu, Influence of temperature on tensile behavior and deformation mechanism of Re-containing single crystal superalloy, Trans. Nonferrous Metals Soc. China 21 (2011) 1518–1523. [28] L. Kovarik, R.R. Unocic, J. Li, P. Sarosi, C. Shen, Y. Wang, M.J. Mills, Microtwinning and other shearing mechanisms at intermediate temperatures in Nibased superalloys, Prog. Mater. Sci. 54 (2009) 839–873. [29] K.J. Hemker, M.J. Mills, W.D. Nix, An investigation of the mechanisms that control intermediate temperature creep of Ni3Al, Acta Metall. Mater. 39 (1991) 1901–1913. [30] Y.L. Cai, C.G. Tian, S.H. Fu, G.M. Han, C.Y. Cui, Q.C. Zhang, Influence of γʹ precipitates on portevin–le chatelier effect of Ni-based superalloys, Mater. Sci. Eng., A 638 (2015) 314–321.
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