& Engineering A 620 (2015) 36–43

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Materials Science & Engineering A

journal homepage: www.elsevier.com/locate/msea

Temperature dependence of deformation behavior in a Co–Al–W-base single crystal

L. Shi, J.J. Yu n, C.Y. Cui, X.F. Sun

Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China article info abstract

Article history: Tensile properties of a single-crystal Co–Al–W–Ni–Cr–Ta alloy with low tungsten content have been Received 18 June 2014 studied within the temperatures ranging from 20 to 1000 1C at a constant strain rate of 1.0 10 4 s1. Received in revised form The alloy exhibits comparable strength with that of Co–Al–W-base alloys containing more 17 September 2014 tungsten. From 600 1Cto8001C, a yield strength anomaly is observed, probably due to the cross- Accepted 18 September 2014 of superdislocations from the octahedral plane to the cube plane. TEM analysis demonstrates that Available online 26 September 2014 stacking faults (SFs) appear both in γ channels and γ0 precipitates in a wide temperature range. These SFs Keywords: are responsible for the obvious strain hardening observed in stress–strain curves. From room – – Co Al W-base alloy temperature to 900 1C, the deformation is dominated by shearing γ0 particles. At 1000 1C, Tensile behavior 0 the main deformation mechanism is dislocations bypassing γ particles. structures & 2014 Elsevier B.V. All rights reserved. Stacking fault

1. Introduction that the addition of Cr makes against the improvement of γ0 solvus temperature. With 8 at% Cr additions, the oxidation resistance of Nickel-base superalloys, possessing exceptional mechanical this new type Co-base superalloys is approaching the level of properties due to the well known strengthening of L12 type MAR-M 509 at 800 1C [6]. Recent investigations by Shinagawa γ0 precipitates, are widely used for manufacturing aircraft and et al. [7] indicate that substitution of Ni for W can stabilize the γ0 power-generation engine turbines. Recently, Co-base superalloys phase and slightly increase the γ0 solvus temperature, which is 0 strengthened by γ (Co3(Al,W)) with L12 structure have gained beneficial to decrease the density. Besides, a combined addition of substantial interest. A series of experimental [1–7] and computa- Cr and Ni to ternary Co–Al–W system can improve the γ0 solvus tional [8,9] efforts have been done to study the effects of alloying temperature [5]. It can be seen that the combination of alloying elements on the microstructure and mechanical property of the elements greatly affects the γ0 solvus temperature of Co–Al–W- new Co-base alloys, suggesting that Co3(Al,W) has some simila- base alloys. It is interesting to know the effect of a combined rities with that of Ni3Al and can be practically used as the addition of Ni, Cr and Ta on the microstructure and property of Co– strengthening phase of Co-base superalloys. Al–W alloy system. Thus, in the present study, a Co–Al–W–Ni–Cr– However, in the Co–Al–W-base alloys, large amount of W is Ta alloy system is tentatively designed, whose tungsten amount is 0 added to stabilize Co3(Al,W), leading to a high density. The γ merely half of those reported Co-base alloys [3,4], and attentions solvus temperature is relatively lower compared with that of Ni- are paid to the microstructure and tensile properties of the alloy base superalloys, a big restriction on high temperature applica- and characterizing the main deformation microstructures with the tions. Efforts have been dedicated to improve the two-phase γ/γ0 aim of comparison to those of Co-base alloys containing high microstructural stability at elevated temperature. Ta is effective to amount of tungsten. improve the γ0 solvus temperature in ternary Co–Al–W system [3,4]. The γ0 solvus temperature of Co–9Al–8W–2Ta–2Cr (at%) alloy is above 1050 1C which is slightly lower than that of Co– 2. Experimental procedure 9Al–8W–2Ta (at%) alloy [3], while the γ0 solvus temperature of Co– 7.8Al–7.8W–2Ta–4.5Cr (at%) alloy is only 960 1C [4].InaCo–7.5Al– The nominal composition (at%) of the alloy studied is as follows: 7W–xCr (x¼13, 17, 21, at%) alloy system [5], the γ0 solvus Al10,W5,Ni17,Cr6,Ta2.7,balancebyCo(namedas5W).The temperature is decreasing as the Cr content increasing. It seems master alloy was melted in a vacuum induction furnace, and then directionally solidified into [001] single crystal rods by Bridgman n Corresponding author. Tel.: þ86 24 2397 1713. technique at a constant withdraw rate of 6 mm/min. The melting 0 E-mail address: [email protected] (J.J. Yu). point and γ solvus temperature of the alloy were determined by http://dx.doi.org/10.1016/j.msea.2014.09.074 0921-5093/& 2014 Elsevier B.V. All rights reserved. L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43 37

Differential Thermal Analysis (DTA) under high purity Ar atmosphere other Co-base superalloys with γ/γ0 microstructure and the Ni- with a heating rate of 10 1C/min. The heat treatments were carried base superalloy CMSX-4 are also given in Table 1. It can be seen out as follows: 1310 1C/10 h, furnace coolingþ1000 1C/36 h, air cool- that the solidus and liquidus of Co-base superalloys are higher ingþ750 1C/24 h, air cooling. Heat-treated samples were polished than those of Ni-base superalloys. This suggests that there is a in a solution of 42 ml H3PO4þ34 ml H2SO4þ24 ml H2Oat10V. possibility for greater temperature capability compared to Ni-base The microstructure was analyzed using a Scanning Electron Micro- alloys. However, the γ0 solvus temperature of Co-base alloys is still scope (SEM). The volume fraction and size distribution of γ0 much lower than that of Ni-base superalloy CMSX-4. By comparing precipitates were analyzed by image analyzer. Co–9Al–9W (at%) alloy and Co–7.3Al–6.8W (at%) alloy, the γ0 Tensile specimens with a nominal 35 mm gage length and a solvus temperature is decreased from 985 1C to 854 1C, indicating diameter of 5 mm were machined from heat-treated samples. a lower W content depresses the stability of Co3(Al,W). With large Tensile tests were conducted at a strain rate of 1 10 4 s1 from amount of Ni additions, the γ0 solvus temperature increases room temperature to 1000 1C with the crystal growth direction slightly. As mentioned earlier, alloying with certain amount of Cr parallel to the tension loading direction. During the test, the will result in the decrease of γ0 solvus temperature. Possessing a temperature variation was maintained within 72 1C. At least higher amount of W and Ta, the γ0 solvus temperature of Co–7.8Al– two identical specimens were tested at each temperature. A 7.8W–4.5Cr–2Ta (at%) alloy is lower than that of Co–9.9Al–4.8W– JMS-6301F field-emission scanning electron microscope (SEM) 1.8Ta (at%) alloy, probably associated with the addition of Cr. Thus, was used to observe the fractures. Transverse sections of the simple substitution of Ni for W or alloying with certain amount of fractured specimens were cut into discs with 0.5 mm in thickness. Cr in Co–Al–W–Ta alloy system is not valid to improve the γ0 These discs were polished to 50 μm, and then subjected to twin- solvus temperature. The 5W alloy exhibits a relative higher γ0 jet polishing in a solution of methanol with 5 vol% perchloric acid solvus temperature and lower density compared with those of Co– at 30 1C and 18–20 V. A JEM 2100 Transmission Electron Micro- 8.8Al–9.8W–2Ta (at%) alloy, suggesting that alloying with high Ni scope (TEM) was used for dislocations analysis. and high Ta can overwhelm negative Cr effect as well as the negative low W effect. The SEM micrographs and frequency size distribution of γ0 3. Results precipitates of the heat-treated sample are shown in Fig. 1. The 5W alloy is only constituted of γ and γ0 phases. The γ0 precipitates 〈 〉 3.1. Microstructures exhibit cuboidal morphology, aligned along the 100 direction, which is similar to that of the typical Ni-base superalloys. It is The transformation temperatures measured by DTA are given possible to note that the size distribution is close to a Gaussian γ0 in Table 1. For comparison, the transformation temperatures of distribution (Fig. 1b). The average size of precipitates is about 310 nm. Based on the results of image analyzer, the γ0 volume

Table 1 fraction of the alloy is about 65%. Liquidus, solidus, γ0-solvus temperatures and density of the investigated alloy, together with those of other Co-base and Ni-base superalloys. 3.2. Tensile behavior Alloy Transformation Density temperature (1C) (g cm3) Fig. 2 shows true stress–strain curves of the alloy tested at Solidus Liquidus γ0 solvus different temperatures. Crystal orientations of four single-crystal bars used in the present study are about 51,61,31,81 away from 5W 1395 1426 1100 9.32 [001], respectively. It can be seen that the alloy exhibits different – – –– Co 7.3Al 6.8W (at%) [5] 854 9.18 tensile behavior over the experimental temperature ranges. That is Co–9.2Al–9W (at%) [1] 1441 1466 985 9.54 Co–8.8Al–9.8W–2Ta (at%) [4] 1407 1451 1079 49.54 at room temperature, a strong strain hardening phenomenon is Co–7.3Al–7.2W–20.2Ni (at%) [5] –– 881 9.29 observed. With the increase of temperature, the degree of strain Co–7.8Al–7.8W–4.5Cr–2Ta (at%) [4] 1412 1453 960 – hardening becomes weak. The partial enlargement of A in the Co–9.9Al–4.8W–1.8Ta (at%) [5] –– 983 9.09 tensile curve tested at room temperature is shown in the top-right CMSX-4 [10] 1326 1370 1309 8.70 corner, in which serrations are observed. Serrations are also

Fig. 1. SEM micrographs of γ0 precipitates after heat treatment (a) and the particle size distribution (b). 38 L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43

Fig. 2. True stain–stress curves of the alloy at different test temperatures and the partial enlargement of A and B in curves tested at room temperature and 900 1C.

Co-base superalloys [4] are also included. It can be seen that the strength of γ0-strengthened Co-base superalloys is higher than that of a conventional Co-base superalloy Mar-M509. At high temperature, the yield strength of γ0-strengthened Co-base super- alloys is even comparable to that of PWA1480 at a similar strain rate of 8.33 10 5 s1 above 871 1C [12], so the Co-base alloy system has a great potential for high-temperature structural applications above 900 1C. However the intermediate- temperature strength of γ0-strengthened Co-base superalloys is still inferior to that of Ni-base superalloys. For Ni-base superalloys, the deformation mechanism below peak temperature is well accepted that γ0 precipitates are cut by a/2〈110〉 dislocations on octahedral slip forming an antiphase boundary (APB). Analogously, in Co-base superalloys, γ0 precipitates are cut by a/2〈110〉 disloca- tions on octahedral as well as cubic slip from 600 1C to the peak

temperature [4]. In the single phase Co3(Al,W) alloys [14], exten- sive APB-coupled dislocations are observed below the peak tem-

Fig. 3. Temperature dependence of the 0.2% yield strength of the alloy: together perature. The associated APB energy, γAPB, represents a barrier with 0.2% yield stresses of the PWA 1480, MarM-509 and two single-crystal γ/γ0 which must be overcome if cutting of particle occurs, and the two-phase Co-base alloys. precipitate-cutting stress is expected to be in the order γAPB/b, where b is the Burgers vector. Thus, for γ0-strengthened Co-base observed at the initial stage after yielding up to 800 1C. Such superalloys the strength can be enhanced by alloying elements similar phenomenon is also observed in several Ni-base single- that increase the APB energy. In view of inadequate understanding crystal superalloys [11], which is due to a difficult dislocation of the alloying effect on APB energy, availability of a thermody- generation step in the γ matrix and then an easy propagation step namic database would provide useful guidance in identifying the in an octahedral slip plane. So far, the authors have studied the key alloying elements. effect of microstructures, strain rate and temperature on the Similar to another two Co-base alloys in Fig. 3, the 5W alloy serrations. The occurrence of serrations observed in the present exhibits a three-stage temperature dependence of the stress. In study will be discussed in another paper. It seems that the curve the low temperature range (below 600 1C), the strength decreases tested at 900 1C exhibits two elastic stages. It is not an accidental with increasing temperature. In intermediate temperature range phenomenon, which is also observed in the other two specimens an anomalous behavior is found which peaks at around 800 1C. In tested in same conditions. The partial enlargement of the kinking the high temperature range (above 800 1C), the strength decreases region in the 900 1C curve is shown in the lower right corner. again. At room temperature, the yield strength of 5W alloy is lower Some small serrations can be observed, while serrations are hardly than that of Co–8.8Al–9.8W–2Ta (at%) alloy. With temperature observed beyond the region. It is not clear whether the serrations increasing, the 5W alloy and Co–8.8Al–9.8W–2Ta (at%) possess are associated with the kinking. Further work needs to be done to similar strength up to 700 1C. Nevertheless, above the peak clarify the uncommon phenomenon. temperature, the strength of the 5W alloy is inferior to that of The 0.2% yield strength as a function of temperature is plotted Co–8.8Al–9.8W–2Ta (at%), exhibiting a slightly rapid decrease in in Fig. 3. For comparison, those of PWA1480 single crystal super- strength. Compared with Co–9.4Al–10.7W (at%), the 5W alloy alloy [12], Mar-M509 superalloy [13] and two γ0-strengthened exhibits an higher yield strength during the testing temperature L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43 39

Fig. 4. SEM images of the fracture surfaces at room temperature (a), 700 1C (c) and 1000 1C (e); SEM images (b), (d) and (f) are the partial enlargement of A, B and C area, respectively. range. This suggests that despite of alloying with a small quantity observed in a single-crystal Ni-base superalloy, which was attrib- of tungsten the strength still possesses superior strength with Ni uted to the deviation of crystal orientation from the ideal [001] [15]. and Ta additions. A similar dislocation activity in parallel or vertical channels is also observed in Co–9Al–9W–0.1B (at%) alloy after tested at 850 1C 3.3. Fracture surfaces [16]. Their view on such behavior is based on Nabarro's predictions [17] that in the first stage of creep the dislocation activity is Fig. 4 shows the fractures of specimens tested at room contained in one family of matrix channels either parallel (vertical) temperature, 700 1C and 1000 1C. Below the peak temperature, or perpendicular (horizontal) to the external applied stress and the the samples exhibit similar fracture characteristics. At room preferred channel is dependent on the mode of creep deformation temperature and at 700 1C, fractures are wedge-shaped, and cross either compressive or tensile and the sign of the lattice mismatch sections exhibit elliptical shape, indicating a predominance of a between the phases γ and γ0. The tensile test is a transient single 〈110〉{111} type slip system. Less river patterns and cleavage deformation process, much similar with the primary stage of creep. ledges are observed in the fracture surfaces. The dimple morphol- Thus, the two conditions mentioned above may result in the ogy shown in Fig. 4b and d confirms that the specimens fractured different local stress states in both orientations of matrix channels. in a ductile mode under tension. At 1000 1C, the fracture surface is SFs crossings and interactions with dislocations in the γ channel are cavernous. Many pores as well as facets can be observed in Fig. 4f. expected to be effective impediments to gliding dislocations on At higher temperature and lower strain rate, the movement of different planes. It can be deduced that the strong strain hardening dislocations as well as pores formed during solidification can lead phenomenon observed at room temperature is related to the to the coalescence of micropores. The continuous coalescence of occurrence of SFs in the γ channel. High density of SFs extending micropores can result in the spread of microcracks and the the entire γ0 phase is frequently observed. In contrast to the single, occurrence of facets. The spread of microcracks will result in the isolated SFs shearing the entire γ0 particles, fault loops are occa- formation of large cracks, which can contribute to the fracture sionally observed within γ0, as shown in Fig. 5. The formation of rapidly. It can be inferred that the specimen tested at 1000 1C three SF configurations will be discussed in Section 4.2.The fractured in such mode. deformation microstructures observed at 300 1C (shown in Fig. 6) are similar to those observations at room temperature. SFs in γ0 3.4. Dislocation structures precipitates and fault loops are also observed.

3.4.1. Low temperature regime (20–600 1C) Fig. 5 shows the deformation microstructures of 5W alloy tested 3.4.2. Intermediate temperature regime (600–800 1C) at room temperature. Three SF configurations are observed in the γ0 Fig. 7 demonstrates the deformation microstructures at 600 1C, precipitates or in γ channels, which are SFs extending from the γ/γ0 700 1C and 800 1C, which correspond to temperatures below, interface to γ0 precipitates, SF loops within γ0 and SFs in γ matrix. It near-peak and at the peak stress. At 600 1C(Fig. 7a), SFs are also can be seen that there are more SFs in the γ channels of the [010] observed both in γ0 precipitates and γ channels. Some straight direction than those in the [100] direction, which means the [010] dislocations in the γ0 precipitates are observed as well (marked by and [100] channels are under different stress states. dark arrows), indicating the shearing mechanism. In the single- The maldistribution of SFs between [010] channels and [100] crystal Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al–10.7W(at%) alloys channels in tensile specimens tested at room temperature was also after compression at same temperature, pairs of a/2〈110〉-type 40 L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43

Fig. 5. Bright-field TEM images of the deformation microstructures after fractured at room temperature (TEM foil perpendicular to external stress axis). (a) SFs in γ0 precipitates and γ channels. (b) Fault loops within γ0 precipitates. Beam direction close to the [001] zone axial.

Fig. 6. Bright-field TEM images of the deformation microstructures after fractured at 300 1C (TEM foil perpendicular to external stress axis). (a) SFs in γ0 precipitates. (b) Fault loops within γ0 precipitates.

Fig. 7. Bright-field TEM image of the deformation microstructure after fractured at: (a) 600 1C, (b) 700 1C, (c) 800 1C and (d) is the partial enlargement of the black box area in (c) (TEM foil perpendicular to external stress axis). Beam direction close to the [001] zone axial. L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43 41 dislocations that sheared into γ0 precipitates were frequently observed [4]. Since the a/2〈110〉 dislocation in the γ0 precipitate can dissociate into a/3〈121〉 dislocation and a/6〈112〉 Shockley dislocation accompanying with SFs, the SFs observed in present alloy can be also formed in such way. From the true stress–strain curve at 600 1C, obviously work hardening can be observed. SFs in the γ channels may partially contribute to the work hardening. At 700 1C(Fig. 7b), a typical feature is dislocations shearing into γ0 precipitates. These dislocations (marked by white arrows) lie in the [010] direction, and TEM analysis demonstrates that these dislocations are the a/3〈121〉-type dislocations lying in the {111} plane in the γ0 precipitate while at two beam conditions SFs are invisible. The dislocations (marked by single white arrow in Fig. 7b and d) slip in different directions within same γ0 precipitate, indicating a possibility which is the of the superdisloca- tions from the octahedral plane to the cube plane. SFs in γ0 precipitates and fault loops are also observed. What interests us Fig. 8. Bright-field TEM image of the deformation microstructure after fractured at 1 more is the occurrence of fault loops within γ0 precipitates at high 900 C (TEM foil perpendicular to external stress axis). Beam direction close to the [001] zone axial. temperature, which are generally observed in low temperature range. We will give the explanation on the abnormal appearance of fault loops in Section 4.2. refractory-element -solution-hardening and the variation in

Fig. 7c shows general features of the dislocation substructures the SISF energy of Co3(Al,W). observed in the 5W alloy after deformation in tension at peak Fig. 9 shows the configuration of the dislocations after tensile temperature. High density of SFs in the γ0 phase without extension test at 1000 1C. In comparison with the γ0 precipitates sheared by across the γ channel is commonly observed. The density of SFs at dislocations below 900 1C, deformation at higher temperatures 800 1C is higher than that at 700 and 600 1C(Fig. 7a and b). In the is mainly via bypassing of γ0 particles by dislocations, which single-crystal alloys Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al– is similar to observations in alloy Co–9.4Al–10.7W (at%) above 10.7W (at%) [4], the deformation occurs initially by shearing of 900 1C. Interrupted tensile test results from Milligan et al. indicate the γ0 precipitates by a/2〈110〉 superpartial dislocations that enter that the first step in deformation was bypassing of the γ0 particles, as pairs of dislocations from the γ phase at 800 1C. Since the a/2 which was followed by shearing of the γ0 particles later during 〈l10〉 partials can dissociate at the γ/γ0 interface or in the γ0 tension [18]. SFs are observed in γ0 precipitates, indicating the precipitates, yielding a superlattice intrinsic stacking fault (SISF), shearing mechanism. Thus, this suggests that the deformation it is difficult to confirm which one takes place during the experi- mechanism in the present alloy at 1000 1C is very similar. ment. The SFs (projected on the (001) plane) extending along Compared with Ni-base superalloys, γ/γ0 interface dislocation mutually perpendicular directions in the same γ0 precipitate are networks are hardly observed, probably due to the large positive also observed, suggesting that dislocations with different slip lattice mismatch, resulting in the rapid decrease in yield strength. systems are activated. In the γ channel lots of entangled dislocations and only few SFs can be seen, while strong interaction of SFs within γ0 precipitates 4. Discussion can be seen, probably associated with the work hardening exhib- ited in the true stress–strain curve. Fault loops are still observed 4.1. The yield strength anomaly within some γ0 precipitates (marked by dark arrows in Fig. 7c). As shown in Fig. 3, from room temperature to 600 1C, the yield strength decreases with increasing temperature, and then the yield stress abnormally increases with temperature up to 800 1C, 3.4.3. High temperature regime (900–1000 1C) and finally the yield stress decreases rapidly above 800 1C. In the The dislocation configuration at 900 1C is shown in Fig. 8. L12 compound Co3(Al,W), the yield strength exhibits a near- A high density of SFs is observed in the γ0 precipitates, similar to plateau from room temperature to 677 1C and an anomalous the observations in single-crystal alloy Co–8.8Al–9.8W–2Ta (at%) increase between 677 and 827 1C [14], which is analogous to the at 890 1C [4], indicating a similar dissociation mechanism. Thus, observations in the present alloy. It is generally understood that the SFs are likely to be generated by the reaction of dislocations at the anomalous temperature dependence of the yield strength in 0 the γ/γ interfaces to form partial dislocations and gliding of the L12 compounds is caused by pinning of cross-slipped screw partial dislocations across the γ0 precipitates. However, the major segments of superdislocations from the octahedral {111} plane to deformation mechanism in the alloy Co–9.4Al–10.7W (at%) is the cube {100} plane, which is driven by elastic anisotropy and/or bypassing of the precipitates by unpaired a/2〈110〉 dislocations lower APB energy on {001} planes. It is reported [14] that the above 800 1C [4]. In the alloys Co–9.4Al–10.7W (at%) and Co– elastic anisotropy factor of Co3(Al,W) is larger than that of Ni3Al by 8.8Al–9.8W–2Ta (at%), the variation of deformation mechanisms 5–10% and the ratio of APB energy on {111} planes and {001} above the peak temperature is associated with the SISF energy of planes is 1.42 at 700 1C which compares well with those observed

Co3(Al,W) impacted by Ta additions. The applied stress that for many other L12 compounds. In view of the yield strength enables the a/3〈112〉 partial dislocations to shear into γ0 precipi- anomaly as well as observations in Fig. 7d, it can be inferred that tates is affected by the SISF energy of Co3(Al,W), a higher SISF the similar strengthening mechanism exists in the present study. energy requiring a higher applied stress and resulting in the increment of strength [9]. Ta is believed to increase the SISF 4.2. Temperature dependence of SFs formation energy of Co3(Al,W). Around 900 1C, the yield strength of 5W alloy is higher than that of Co–9.4Al–10.7W (at%) and lower than With respect to Ni-base superalloys, the deformation mechan- that of Co–8.8Al–9.8W–2Ta (at%), probably resulting from the ism in tensile tests consists of: (i) shearing of γ0 precipitates by 42 L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43

Fig. 9. TEM image of the deformation microstructures after fractured at 1000 1C (TEM foil perpendicular to external stress axis). (a) Dislocations bypassing γ0 precipitates. (b) SFs in γ0 precipitates. pairs of a/2〈110〉dislocations on {111} planes in low temperatures in Ref. [24] as follows: 0 range (20–760 1C), (ii) shearing of γ precipitates by the 〈112〉{111} =  3 7 ÀÁ 22=203 ÀÁ γ d1 slip system at intermediate temperatures (around 760 1C), and APB 4 : 2 β þ : d1 2 β ; If 2 49 cos 1 3 b sin 0 γ b ð3Þ (iii) bypassing followed by shearing of the γ precipitates at high SISF o ; temperatures (above 800 1C). SFs are commonly observed at then E2 E1 intermediate-temperatures deformations. A typical reaction might where d1 is the equilibrium spacing of the component partials in then be [19]: these two schemes, b is the magnitudes of the Burgers vector of 0 the component partials, β is the angle between the burgers vector a=2〈011〉þa=2〈101〉-a=6〈112〉þa=3〈112〉þSF in γ ð1Þ and the dislocation line. Thus, it can be inferred that large ratio of If the applied stress is sufficient, then the a/3〈112〉 dislocation is APB energy and SFE results in the formation of fault loops at room 0 able to enter the γ precipitate leaving a SISF behind it, and the a/6 temperature. In single phase Co3(Al,W), the size and density of the 〈112〉 remains at the γ/γ0 interface, relaxing the coherent stresses. faulted dipoles decrease with the increasing deformation tem- However, in the present alloy 5W, SFs are observed either in γ0 perature up to 300 1C, where the faulted dipoles disappear [14]. precipitates or in γ channel with various configurations in a wide This suggests that the ratio of APB energy and SFE decrease with temperature range after deformation. SFs in γ0 precipitate are temperature. However, at 700 1C and 800 1C, fault loops are still commonly observed at 800 and 900 1C, which can be observed in observed. The 5W alloy is more complex alloying and certain Ni-base superalloys at lower temperature during tension. In a Ta- solute atoms may segregate to defects. The experimental [26] and doped γ0-strengthened Co-base superalloys [4], high density of SFs computational [27] results have proved that the SFE can be that originate from dislocation segments on the γ/γ0 interfaces are reduced by the segregation of solute atoms (i.e. Suzuki segrega- observed at 890 1C after compression test, which are also formed tion). The simulation results [27] revealed that Cr segregated at SFs by the way of Eq. (1). In the present alloy, SFs observed at high and Nb additions enhanced the Cr segregation significantly in a temperature are probably generated in a similar way. In a Ru- single phase Co-base alloy, which resulted in a further decrease of containing single-crystal Ni-base superalloys [15], SFs appear in SFE. Thus, there is a possibility that the SFE of sites which solute both γ0 precipitates and γ channels after tensile testing at room atoms segregate to is decreased and the small ratio of APB energy temperature, which is attributed to a lower SF energy. As for and SFE leads to the formation of fault loops. So far, it is not clear conventional Co-base superalloys, owning a lower SF energy, which element segregates at SFs and which element can enhance 0 dislocations frequently dissociate into Shockley partial dislocations or decrease the SFE of γ phase in Co–Al–W-base alloys. Thus, bounding SFs even without deformation [20]. Therefore, the further work needs to be done to clarify the elements segregation formation of SFs in the γ channel is probable due to its lower SF behavior at SFs and its effect on SFE. energy. It is interesting to note that fault loops are observed both at low 5. Conclusion and high temperature. These defects are commonly observed in many alloys such as PWA 1480 [21] and (Co,Ni) Ti [22] after 3 The microstructures of the single-crystal Co–Al–W-base super- deformation in the low temperature range. A possible formation alloy have been characterized by TEM after tensile tests at way is given [21]: different temperatures. The deformation mechanisms under dif- a〈101〉-a=3〈112〉þa=3〈211〉þSF in γ0 ð2Þ ferent test conditions have been analyzed. The following conclu- sions can be drawn:

In the compression test of Co3(Al,W) at low temperature range (from 196 1Cto3001C), fault loops were continually observed, (1) Alloying is an effective way to improve the mechanical proper- and the TEM analysis suggested these were formed by Eq. (2) [14] ties of the new generation Co-base superalloys. With the great as well. A series of experiments [23–25] indicates that the SISF can reduction of tungsten amount, the alloy still possesses higher be switched from the APB. Assuming that the total energy of an γ0 solvus temperature and yield strength by additions of Ta unit a〈011〉 dislocation dissociation into two partial a/2〈011〉 dis- and Ni. locations coupled by APB and Eq. (2) is E1 and E2, respectively, then (2) From room temperature to 900 1C, the deformation mechan- whether this transition happens or not depends on the relative ism is the shearing of γ particles by dislocations. At 1000 1C, 0 magnitude of total energies E1 and E2 of the two dissociation the deformation mechanism is the bypassing of γ particles schemes. The criterion for the transition from APB to SISF is given by dislocations followed by shearing of γ0 particles. The L. Shi et al. / Materials Science & Engineering A 620 (2015) 36–43 43

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