- Email: [email protected]
Temperature dependence of deformation behavior in a Co–Al–W-base single crystal superalloy
Materials Science & Engineering A 620 (2014) 36–43
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Temperature dependence of deformation behavior in a Co–Al–W-base single crystal superalloy L. Shi, J.J. Yu n, C.Y. Cui, X.F. Sun Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
art ic l e i nf o
a b s t r a c t
Article history: Received 18 June 2014 Received in revised form 17 September 2014 Accepted 18 September 2014 Available online 26 September 2014
Tensile properties of a single-crystal Co–Al–W–Ni–Cr–Ta alloy with low tungsten content have been studied within the temperatures ranging from 20 to 1000 1C at a constant strain rate of 1.0 10 4 s 1. The alloy exhibits comparable yield strength with that of Co–Al–W-base alloys containing more tungsten. From 600 1C to 800 1C, a yield strength anomaly is observed, probably due to the cross-slip of superdislocations from the octahedral plane to the cube plane. TEM analysis demonstrates that stacking faults (SFs) appear both in γ channels and γ0 precipitates in a wide temperature range. These SFs are responsible for the obvious strain hardening observed in stress–strain curves. From room temperature to 900 1C, the deformation is dominated by dislocations shearing γ0 particles. At 1000 1C, the main deformation mechanism is dislocations bypassing γ0 particles. & 2014 Elsevier B.V. All rights reserved.
Keywords: Co–Al–W-base alloy Tensile behavior Dislocation structures Stacking fault
1. Introduction Nickel-base superalloys, possessing exceptional mechanical properties due to the well known strengthening of L12 type γ0 precipitates, are widely used for manufacturing aircraft and power-generation engine turbines. Recently, Co-base superalloys strengthened by γ0 (Co3(Al,W)) with L12 structure have gained substantial interest. A series of experimental [1–7] and computational [8,9] efforts have been done to study the effects of alloying elements on the microstructure and mechanical property of the new Co-base alloys, suggesting that Co3(Al,W) has some similarities with that of Ni3Al and can be practically used as the strengthening phase of Co-base superalloys. However, in the Co–Al–W-base alloys, large amount of W is added to stabilize Co3(Al,W), leading to a high density. The γ0 solvus temperature is relatively lower compared with that of Nibase superalloys, a big restriction on high temperature applications. Efforts have been dedicated to improve the two-phase γ/γ0 microstructural stability at elevated temperature. Ta is effective to 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– 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]. In a Co–7.5Al– 7W–xCr (x¼ 13, 17, 21, at%) alloy system [5], the γ0 solvus temperature is decreasing as the Cr content increasing. It seems n
Corresponding author. Tel.: þ 86 24 2397 1713. E-mail address: [email protected] (J.J. Yu).
http://dx.doi.org/10.1016/j.msea.2014.09.074 0921-5093/& 2014 Elsevier B.V. All rights reserved.
that the addition of Cr makes against the improvement of γ0 solvus temperature. With 8 at% Cr additions, the oxidation resistance of this new type Co-base superalloys is approaching the level of MAR-M 509 at 800 1C [6]. Recent investigations by Shinagawa et al. [7] indicate that substitution of Ni for W can stabilize the γ0 phase and slightly increase the γ0 solvus temperature, which is beneficial to decrease the density. Besides, a combined addition of Cr and Ni to ternary Co–Al–W system can improve the γ0 solvus temperature [5]. It can be seen that the combination of alloying elements greatly affects the γ0 solvus temperature of Co–Al–Wbase alloys. It is interesting to know the effect of a combined addition of Ni, Cr and Ta on the microstructure and property of Co– Al–W alloy system. Thus, in the present study, a Co–Al–W–Ni–Cr– Ta alloy system is tentatively designed, whose tungsten amount is merely half of those reported Co-base alloys [3,4], and attentions are paid to the microstructure and tensile properties of the alloy and characterizing the main deformation microstructures with the aim of comparison to those of Co-base alloys containing high amount of tungsten.
2. Experimental procedure The nominal composition (at%) of the alloy studied is as follows: Al 10, W 5, Ni 17,Cr 6, Ta 2.7, balance by Co (named as 5W). The master alloy was melted in a vacuum induction furnace, and then directionally solidified into [001] single crystal rods by Bridgman technique at a constant withdraw rate of 6 mm/min. The melting point and γ0 solvus temperature of the alloy were determined by
L. Shi et al. / Materials Science & Engineering A 620 (2014) 36–43
Differential Thermal Analysis (DTA) under high purity Ar atmosphere with a heating rate of 10 1C/min. The heat treatments were carried out as follows: 1310 1C/10 h, furnace coolingþ1000 1C/36 h, air coolingþ 750 1C/24 h, air cooling. Heat-treated samples were polished in a solution of 42 ml H3PO4 þ 34 ml H2SO4 þ24 ml H2O at 10 V. The microstructure was analyzed using a Scanning Electron Microscope (SEM). The volume fraction and size distribution of γ0 precipitates were analyzed by image analyzer. Tensile specimens with a nominal 35 mm gage length and a diameter of 5 mm were machined from heat-treated samples. Tensile tests were conducted at a strain rate of 1 10 4 s 1 from room temperature to 1000 1C with the crystal growth direction parallel to the tension loading direction. During the test, the temperature variation was maintained within 72 1C. At least two identical specimens were tested at each temperature. A JMS-6301F field-emission scanning electron microscope (SEM) was used to observe the fractures. Transverse sections of the fractured specimens were cut into discs with 0.5 mm in thickness. These discs were polished to 50 μm, and then subjected to twinjet polishing in a solution of methanol with 5 vol% perchloric acid at 30 1C and 18–20 V. A JEM 2100 Transmission Electron Microscope (TEM) was used for dislocations analysis.
3. Results 3.1. Microstructures The transformation temperatures measured by DTA are given in Table 1. For comparison, the transformation temperatures of Table 1 Liquidus, solidus, γ0 -solvus temperatures and density of the investigated alloy, together with those of other Co-base and Ni-base superalloys.
37
other Co-base superalloys with γ/γ0 microstructure and the Nibase superalloy CMSX-4 are also given in Table 1. It can be seen that the solidus and liquidus of Co-base superalloys are higher than those of Ni-base superalloys. This suggests that there is a possibility for greater temperature capability compared to Ni-base alloys. However, the γ0 solvus temperature of Co-base alloys is still much lower than that of Ni-base superalloy CMSX-4. By comparing Co–9Al–9W (at%) alloy and Co–7.3Al–6.8W (at%) alloy, the γ0 solvus temperature is decreased from 985 1C to 854 1C, indicating a lower W content depresses the stability of Co3(Al,W). With large amount of Ni additions, the γ0 solvus temperature increases slightly. As mentioned earlier, alloying with certain amount of Cr will result in the decrease of γ0 solvus temperature. Possessing a higher amount of W and Ta, the γ0 solvus temperature of Co–7.8Al– 7.8W–4.5Cr–2Ta (at%) alloy is lower than that of Co–9.9Al–4.8W– 1.8Ta (at%) alloy, probably associated with the addition of Cr. Thus, simple substitution of Ni for W or alloying with certain amount of Cr in Co–Al–W–Ta alloy system is not valid to improve the γ0 solvus temperature. The 5W alloy exhibits a relative higher γ0 solvus temperature and lower density compared with those of Co– 8.8Al–9.8W–2Ta (at%) alloy, suggesting that alloying with high Ni 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 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 exhibit cuboidal morphology, aligned along the 〈100〉 direction, which is similar to that of the typical Ni-base superalloys. It is possible to note that the size distribution is close to a Gaussian distribution (Fig. 1b). The average size of γ0 precipitates is about 310 nm. Based on the results of image analyzer, the γ0 volume fraction of the alloy is about 65%.
3.2. Tensile behavior Alloy
Transformation temperature (1C)
5W Co–7.3Al–6.8W (at%) [5] Co–9.2Al–9W (at%) [1] Co–8.8Al–9.8W–2Ta (at%) [4] Co–7.3Al–7.2W–20.2Ni (at%) [5] Co–7.8Al–7.8W–4.5Cr–2Ta (at%) [4] Co–9.9Al–4.8W–1.8Ta (at%) [5] CMSX-4 [10]
Density (g cm 3) 0
Solidus
Liquidus
γ solvus
1395 – 1441 1407 – 1412 – 1326
1426 – 1466 1451 – 1453 – 1370
1100 854 985 1079 881 960 983 1309
9.32 9.18 9.54 49.54 9.29 – 9.09 8.70
Fig. 2 shows true stress–strain curves of the alloy tested at different temperatures. Crystal orientations of four single-crystal bars used in the present study are about 51, 61, 31, 81 away from [001], respectively. It can be seen that the alloy exhibits different tensile behavior over the experimental temperature ranges. That is at room temperature, a strong strain hardening phenomenon is observed. With the increase of temperature, the degree of strain hardening becomes weak. The partial enlargement of A in the tensile curve tested at room temperature is shown in the top-right 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 (2014) 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.
Fig. 3. Temperature dependence of the 0.2% yield strength of the alloy: together with 0.2% yield stresses of the PWA 1480, MarM-509 and two single-crystal γ/γ0 two-phase Co-base alloys.
observed at the initial stage after yielding up to 800 1C. Such similar phenomenon is also observed in several Ni-base singlecrystal superalloys [11], which is due to a difficult dislocation generation step in the γ matrix and then an easy propagation step in an octahedral slip plane. So far, the authors have studied the effect of microstructures, strain rate and temperature on the serrations. The occurrence of serrations observed in the present study will be discussed in another paper. It seems that the curve tested at 900 1C exhibits two elastic stages. It is not an accidental phenomenon, which is also observed in the other two specimens tested in same conditions. The partial enlargement of the kinking region in the 900 1C curve is shown in the lower right corner. Some small serrations can be observed, while serrations are hardly observed beyond the region. It is not clear whether the serrations are associated with the kinking. Further work needs to be done to clarify the uncommon phenomenon. The 0.2% yield strength as a function of temperature is plotted in Fig. 3. For comparison, those of PWA1480 single crystal superalloy [12], Mar-M509 superalloy [13] and two γ0 -strengthened
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 superalloys is even comparable to that of PWA1480 at a similar strain rate of 8.33 10 5 s 1 above 871 1C [12], so the Co-base alloy system has a great potential for high-temperature structural applications above 900 1C. However the intermediatetemperature 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〉 dislocations on octahedral as well as cubic slip from 600 1C to the peak temperature [4]. In the single phase Co3(Al,W) alloys [14], extensive APB-coupled dislocations are observed below the peak temperature. The associated APB energy, γAPB, represents a barrier which must be overcome if cutting of particle occurs, and the precipitate-cutting stress is expected to be in the order γAPB/b, where b is the Burgers vector. Thus, for γ0 -strengthened Co-base superalloys the strength can be enhanced by alloying elements that increase the APB energy. In view of inadequate understanding of the alloying effect on APB energy, availability of a thermodynamic database would provide useful guidance in identifying the key alloying elements. Similar to another two Co-base alloys in Fig. 3, the 5W alloy exhibits a three-stage temperature dependence of the stress. In the low temperature range (below 600 1C), the strength decreases with increasing temperature. In intermediate temperature range an anomalous behavior is found which peaks at around 800 1C. In the high temperature range (above 800 1C), the strength decreases again. At room temperature, the yield strength of 5W alloy is lower than that of Co–8.8Al–9.8W–2Ta (at%) alloy. With temperature increasing, the 5W alloy and Co–8.8Al–9.8W–2Ta (at%) possess similar strength up to 700 1C. Nevertheless, above the peak temperature, the strength of the 5W alloy is inferior to that of Co–8.8Al–9.8W–2Ta (at%), exhibiting a slightly rapid decrease in strength. Compared with Co–9.4Al–10.7W (at%), the 5W alloy exhibits an higher yield strength during the testing temperature
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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.
3.4. Dislocation structures
observed in a single-crystal Ni-base superalloy, which was attributed to the deviation of crystal orientation from the ideal [001] [15]. A similar dislocation activity in parallel or vertical channels is also observed in Co–9Al–9W–0.1B (at%) alloy after creep tested at 850 1C [16]. Their view on such behavior is based on Nabarro's predictions [17] that in the first stage of creep the dislocation activity is contained in one family of matrix channels either parallel (vertical) or perpendicular (horizontal) to the external applied stress and the preferred channel is dependent on the mode of creep deformation either compressive or tensile and the sign of the lattice mismatch between the phases γ and γ0 . The tensile test is a transient deformation process, much similar with the primary stage of creep. Thus, the two conditions mentioned above may result in the different local stress states in both orientations of matrix channels. SFs crossings and interactions with dislocations in the γ channel are expected to be effective impediments to gliding dislocations on different planes. It can be deduced that the strong strain hardening phenomenon observed at room temperature is related to the occurrence of SFs in the γ channel. High density of SFs extending the entire γ0 phase is frequently observed. In contrast to the single, isolated SFs shearing the entire γ0 particles, fault loops are occasionally observed within γ0 , as shown in Fig. 5. The formation of three SF configurations will be discussed in Section 4.2. The deformation microstructures observed at 300 1C (shown in Fig. 6) are similar to those observations at room temperature. SFs in γ0 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 at room temperature. Three SF configurations are observed in the γ0 precipitates or in γ channels, which are SFs extending from the γ/γ0 interface to γ0 precipitates, SF loops within γ0 and SFs in γ matrix. It can be seen that there are more SFs in the γ channels of the [010] direction than those in the [100] direction, which means the [010] and [100] channels are under different stress states. The maldistribution of SFs between [010] channels and [100] channels in tensile specimens tested at room temperature was also
3.4.2. Intermediate temperature regime (600–800 1C) Fig. 7 demonstrates the deformation microstructures at 600 1C, 700 1C and 800 1C, which correspond to temperatures below, near-peak and at the peak stress. At 600 1C (Fig. 7a), SFs are also observed both in γ0 precipitates and γ channels. Some straight dislocations in the γ0 precipitates are observed as well (marked by dark arrows), indicating the shearing mechanism. In the singlecrystal Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al–10.7W(at%) alloys after compression at same temperature, pairs of a/2〈110〉-type
range. This suggests that despite of alloying with a small quantity of tungsten the strength still possesses superior strength with Ni and Ta additions. 3.3. Fracture surfaces Fig. 4 shows the fractures of specimens tested at room temperature, 700 1C and 1000 1C. Below the peak temperature, the samples exhibit similar fracture characteristics. At room temperature and at 700 1C, fractures are wedge-shaped, and cross sections exhibit elliptical shape, indicating a predominance of a single 〈110〉{111} type slip system. Less river patterns and cleavage ledges are observed in the fracture surfaces. The dimple morphology shown in Fig. 4b and d confirms that the specimens fractured in a ductile mode under tension. At 1000 1C, the fracture surface is cavernous. Many pores as well as facets can be observed in Fig. 4f. At higher temperature and lower strain rate, the movement of dislocations as well as pores formed during solidification can lead to the coalescence of micropores. The continuous coalescence of micropores can result in the spread of microcracks and the occurrence of facets. The spread of microcracks will result in the formation of large cracks, which can contribute to the fracture rapidly. It can be inferred that the specimen tested at 1000 1C fractured in such mode.
40
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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 precipitates and γ channels. (b) Fault loops within γ0 precipitates. Beam direction close to the [001] zone axial.
γ0
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 (2014) 36–43
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 cross slip of the superdislocations from the octahedral plane to the cube plane. SFs in γ0 precipitates and fault loops are also observed. What interests us more is the occurrence of fault loops within γ0 precipitates at high 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. Fig. 7c shows general features of the dislocation substructures observed in the 5W alloy after deformation in tension at peak temperature. High density of SFs in the γ0 phase without extension across the γ channel is commonly observed. The density of SFs at 800 1C is higher than that at 700 and 600 1C (Fig. 7a and b). In the single-crystal alloys Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al– 10.7W (at%) [4], the deformation occurs initially by shearing of the γ0 precipitates by a/2〈110〉 superpartial dislocations that enter as pairs of dislocations from the γ phase at 800 1C. Since the a/2 〈l10〉 partials can dissociate at the γ/γ0 interface or in the γ0 precipitates, yielding a superlattice intrinsic stacking fault (SISF), it is difficult to confirm which one takes place during the experiment. The SFs (projected on the (001) plane) extending along mutually perpendicular directions in the same γ0 precipitate are also observed, suggesting that dislocations with different slip 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 can be seen, probably associated with the work hardening exhibited in the true stress–strain curve. Fault loops are still observed within some γ0 precipitates (marked by dark arrows in Fig. 7c).
3.4.3. High temperature regime (900–1000 1C) The dislocation configuration at 900 1C is shown in Fig. 8. A high density of SFs is observed in the γ0 precipitates, similar to the observations in single-crystal alloy Co–8.8Al–9.8W–2Ta (at%) at 890 1C [4], indicating a similar dissociation mechanism. Thus, the SFs are likely to be generated by the reaction of dislocations at the γ/γ0 interfaces to form partial dislocations and gliding of the partial dislocations across the γ0 precipitates. However, the major deformation mechanism in the alloy Co–9.4Al–10.7W (at%) is bypassing of the precipitates by unpaired a/2〈110〉 dislocations above 800 1C [4]. In the alloys Co–9.4Al–10.7W (at%) and Co– 8.8Al–9.8W–2Ta (at%), the variation of deformation mechanisms above the peak temperature is associated with the SISF energy of Co3(Al,W) impacted by Ta additions. The applied stress that enables the a/3〈112〉 partial dislocations to shear into γ0 precipitates is affected by the SISF energy of Co3(Al,W), a higher SISF energy requiring a higher applied stress and resulting in the increment of strength [9]. Ta is believed to increase the SISF 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 that of Co–8.8Al–9.8W–2Ta (at%), probably resulting from the
41
Fig. 8. Bright-field TEM image of the deformation microstructure after fractured at 900 1C (TEM foil perpendicular to external stress axis). Beam direction close to the [001] zone axial.
refractory-element solid-solution-hardening and the variation in the SISF energy of Co3(Al,W). Fig. 9 shows the configuration of the dislocations after tensile test at 1000 1C. In comparison with the γ0 precipitates sheared by dislocations below 900 1C, deformation at higher temperatures is mainly via bypassing of γ0 particles by dislocations, which is similar to observations in alloy Co–9.4Al–10.7W (at%) above 900 1C. Interrupted tensile test results from Milligan et al. indicate that the first step in deformation was bypassing of the γ0 particles, which was followed by shearing of the γ0 particles later during tension [18]. SFs are observed in γ0 precipitates, indicating the shearing mechanism. Thus, this suggests that the deformation mechanism in the present alloy at 1000 1C is very similar. Compared with Ni-base superalloys, γ/γ0 interface dislocation networks are hardly observed, probably due to the large positive lattice mismatch, resulting in the rapid decrease in yield strength.
4. Discussion 4.1. The yield strength anomaly 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, and finally the yield stress decreases rapidly above 800 1C. In the L12 compound Co3(Al,W), the yield strength exhibits a nearplateau from room temperature to 677 1C and an anomalous increase between 677 and 827 1C [14], which is analogous to the observations in the present alloy. It is generally understood that the anomalous temperature dependence of the yield strength in L12 compounds is caused by pinning of cross-slipped screw segments of superdislocations from the octahedral {111} plane to the cube {100} plane, which is driven by elastic anisotropy and/or lower APB energy on {001} planes. It is reported [14] that the elastic anisotropy factor of Co3(Al,W) is larger than that of Ni3Al by 5–10% and the ratio of APB energy on {111} planes and {001} planes is 1.42 at 700 1C which compares well with those observed for many other L12 compounds. In view of the yield strength anomaly as well as observations in Fig. 7d, it can be inferred that the similar strengthening mechanism exists in the present study. 4.2. Temperature dependence of SFs formation With respect to Ni-base superalloys, the deformation mechanism in tensile tests consists of: (i) shearing of γ0 precipitates by
42
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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 range (20–760 1C), (ii) shearing of γ0 precipitates by the 〈112〉{111} slip system at intermediate temperatures (around 760 1C), and (iii) bypassing followed by shearing of the γ0 precipitates at high temperatures (above 800 1C). SFs are commonly observed at intermediate-temperatures deformations. A typical reaction might then be [19]: a=2〈011〉 þ a=2〈101〉-a=6〈112〉 þ a=3〈112〉 þ SF in γ0
ð1Þ
If the applied stress is sufficient, then the a/3〈112〉 dislocation is able to enter the γ0 precipitate leaving a SISF behind it, and the a/6 〈112〉 remains at the γ/γ0 interface, relaxing the coherent stresses. However, in the present alloy 5W, SFs are observed either in γ0 precipitates or in γ channel with various configurations in a wide temperature range after deformation. SFs in γ0 precipitate are commonly observed at 800 and 900 1C, which can be observed in Ni-base superalloys at lower temperature during tension. In a Tadoped γ0 -strengthened Co-base superalloys [4], high density of SFs that originate from dislocation segments on the γ/γ0 interfaces are observed at 890 1C after compression test, which are also formed by the way of Eq. (1). In the present alloy, SFs observed at high temperature are probably generated in a similar way. In a Rucontaining single-crystal Ni-base superalloys [15], SFs appear in both γ0 precipitates and γ channels after tensile testing at room temperature, which is attributed to a lower SF energy. As for conventional Co-base superalloys, owning a lower SF energy, dislocations frequently dissociate into Shockley partial dislocations bounding SFs even without deformation [20]. Therefore, the formation of SFs in the γ channel is probable due to its lower SF energy. It is interesting to note that fault loops are observed both at low and high temperature. These defects are commonly observed in many alloys such as PWA 1480 [21] and (Co,Ni)3Ti [22] after deformation in the low temperature range. A possible formation way is given [21]: a〈101〉-a=3〈112〉 þ a=3〈211〉þ SF in γ0
ð2Þ
In the compression test of Co3(Al,W) at low temperature range (from 196 1C to 300 1C), fault loops were continually observed, and the TEM analysis suggested these were formed by Eq. (2) [14] as well. A series of experiments [23–25] indicates that the SISF can be switched from the APB. Assuming that the total energy of an unit a〈011〉 dislocation dissociation into two partial a/2〈011〉 dislocations coupled by APB and Eq. (2) is E1 and E2, respectively, then whether this transition happens or not depends on the relative magnitude of total energies E1 and E2 of the two dissociation schemes. The criterion for the transition from APB to SISF is given
in Ref. [24] as follows: 3=7 22=203 γ d1 If APB 4 2:49 cos 2 β þ 1:3 db1 sin 2 β ; γ SISF b then E2 oE1 ;
ð3Þ
where d1 is the equilibrium spacing of the component partials in these two schemes, b is the magnitudes of the Burgers vector of the component partials, β is the angle between the burgers vector and the dislocation line. Thus, it can be inferred that large ratio of APB energy and SFE results in the formation of fault loops at room temperature. In single phase Co3(Al,W), the size and density of the faulted dipoles decrease with the increasing deformation temperature up to 300 1C, where the faulted dipoles disappear [14]. This suggests that the ratio of APB energy and SFE decrease with temperature. However, at 700 1C and 800 1C, fault loops are still observed. The 5W alloy is more complex alloying and certain solute atoms may segregate to defects. The experimental [26] and computational [27] results have proved that the SFE can be reduced by the segregation of solute atoms (i.e. Suzuki segregation). The simulation results [27] revealed that Cr segregated at SFs and Nb additions enhanced the Cr segregation significantly in a single phase Co-base alloy, which resulted in a further decrease of SFE. Thus, there is a possibility that the SFE of sites which solute atoms segregate to is decreased and the small ratio of APB energy and SFE leads to the formation of fault loops. So far, it is not clear which element segregates at SFs and which element can enhance or decrease the SFE of γ0 phase in Co–Al–W-base alloys. Thus, further work needs to be done to clarify the elements segregation behavior at SFs and its effect on SFE.
5. Conclusion The microstructures of the single-crystal Co–Al–W-base superalloy have been characterized by TEM after tensile tests at different temperatures. The deformation mechanisms under different test conditions have been analyzed. The following conclusions can be drawn: (1) Alloying is an effective way to improve the mechanical properties of the new generation Co-base superalloys. With the great reduction of tungsten amount, the alloy still possesses higher γ0 solvus temperature and yield strength by additions of Ta and Ni. (2) From room temperature to 900 1C, the deformation mechanism is the shearing of γ particles by dislocations. At 1000 1C, the deformation mechanism is the bypassing of γ0 particles by dislocations followed by shearing of γ0 particles. The
L. Shi et al. / Materials Science & Engineering A 620 (2014) 36–43
anomalous yield behavior above 600 1C can be attributed to the paired 1/2〈110〉 dislocation slip on both octahedral and cube planes in the γ0 precipitates. (3) Fault loops are observed both at low and high temperature. At low temperature, the occurrence of fault loops is associated with the relative high ratio of APB energy and SFE. At high temperature, one possibility of the formation fault loops is that a lower SFE caused by Suzuki segregation results in the relative high ratio of APB energy and SFE.
[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
Acknowledgments
[15] [16]
This work was partly supported by the National Basic Research Program (973 Program) of China under Grant no. 2010CB631206 and the National Natural Science Foundation of China (NSFC) under Grant nos. 51171179, 51271174, 51331005, U1037601 and 11332010.
[17] [18] [19] [20]
Reference [1] J. Sato, T. Omori, K. Oikawa, I. Ohnuma, R. Kainuma, K. Ishida, Science 312 (2006) 90. [2] T.M. Pollock, J. Dibbern, M. Tsunekane, J. Zhu, A. Suzuki, J. Mater. 62 (2010) 58.
[21] [22] [23] [24] [25] [26] [27]
43
A. Bauer, S. Neumeier, F. Pyczak, M. Göken, Scr. Mater. 63 (2010) 1197. Akane Suzuki, Tresa M. Pollock, Acta Mater. 56 (2008) 1288. H.Y. Yan, V.A. Vorontsov, D. Dye, Intermetallics 48 (2014) 44. L. Klein, Y. Shen, M.S. Killian, S. Virtanen, Corros. Sci. 53 (2011) 2713. Kazuya Shinagawa, et al., Mater. Trans. 49 (2008) 1474. Masahiro Ooshima, Katsushi Tanaka, Norihiko L. Okamoto, Kyosuke Kishida, Haruyuki Inui, J. Alloys Compd. 508 (2010) 71. Alessandro Mottura, Anderson Janotti, Tresa M. Pollock, Intermetallics 28 (2012) 138. C. Walter, B. Hallstedt, N. Warnken, Mater. Sci. Eng. A 397 (2005) 385. L.P. Kubin, B. Lisiecki, P. Caron, Philos. Mag. A 71 (1995) 991. Walter W. Milligan, Stephen D. Antolovich, Metall. Trans. A 18 (1987) 85. H.L. Wheaton, Cobalt 29 (1965) 163. Norihiko L. Okamoto, Takashi Oohashi, Hiroki Adachi, et al., Philos. Mag. 91 (2011) 3667. X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Mater. Sci. Eng. A 598 (2014) 154. F. Pyczak, A. Bauer, M. Göken, S. Neumeier, U. Lorenz, M. Oehring, N. Schell, A. Schreyer, A. Stark, F. Symanzik, Mater. Sci. Eng. A 571 (2013) 13. F.R.N. Nabarro, Metall. Mater. Trans. 27A (1996) 513. W. Milligan, S. Antolovich, Metall. Mater. Trans. A 18 (1987) 85. C.M.F. Rae, R.C. Reed, Acta Mater. 55 (2007) 1067. Kenta Yamanaka, Manami Mori, Akihiko Chiba, Mater. Sci. Eng. A 552 (2012) 69. Walter W. Milligan, Stephen D. Antolovich, Metall. Trans. A 22 (1991) 2309. Y. Liu, T. Takasugi, O. Izumi, T. Takahashi, Acta Metall. 36 (1988) 2959. H. Pak, T. Saburi, S. Nenno, Scr. Mater. 10 (1976) 1081. K. Suzuki, M. Ichihara, S. Takeuchi, Acta Metall. 27 (1979) 193. A. Chiba, S. Hanada, Philos. Mag. A 69 (1994) 751. G.W. Han, I.P. Jones, R.E. Smallman, Acta Mater. 51 (2003) 2731. Yuichiro Koizumi, Takeshi Nukaya, et al., Acta Mater. 60 (2012) 2901.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Temperature dependence of deformation behavior in a Co–Al–W-base single crystal superalloy L. Shi, J.J. Yu n, C.Y. Cui, X.F. Sun Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
art ic l e i nf o
a b s t r a c t
Article history: Received 18 June 2014 Received in revised form 17 September 2014 Accepted 18 September 2014 Available online 26 September 2014
Tensile properties of a single-crystal Co–Al–W–Ni–Cr–Ta alloy with low tungsten content have been studied within the temperatures ranging from 20 to 1000 1C at a constant strain rate of 1.0 10 4 s 1. The alloy exhibits comparable yield strength with that of Co–Al–W-base alloys containing more tungsten. From 600 1C to 800 1C, a yield strength anomaly is observed, probably due to the cross-slip of superdislocations from the octahedral plane to the cube plane. TEM analysis demonstrates that stacking faults (SFs) appear both in γ channels and γ0 precipitates in a wide temperature range. These SFs are responsible for the obvious strain hardening observed in stress–strain curves. From room temperature to 900 1C, the deformation is dominated by dislocations shearing γ0 particles. At 1000 1C, the main deformation mechanism is dislocations bypassing γ0 particles. & 2014 Elsevier B.V. All rights reserved.
Keywords: Co–Al–W-base alloy Tensile behavior Dislocation structures Stacking fault
1. Introduction Nickel-base superalloys, possessing exceptional mechanical properties due to the well known strengthening of L12 type γ0 precipitates, are widely used for manufacturing aircraft and power-generation engine turbines. Recently, Co-base superalloys strengthened by γ0 (Co3(Al,W)) with L12 structure have gained substantial interest. A series of experimental [1–7] and computational [8,9] efforts have been done to study the effects of alloying elements on the microstructure and mechanical property of the new Co-base alloys, suggesting that Co3(Al,W) has some similarities with that of Ni3Al and can be practically used as the strengthening phase of Co-base superalloys. However, in the Co–Al–W-base alloys, large amount of W is added to stabilize Co3(Al,W), leading to a high density. The γ0 solvus temperature is relatively lower compared with that of Nibase superalloys, a big restriction on high temperature applications. Efforts have been dedicated to improve the two-phase γ/γ0 microstructural stability at elevated temperature. Ta is effective to 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– 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]. In a Co–7.5Al– 7W–xCr (x¼ 13, 17, 21, at%) alloy system [5], the γ0 solvus temperature is decreasing as the Cr content increasing. It seems n
Corresponding author. Tel.: þ 86 24 2397 1713. E-mail address: [email protected] (J.J. Yu).
http://dx.doi.org/10.1016/j.msea.2014.09.074 0921-5093/& 2014 Elsevier B.V. All rights reserved.
that the addition of Cr makes against the improvement of γ0 solvus temperature. With 8 at% Cr additions, the oxidation resistance of this new type Co-base superalloys is approaching the level of MAR-M 509 at 800 1C [6]. Recent investigations by Shinagawa et al. [7] indicate that substitution of Ni for W can stabilize the γ0 phase and slightly increase the γ0 solvus temperature, which is beneficial to decrease the density. Besides, a combined addition of Cr and Ni to ternary Co–Al–W system can improve the γ0 solvus temperature [5]. It can be seen that the combination of alloying elements greatly affects the γ0 solvus temperature of Co–Al–Wbase alloys. It is interesting to know the effect of a combined addition of Ni, Cr and Ta on the microstructure and property of Co– Al–W alloy system. Thus, in the present study, a Co–Al–W–Ni–Cr– Ta alloy system is tentatively designed, whose tungsten amount is merely half of those reported Co-base alloys [3,4], and attentions are paid to the microstructure and tensile properties of the alloy and characterizing the main deformation microstructures with the aim of comparison to those of Co-base alloys containing high amount of tungsten.
2. Experimental procedure The nominal composition (at%) of the alloy studied is as follows: Al 10, W 5, Ni 17,Cr 6, Ta 2.7, balance by Co (named as 5W). The master alloy was melted in a vacuum induction furnace, and then directionally solidified into [001] single crystal rods by Bridgman technique at a constant withdraw rate of 6 mm/min. The melting point and γ0 solvus temperature of the alloy were determined by
L. Shi et al. / Materials Science & Engineering A 620 (2014) 36–43
Differential Thermal Analysis (DTA) under high purity Ar atmosphere with a heating rate of 10 1C/min. The heat treatments were carried out as follows: 1310 1C/10 h, furnace coolingþ1000 1C/36 h, air coolingþ 750 1C/24 h, air cooling. Heat-treated samples were polished in a solution of 42 ml H3PO4 þ 34 ml H2SO4 þ24 ml H2O at 10 V. The microstructure was analyzed using a Scanning Electron Microscope (SEM). The volume fraction and size distribution of γ0 precipitates were analyzed by image analyzer. Tensile specimens with a nominal 35 mm gage length and a diameter of 5 mm were machined from heat-treated samples. Tensile tests were conducted at a strain rate of 1 10 4 s 1 from room temperature to 1000 1C with the crystal growth direction parallel to the tension loading direction. During the test, the temperature variation was maintained within 72 1C. At least two identical specimens were tested at each temperature. A JMS-6301F field-emission scanning electron microscope (SEM) was used to observe the fractures. Transverse sections of the fractured specimens were cut into discs with 0.5 mm in thickness. These discs were polished to 50 μm, and then subjected to twinjet polishing in a solution of methanol with 5 vol% perchloric acid at 30 1C and 18–20 V. A JEM 2100 Transmission Electron Microscope (TEM) was used for dislocations analysis.
3. Results 3.1. Microstructures The transformation temperatures measured by DTA are given in Table 1. For comparison, the transformation temperatures of Table 1 Liquidus, solidus, γ0 -solvus temperatures and density of the investigated alloy, together with those of other Co-base and Ni-base superalloys.
37
other Co-base superalloys with γ/γ0 microstructure and the Nibase superalloy CMSX-4 are also given in Table 1. It can be seen that the solidus and liquidus of Co-base superalloys are higher than those of Ni-base superalloys. This suggests that there is a possibility for greater temperature capability compared to Ni-base alloys. However, the γ0 solvus temperature of Co-base alloys is still much lower than that of Ni-base superalloy CMSX-4. By comparing Co–9Al–9W (at%) alloy and Co–7.3Al–6.8W (at%) alloy, the γ0 solvus temperature is decreased from 985 1C to 854 1C, indicating a lower W content depresses the stability of Co3(Al,W). With large amount of Ni additions, the γ0 solvus temperature increases slightly. As mentioned earlier, alloying with certain amount of Cr will result in the decrease of γ0 solvus temperature. Possessing a higher amount of W and Ta, the γ0 solvus temperature of Co–7.8Al– 7.8W–4.5Cr–2Ta (at%) alloy is lower than that of Co–9.9Al–4.8W– 1.8Ta (at%) alloy, probably associated with the addition of Cr. Thus, simple substitution of Ni for W or alloying with certain amount of Cr in Co–Al–W–Ta alloy system is not valid to improve the γ0 solvus temperature. The 5W alloy exhibits a relative higher γ0 solvus temperature and lower density compared with those of Co– 8.8Al–9.8W–2Ta (at%) alloy, suggesting that alloying with high Ni 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 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 exhibit cuboidal morphology, aligned along the 〈100〉 direction, which is similar to that of the typical Ni-base superalloys. It is possible to note that the size distribution is close to a Gaussian distribution (Fig. 1b). The average size of γ0 precipitates is about 310 nm. Based on the results of image analyzer, the γ0 volume fraction of the alloy is about 65%.
3.2. Tensile behavior Alloy
Transformation temperature (1C)
5W Co–7.3Al–6.8W (at%) [5] Co–9.2Al–9W (at%) [1] Co–8.8Al–9.8W–2Ta (at%) [4] Co–7.3Al–7.2W–20.2Ni (at%) [5] Co–7.8Al–7.8W–4.5Cr–2Ta (at%) [4] Co–9.9Al–4.8W–1.8Ta (at%) [5] CMSX-4 [10]
Density (g cm 3) 0
Solidus
Liquidus
γ solvus
1395 – 1441 1407 – 1412 – 1326
1426 – 1466 1451 – 1453 – 1370
1100 854 985 1079 881 960 983 1309
9.32 9.18 9.54 49.54 9.29 – 9.09 8.70
Fig. 2 shows true stress–strain curves of the alloy tested at different temperatures. Crystal orientations of four single-crystal bars used in the present study are about 51, 61, 31, 81 away from [001], respectively. It can be seen that the alloy exhibits different tensile behavior over the experimental temperature ranges. That is at room temperature, a strong strain hardening phenomenon is observed. With the increase of temperature, the degree of strain hardening becomes weak. The partial enlargement of A in the tensile curve tested at room temperature is shown in the top-right 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 (2014) 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.
Fig. 3. Temperature dependence of the 0.2% yield strength of the alloy: together with 0.2% yield stresses of the PWA 1480, MarM-509 and two single-crystal γ/γ0 two-phase Co-base alloys.
observed at the initial stage after yielding up to 800 1C. Such similar phenomenon is also observed in several Ni-base singlecrystal superalloys [11], which is due to a difficult dislocation generation step in the γ matrix and then an easy propagation step in an octahedral slip plane. So far, the authors have studied the effect of microstructures, strain rate and temperature on the serrations. The occurrence of serrations observed in the present study will be discussed in another paper. It seems that the curve tested at 900 1C exhibits two elastic stages. It is not an accidental phenomenon, which is also observed in the other two specimens tested in same conditions. The partial enlargement of the kinking region in the 900 1C curve is shown in the lower right corner. Some small serrations can be observed, while serrations are hardly observed beyond the region. It is not clear whether the serrations are associated with the kinking. Further work needs to be done to clarify the uncommon phenomenon. The 0.2% yield strength as a function of temperature is plotted in Fig. 3. For comparison, those of PWA1480 single crystal superalloy [12], Mar-M509 superalloy [13] and two γ0 -strengthened
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 superalloys is even comparable to that of PWA1480 at a similar strain rate of 8.33 10 5 s 1 above 871 1C [12], so the Co-base alloy system has a great potential for high-temperature structural applications above 900 1C. However the intermediatetemperature 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〉 dislocations on octahedral as well as cubic slip from 600 1C to the peak temperature [4]. In the single phase Co3(Al,W) alloys [14], extensive APB-coupled dislocations are observed below the peak temperature. The associated APB energy, γAPB, represents a barrier which must be overcome if cutting of particle occurs, and the precipitate-cutting stress is expected to be in the order γAPB/b, where b is the Burgers vector. Thus, for γ0 -strengthened Co-base superalloys the strength can be enhanced by alloying elements that increase the APB energy. In view of inadequate understanding of the alloying effect on APB energy, availability of a thermodynamic database would provide useful guidance in identifying the key alloying elements. Similar to another two Co-base alloys in Fig. 3, the 5W alloy exhibits a three-stage temperature dependence of the stress. In the low temperature range (below 600 1C), the strength decreases with increasing temperature. In intermediate temperature range an anomalous behavior is found which peaks at around 800 1C. In the high temperature range (above 800 1C), the strength decreases again. At room temperature, the yield strength of 5W alloy is lower than that of Co–8.8Al–9.8W–2Ta (at%) alloy. With temperature increasing, the 5W alloy and Co–8.8Al–9.8W–2Ta (at%) possess similar strength up to 700 1C. Nevertheless, above the peak temperature, the strength of the 5W alloy is inferior to that of Co–8.8Al–9.8W–2Ta (at%), exhibiting a slightly rapid decrease in strength. Compared with Co–9.4Al–10.7W (at%), the 5W alloy exhibits an higher yield strength during the testing temperature
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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.
3.4. Dislocation structures
observed in a single-crystal Ni-base superalloy, which was attributed to the deviation of crystal orientation from the ideal [001] [15]. A similar dislocation activity in parallel or vertical channels is also observed in Co–9Al–9W–0.1B (at%) alloy after creep tested at 850 1C [16]. Their view on such behavior is based on Nabarro's predictions [17] that in the first stage of creep the dislocation activity is contained in one family of matrix channels either parallel (vertical) or perpendicular (horizontal) to the external applied stress and the preferred channel is dependent on the mode of creep deformation either compressive or tensile and the sign of the lattice mismatch between the phases γ and γ0 . The tensile test is a transient deformation process, much similar with the primary stage of creep. Thus, the two conditions mentioned above may result in the different local stress states in both orientations of matrix channels. SFs crossings and interactions with dislocations in the γ channel are expected to be effective impediments to gliding dislocations on different planes. It can be deduced that the strong strain hardening phenomenon observed at room temperature is related to the occurrence of SFs in the γ channel. High density of SFs extending the entire γ0 phase is frequently observed. In contrast to the single, isolated SFs shearing the entire γ0 particles, fault loops are occasionally observed within γ0 , as shown in Fig. 5. The formation of three SF configurations will be discussed in Section 4.2. The deformation microstructures observed at 300 1C (shown in Fig. 6) are similar to those observations at room temperature. SFs in γ0 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 at room temperature. Three SF configurations are observed in the γ0 precipitates or in γ channels, which are SFs extending from the γ/γ0 interface to γ0 precipitates, SF loops within γ0 and SFs in γ matrix. It can be seen that there are more SFs in the γ channels of the [010] direction than those in the [100] direction, which means the [010] and [100] channels are under different stress states. The maldistribution of SFs between [010] channels and [100] channels in tensile specimens tested at room temperature was also
3.4.2. Intermediate temperature regime (600–800 1C) Fig. 7 demonstrates the deformation microstructures at 600 1C, 700 1C and 800 1C, which correspond to temperatures below, near-peak and at the peak stress. At 600 1C (Fig. 7a), SFs are also observed both in γ0 precipitates and γ channels. Some straight dislocations in the γ0 precipitates are observed as well (marked by dark arrows), indicating the shearing mechanism. In the singlecrystal Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al–10.7W(at%) alloys after compression at same temperature, pairs of a/2〈110〉-type
range. This suggests that despite of alloying with a small quantity of tungsten the strength still possesses superior strength with Ni and Ta additions. 3.3. Fracture surfaces Fig. 4 shows the fractures of specimens tested at room temperature, 700 1C and 1000 1C. Below the peak temperature, the samples exhibit similar fracture characteristics. At room temperature and at 700 1C, fractures are wedge-shaped, and cross sections exhibit elliptical shape, indicating a predominance of a single 〈110〉{111} type slip system. Less river patterns and cleavage ledges are observed in the fracture surfaces. The dimple morphology shown in Fig. 4b and d confirms that the specimens fractured in a ductile mode under tension. At 1000 1C, the fracture surface is cavernous. Many pores as well as facets can be observed in Fig. 4f. At higher temperature and lower strain rate, the movement of dislocations as well as pores formed during solidification can lead to the coalescence of micropores. The continuous coalescence of micropores can result in the spread of microcracks and the occurrence of facets. The spread of microcracks will result in the formation of large cracks, which can contribute to the fracture rapidly. It can be inferred that the specimen tested at 1000 1C fractured in such mode.
40
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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 precipitates and γ channels. (b) Fault loops within γ0 precipitates. Beam direction close to the [001] zone axial.
γ0
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 (2014) 36–43
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 cross slip of the superdislocations from the octahedral plane to the cube plane. SFs in γ0 precipitates and fault loops are also observed. What interests us more is the occurrence of fault loops within γ0 precipitates at high 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. Fig. 7c shows general features of the dislocation substructures observed in the 5W alloy after deformation in tension at peak temperature. High density of SFs in the γ0 phase without extension across the γ channel is commonly observed. The density of SFs at 800 1C is higher than that at 700 and 600 1C (Fig. 7a and b). In the single-crystal alloys Co–8.8Al–9.8W–2Ta (at%) and Co–9.4Al– 10.7W (at%) [4], the deformation occurs initially by shearing of the γ0 precipitates by a/2〈110〉 superpartial dislocations that enter as pairs of dislocations from the γ phase at 800 1C. Since the a/2 〈l10〉 partials can dissociate at the γ/γ0 interface or in the γ0 precipitates, yielding a superlattice intrinsic stacking fault (SISF), it is difficult to confirm which one takes place during the experiment. The SFs (projected on the (001) plane) extending along mutually perpendicular directions in the same γ0 precipitate are also observed, suggesting that dislocations with different slip 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 can be seen, probably associated with the work hardening exhibited in the true stress–strain curve. Fault loops are still observed within some γ0 precipitates (marked by dark arrows in Fig. 7c).
3.4.3. High temperature regime (900–1000 1C) The dislocation configuration at 900 1C is shown in Fig. 8. A high density of SFs is observed in the γ0 precipitates, similar to the observations in single-crystal alloy Co–8.8Al–9.8W–2Ta (at%) at 890 1C [4], indicating a similar dissociation mechanism. Thus, the SFs are likely to be generated by the reaction of dislocations at the γ/γ0 interfaces to form partial dislocations and gliding of the partial dislocations across the γ0 precipitates. However, the major deformation mechanism in the alloy Co–9.4Al–10.7W (at%) is bypassing of the precipitates by unpaired a/2〈110〉 dislocations above 800 1C [4]. In the alloys Co–9.4Al–10.7W (at%) and Co– 8.8Al–9.8W–2Ta (at%), the variation of deformation mechanisms above the peak temperature is associated with the SISF energy of Co3(Al,W) impacted by Ta additions. The applied stress that enables the a/3〈112〉 partial dislocations to shear into γ0 precipitates is affected by the SISF energy of Co3(Al,W), a higher SISF energy requiring a higher applied stress and resulting in the increment of strength [9]. Ta is believed to increase the SISF 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 that of Co–8.8Al–9.8W–2Ta (at%), probably resulting from the
41
Fig. 8. Bright-field TEM image of the deformation microstructure after fractured at 900 1C (TEM foil perpendicular to external stress axis). Beam direction close to the [001] zone axial.
refractory-element solid-solution-hardening and the variation in the SISF energy of Co3(Al,W). Fig. 9 shows the configuration of the dislocations after tensile test at 1000 1C. In comparison with the γ0 precipitates sheared by dislocations below 900 1C, deformation at higher temperatures is mainly via bypassing of γ0 particles by dislocations, which is similar to observations in alloy Co–9.4Al–10.7W (at%) above 900 1C. Interrupted tensile test results from Milligan et al. indicate that the first step in deformation was bypassing of the γ0 particles, which was followed by shearing of the γ0 particles later during tension [18]. SFs are observed in γ0 precipitates, indicating the shearing mechanism. Thus, this suggests that the deformation mechanism in the present alloy at 1000 1C is very similar. Compared with Ni-base superalloys, γ/γ0 interface dislocation networks are hardly observed, probably due to the large positive lattice mismatch, resulting in the rapid decrease in yield strength.
4. Discussion 4.1. The yield strength anomaly 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, and finally the yield stress decreases rapidly above 800 1C. In the L12 compound Co3(Al,W), the yield strength exhibits a nearplateau from room temperature to 677 1C and an anomalous increase between 677 and 827 1C [14], which is analogous to the observations in the present alloy. It is generally understood that the anomalous temperature dependence of the yield strength in L12 compounds is caused by pinning of cross-slipped screw segments of superdislocations from the octahedral {111} plane to the cube {100} plane, which is driven by elastic anisotropy and/or lower APB energy on {001} planes. It is reported [14] that the elastic anisotropy factor of Co3(Al,W) is larger than that of Ni3Al by 5–10% and the ratio of APB energy on {111} planes and {001} planes is 1.42 at 700 1C which compares well with those observed for many other L12 compounds. In view of the yield strength anomaly as well as observations in Fig. 7d, it can be inferred that the similar strengthening mechanism exists in the present study. 4.2. Temperature dependence of SFs formation With respect to Ni-base superalloys, the deformation mechanism in tensile tests consists of: (i) shearing of γ0 precipitates by
42
L. Shi et al. / Materials Science & Engineering A 620 (2014) 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 range (20–760 1C), (ii) shearing of γ0 precipitates by the 〈112〉{111} slip system at intermediate temperatures (around 760 1C), and (iii) bypassing followed by shearing of the γ0 precipitates at high temperatures (above 800 1C). SFs are commonly observed at intermediate-temperatures deformations. A typical reaction might then be [19]: a=2〈011〉 þ a=2〈101〉-a=6〈112〉 þ a=3〈112〉 þ SF in γ0
ð1Þ
If the applied stress is sufficient, then the a/3〈112〉 dislocation is able to enter the γ0 precipitate leaving a SISF behind it, and the a/6 〈112〉 remains at the γ/γ0 interface, relaxing the coherent stresses. However, in the present alloy 5W, SFs are observed either in γ0 precipitates or in γ channel with various configurations in a wide temperature range after deformation. SFs in γ0 precipitate are commonly observed at 800 and 900 1C, which can be observed in Ni-base superalloys at lower temperature during tension. In a Tadoped γ0 -strengthened Co-base superalloys [4], high density of SFs that originate from dislocation segments on the γ/γ0 interfaces are observed at 890 1C after compression test, which are also formed by the way of Eq. (1). In the present alloy, SFs observed at high temperature are probably generated in a similar way. In a Rucontaining single-crystal Ni-base superalloys [15], SFs appear in both γ0 precipitates and γ channels after tensile testing at room temperature, which is attributed to a lower SF energy. As for conventional Co-base superalloys, owning a lower SF energy, dislocations frequently dissociate into Shockley partial dislocations bounding SFs even without deformation [20]. Therefore, the formation of SFs in the γ channel is probable due to its lower SF energy. It is interesting to note that fault loops are observed both at low and high temperature. These defects are commonly observed in many alloys such as PWA 1480 [21] and (Co,Ni)3Ti [22] after deformation in the low temperature range. A possible formation way is given [21]: a〈101〉-a=3〈112〉 þ a=3〈211〉þ SF in γ0
ð2Þ
In the compression test of Co3(Al,W) at low temperature range (from 196 1C to 300 1C), fault loops were continually observed, and the TEM analysis suggested these were formed by Eq. (2) [14] as well. A series of experiments [23–25] indicates that the SISF can be switched from the APB. Assuming that the total energy of an unit a〈011〉 dislocation dissociation into two partial a/2〈011〉 dislocations coupled by APB and Eq. (2) is E1 and E2, respectively, then whether this transition happens or not depends on the relative magnitude of total energies E1 and E2 of the two dissociation schemes. The criterion for the transition from APB to SISF is given
in Ref. [24] as follows: 3=7 22=203 γ d1 If APB 4 2:49 cos 2 β þ 1:3 db1 sin 2 β ; γ SISF b then E2 oE1 ;
ð3Þ
where d1 is the equilibrium spacing of the component partials in these two schemes, b is the magnitudes of the Burgers vector of the component partials, β is the angle between the burgers vector and the dislocation line. Thus, it can be inferred that large ratio of APB energy and SFE results in the formation of fault loops at room temperature. In single phase Co3(Al,W), the size and density of the faulted dipoles decrease with the increasing deformation temperature up to 300 1C, where the faulted dipoles disappear [14]. This suggests that the ratio of APB energy and SFE decrease with temperature. However, at 700 1C and 800 1C, fault loops are still observed. The 5W alloy is more complex alloying and certain solute atoms may segregate to defects. The experimental [26] and computational [27] results have proved that the SFE can be reduced by the segregation of solute atoms (i.e. Suzuki segregation). The simulation results [27] revealed that Cr segregated at SFs and Nb additions enhanced the Cr segregation significantly in a single phase Co-base alloy, which resulted in a further decrease of SFE. Thus, there is a possibility that the SFE of sites which solute atoms segregate to is decreased and the small ratio of APB energy and SFE leads to the formation of fault loops. So far, it is not clear which element segregates at SFs and which element can enhance or decrease the SFE of γ0 phase in Co–Al–W-base alloys. Thus, further work needs to be done to clarify the elements segregation behavior at SFs and its effect on SFE.
5. Conclusion The microstructures of the single-crystal Co–Al–W-base superalloy have been characterized by TEM after tensile tests at different temperatures. The deformation mechanisms under different test conditions have been analyzed. The following conclusions can be drawn: (1) Alloying is an effective way to improve the mechanical properties of the new generation Co-base superalloys. With the great reduction of tungsten amount, the alloy still possesses higher γ0 solvus temperature and yield strength by additions of Ta and Ni. (2) From room temperature to 900 1C, the deformation mechanism is the shearing of γ particles by dislocations. At 1000 1C, the deformation mechanism is the bypassing of γ0 particles by dislocations followed by shearing of γ0 particles. The
L. Shi et al. / Materials Science & Engineering A 620 (2014) 36–43
anomalous yield behavior above 600 1C can be attributed to the paired 1/2〈110〉 dislocation slip on both octahedral and cube planes in the γ0 precipitates. (3) Fault loops are observed both at low and high temperature. At low temperature, the occurrence of fault loops is associated with the relative high ratio of APB energy and SFE. At high temperature, one possibility of the formation fault loops is that a lower SFE caused by Suzuki segregation results in the relative high ratio of APB energy and SFE.
[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
Acknowledgments
[15] [16]
This work was partly supported by the National Basic Research Program (973 Program) of China under Grant no. 2010CB631206 and the National Natural Science Foundation of China (NSFC) under Grant nos. 51171179, 51271174, 51331005, U1037601 and 11332010.
[17] [18] [19] [20]
Reference [1] J. Sato, T. Omori, K. Oikawa, I. Ohnuma, R. Kainuma, K. Ishida, Science 312 (2006) 90. [2] T.M. Pollock, J. Dibbern, M. Tsunekane, J. Zhu, A. Suzuki, J. Mater. 62 (2010) 58.
[21] [22] [23] [24] [25] [26] [27]
43
A. Bauer, S. Neumeier, F. Pyczak, M. Göken, Scr. Mater. 63 (2010) 1197. Akane Suzuki, Tresa M. Pollock, Acta Mater. 56 (2008) 1288. H.Y. Yan, V.A. Vorontsov, D. Dye, Intermetallics 48 (2014) 44. L. Klein, Y. Shen, M.S. Killian, S. Virtanen, Corros. Sci. 53 (2011) 2713. Kazuya Shinagawa, et al., Mater. Trans. 49 (2008) 1474. Masahiro Ooshima, Katsushi Tanaka, Norihiko L. Okamoto, Kyosuke Kishida, Haruyuki Inui, J. Alloys Compd. 508 (2010) 71. Alessandro Mottura, Anderson Janotti, Tresa M. Pollock, Intermetallics 28 (2012) 138. C. Walter, B. Hallstedt, N. Warnken, Mater. Sci. Eng. A 397 (2005) 385. L.P. Kubin, B. Lisiecki, P. Caron, Philos. Mag. A 71 (1995) 991. Walter W. Milligan, Stephen D. Antolovich, Metall. Trans. A 18 (1987) 85. H.L. Wheaton, Cobalt 29 (1965) 163. Norihiko L. Okamoto, Takashi Oohashi, Hiroki Adachi, et al., Philos. Mag. 91 (2011) 3667. X.G. Wang, J.L. Liu, T. Jin, X.F. Sun, Mater. Sci. Eng. A 598 (2014) 154. F. Pyczak, A. Bauer, M. Göken, S. Neumeier, U. Lorenz, M. Oehring, N. Schell, A. Schreyer, A. Stark, F. Symanzik, Mater. Sci. Eng. A 571 (2013) 13. F.R.N. Nabarro, Metall. Mater. Trans. 27A (1996) 513. W. Milligan, S. Antolovich, Metall. Mater. Trans. A 18 (1987) 85. C.M.F. Rae, R.C. Reed, Acta Mater. 55 (2007) 1067. Kenta Yamanaka, Manami Mori, Akihiko Chiba, Mater. Sci. Eng. A 552 (2012) 69. Walter W. Milligan, Stephen D. Antolovich, Metall. Trans. A 22 (1991) 2309. Y. Liu, T. Takasugi, O. Izumi, T. Takahashi, Acta Metall. 36 (1988) 2959. H. Pak, T. Saburi, S. Nenno, Scr. Mater. 10 (1976) 1081. K. Suzuki, M. Ichihara, S. Takeuchi, Acta Metall. 27 (1979) 193. A. Chiba, S. Hanada, Philos. Mag. A 69 (1994) 751. G.W. Han, I.P. Jones, R.E. Smallman, Acta Mater. 51 (2003) 2731. Yuichiro Koizumi, Takeshi Nukaya, et al., Acta Mater. 60 (2012) 2901.