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Effect of ruthenium on high-temperature creep rupture life of a single crystal nickel-based superalloy
Materials Science and Engineering A 528 (2011) 8381–8388
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Effect of ruthenium on high-temperature creep rupture life of a single crystal nickel-based superalloy X.P. Tan a , J.L. Liu a , T. Jin a,∗ , Z.Q. Hu a , H.U. Hong b , B.G. Choi b , I.S. Kim b , C.Y. Jo b a b
Superalloy Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, People’s Republic of China High Temperature Materials Research Group, Korea Institute of Materials Science, 531 Changwondaero, Changwon, Gyeongnam 641-831, South Korea
a r t i c l e
i n f o
Article history: Received 28 March 2011 Received in revised form 17 June 2011 Accepted 28 July 2011 Available online 9 August 2011 Keywords: Nickel-based superalloy Ruthenium (Ru) Microstructural evolution Creep Misfit
a b s t r a c t The addition of 3 wt% ruthenium (Ru) has been found to improve the creep rupture lives of a single crystal Ni-based superalloy under both conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa. Creep curve analysis indicates that the creep mechanisms are different from each condition. The improvement of creep rupture lives by 3 wt% Ru addition is discussed not only from the view of dislocation movement but also the ␥ phase evolution. The change of ␥/␥ lattice misfit in the initial microstructure is believed to be the key role of Ru on the high-temperature creep deformation. The larger negative lattice misfit caused by an addition of 3 wt% Ru induces smaller and more regular ␥ particles in the initial state, as well as denser dislocation networks at the ␥/␥ interface during creep. These two aspects are crucial to the high-temperature creep rupture life. In addition, a little topologically close-packed (TCP) phases are observed after creep rupture in the two experimental alloys. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Single crystal Ni-based superalloys have become the primary candidate materials for blades and vanes in aero-engine or industrial gas turbines due to their high temperature capability, superior creep resistance and good environment resistance. Hightemperature creep resistance of single crystal Ni-based superalloys has been the most concerned point in this field [1–5]. Several microstructural features have been shown to remarkably influence the creep properties of single crystal Ni-based superalloys. The sign and magnitude of ␥/␥ lattice misfit ı, usually defined as ı = 2(a␥ − a␥ )/(a␥ + a␥ ), where a␥ and a␥ are the lattice constants of ␥ and ␥ phases respectively, can affect the rate of directional coarsening as well as the orientation of the ␥ lamellae with respect to the applied stress. It has been generally accepted that the larger negative misfit is beneficial to the high-temperature creep properties of single crystal Ni-based superalloys [6,7]. In addition, the size and morphology of the initial ␥ particles are also important. However, their influences are still controversial. It is believed that both the thickness and degree of perfection of the lamellar structure are important factors influencing the high-temperature creep properties of single crystal superalloys [8–10]. On the development of single crystal Ni-based superalloys, of note is the introduction of significant additions of Re, and the
∗ Corresponding author. Tel.: +86 24 2397 1757; fax: +86 24 23992092. E-mail address: [email protected] (T. Jin). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.07.075
reduction of Ti and Mo to low concentrations [11,12]. It is well accepted that Re plays a key role in improving the creep properties of these materials. Thus, Re has been regarded as the significant symbol to distinguish the first three generations single crystal superalloys. On the one hand, the addition of Re not only impedes the coarsening of ␥ particles, but also stabilizes the ␥ rafted structures due to its sluggish diffusion [13–16]. Moreover, the addition of Re also makes the lattice misfit more negative which is recognized to accelerate the formation of dense dislocation networks at ␥/␥ interfaces due to its preference into ␥ matrix [13,17–19]. All of these are helpful to improve the creep properties. On the other hand, the addition of Re makes the microstructural homogenization during solution heat treatment more difficult due to its low diffusivity and severe microsegregation [20,21]. The precipitation of topologically close-packed (TCP) phases caused by excessive Re additions is the most concerned. It is known that many detrimental effects can be caused by the precipitation of TCP phases, i.e., the depletion of important strengthening elements from the matrix, the cracks initiation induced by the stress concentration, the interruption of the microstructural continuity, etc., which would degrade the creep properties apparently [22–25]. It has been known that the microstructural instability caused by the TCP precipitation limits the further development of single crystal superalloys to fulfill higher requirements. In recent years, however, another significant alloying element Ru has been added in order to further improve the creep properties based on the third generation single crystal superalloys [26,27]. Thus, Ru has become the symbol element for the new generation
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Fig. 1. Initial microstructures of (a) alloy 0Ru and (b) alloy 3Ru.
Table 1 Nominal compositions of the two experimental alloys (wt%). Alloy
Ni
Al
Co
Cr + W + Mo + Ta
Re
Ru
0Ru 3Ru
Bal. Bal.
6 6
12 12
20 20
5 5
0 3
single crystal superalloys. It is generally accepted that the key roles displayed by Ru additions are the promotion of the microstructural stability and improvement of the creep resistance. So far it should say that the knowledge concerning the effects of Ru is far from enough, especially for the mechanisms of effects, and the acquired results still remain debatable [28–34]. This paper aims at elucidating the effect of Ru on the high-temperature creep rupture lives and giving some helpful guidelines for designing new generations single crystal superalloys. 2. Experimental Two single crystal Ni-based superalloys with and without Ru addition were designed to study the effect of Ru on the high-temperature creep rupture lives. The nominal chemical compositions of the alloys are listed in Table 1. According to their different Ru content, these two alloys are named as 0Ru and 3Ru respectively. The master alloys were melted by vacuum induction technique, and then directionally solidified into cylindrical bars (16 mm in diameter and 220 mm long) in an investment casting cluster mold in a Bridgman furnace with a withdrawal rate of 6 mm/min. Conventional helical starters were utilized to initiate single crystal growth. The heat treatment schemes for both alloys are given in Table 2. Specimens for creep tests with a diameter of 5 mm and a length of 25 mm parallel to the solidification direction [0 0 1] were machined from the heat treated alloys by means of electrosparking. The creep tests were carried out at 1000 ◦ C/310 MPa and 1100 ◦ C/150 MPa. Additionally, in an attempt to obtain the equilibrium dislocation networks for the estimation of unconstrained misfit at high temperature, long-term isothermal aging experiments were carried out. Both alloys were exposed at 1100 ◦ C for 1000 h to fully relieve the misfit stress and form equilibrium dislocation networks. The etchant used for microstructural observation is HCl (80 ml) + CuSO4 (20 g) + H2 O (100 ml) solution. A JMS-6301F fieldemission scanning electron microscope (SEM) was used to examine
the microstructure after full heat treatment and creep rupture tests. The fully heat treated and long-term aging specimens were cut into 3 mm discs parallel to (0 0 1) plane with 500 m in thickness and thinned down to 50 m mechanically. After creep tests, the specimens were cut into 3 mm discs of 500 m in thickness normal to the [0 0 1] direction about 3 mm away from the fracture surface and thinned down to 50 m mechanically. They were then electrochemically polished by the twin-jet method, in an solution of 8% perchloric acid and 92% ethanol at −10 ◦ C and 20–30 mA. A TECNAI 20 transmission electron microscope (TEM) equipped with energy dispersive spectrometer (EDS) was used to examine the dislocation configuration after long-term aging and creep rupture. Furthermore, the partitioning behavior of various alloying elements between ␥ matrix and ␥ precipitate in the initial microstructure was characterized by measuring the local chemical composition of ␥ and ␥ in fully heat treated specimens using an EDS in TEM. An average value for three times measurements was adopted. The partitioning ratio k or k defined as ki = Ci␥ /C␥ or ki = Ci /Ci was widely used. Therefore, it can be known that ki > 1 indicates the preferential partition to ␥ matrix for the alloying element, conversely, ki > 1 indicates the preferential partition to ␥ precipitate. 3. Results 3.1. Initial microstructure Fig. 1 shows the microstructures of both alloys after full heat treatment. The differences between alloy 0Ru and 3Ru on the morphology of the ␥ particles can be observed clearly. According to the image analysis results (Fig. 2), the ␥ size and its volume fraction
Table 2 Heat treatment schemes of the two experimental alloys. Alloy
Full heat treatment scheme
0Ru 3Ru
1325 ◦ C/16 h + 1335 ◦ C/16 h, ACa + 1150 ◦ C/4 h, AC + 870 ◦ C/24 h 1315 ◦ C/16 h + 1325 ◦ C/16 h, AC + 1150 ◦ C/4 h, AC + 870 ◦ C/24 h
a
The air cooling.
Fig. 2. Comparison of ␥ size and its volume fraction in the two experimental alloys.
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Fig. 3. Variation of the ␥/␥ partitioning ratio of each alloying element in the two experimental alloys.
were decreased, the ␥ shape more regular and the distribution of ␥ particles more uniform with 3 wt% Ru addition. Fig. 3 illustrates that all the alloying elements show the so-called “reverse partitioning” behavior with 3 wt% Ru addition. It means that Ru promotes the TCP-forming elements, such as Re, W, Mo, Cr etc., to partition to ␥ phase. It is believed that this effect may be the primary reason for the suppression of TCP phases by Ru additions [35,36]. As measured by the X-ray diffraction method, the average ␥/␥ lattice misfits of alloy 0Ru and 3Ru are −0.15% and −0.25% at room temperature, respectively. The detailed measurement can be seen in [37]. There is no doubt that the addition of 3 wt% Ru made the lattice misfit more negative. 3.2. Creep rupture life Fig. 4 shows the creep curves of both alloys under the two conditions. It shows typical high-temperature creep curves with three obvious stages under the condition of 1100 ◦ C/150 MPa; however, it differs at 1000 ◦ C/310 MPa, which shows shorter primary and secondary stages compared with its tertiary stage. It is clear that alloy 3Ru has longer creep rupture lives compared with alloy 0Ru under both conditions. 3.3. Microstructure after creep rupture Due to the short creep rupture lives less than 200 h for all the cases, only a little small TCP phases were observed after creep rupture. As shown in Fig. 5, there exist a certain amount of rectangular
Fig. 4. Creep rupture curves of the two experimental alloys at (a) 1100 ◦ C/150 MPa and (b) 1000 ◦ C/310 MPa.
pores in the cross-section microstructure near the fracture surfaces. Fig. 6 shows these pores with the shape of square on transversesection and rectangular on cross-section in detail. Therefore, they have a shape of rectangular parallelepiped. In addition, it was found that there must be some small TCP phases around every pore. TEM results for the phase identification indicated that these small TCP particles are phase. Thus, it is believed that the formation of these pores should be associated with TCP phases. However, it shows that the creep rupture lives for both the alloys were less dependent on these rectangular pores. It may be due to the fact that most of the
Fig. 5. Longitudinal microstructures near the fracture surface at1100 ◦ C/150 MPa of (a) alloy 0Ru and (b) alloy 3Ru.
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Fig. 6. Backscattered electron micrographs showing microstructures of alloy 0Ru after creep rupture at 1100 ◦ C/150 MPa of (a) transverse-section and (b) cross-section.
fatal cracks were not directly induced by these small rectangular parallelepiped pores. The rafted microstructures of both alloys after creep rupture under the conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa are shown in Fig. 7. It shows an evident “topological inversion” phenomenon [38], namely ␥ phase acts as the matrix and ␥ phase was surrounded by them. All the rafted ␥ lamellae were normal to the applied stress. It indicates that the lattice misfits of both alloys are negative at high temperature. The dislocation networks of both alloys after creep rupture under the conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa are shown in Fig. 8. It can be clearly seen that the dislocation networks were denser in alloy 3Ru than in 0Ru under both conditions. Furthermore, Ru promotes the formation of square dislocation networks. 4. Discussion 4.1. / portioning ratio and / lattice misfit When Ru caused a small amount of ␥-forming elements to partition to ␥ phase, they might replace some ␥ -forming elements in ␥
phase. Some ␥ -forming elements, in the meantime, would enter ␥ phase. Thus, it is reasonable for almost all the elements to show the behavior of reverse partitioning with Ru additions. As known, the ␥/␥ lattice misfit depends on not only the partitioning ratio but also the Vegard law [12]. Re does not have a larger Vegard coefficient [12]; however, it has a larger partitioning ratio as shown in Fig. 3. Therefore, Re acted as an effective element to make the ␥/␥ lattice misfit larger negative. It should be noted that Ru slightly preferentially partitioned to ␥ phase, which means that the addition of Ru must make the lattice misfit more negative. On the other hand, Ta has a much larger Vegard coefficient in ␥ phase than ␥ phase [12]. In the present work, Ru made slightly more Ta enter ␥ phase, which is also believed to make the lattice misfit more negative. Based on the two effects mentioned above, the lattice misfit became more negative with 3 wt% Ru addition. Hence, the influence of Ru on the lattice misfit is achieved by changing the partitioning behavior of other alloying elements. As reported in the previous study [37], there will exist obvious lattice distortion at the ␥/␥ interfaces in the initial microstructure of alloy 3Ru. Therefore, the lattice misfit mentioned here is the constrained misfit which would be highly dependent on the partitioning behavior of alloying elements.
Fig. 7. Secondary electron micrographs showing ␥/␥ microstructures near the fracture surface after creep rupture in (a) alloy 0Ru and (b) alloy 3Ru at 1100 ◦ C/150 MPa; and (c) alloy 0Ru and (d) alloy 3Ru at 1000 ◦ C/310 MPa.
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Fig. 8. TEM images of the interfacial dislocation networks after creep rupture in (a) alloy 0Ru and (b) alloy 3Ru at 1100 ◦ C/150 MPa; and (c) alloy 0Ru and (d) alloy 3Ru at 1000 ◦ C/310 MPa.
4.2. size, volume fraction and / lattice misfit The morphology of ␥ particles depends on both the interfacial energy and elastic strain energy, which are related with the interfacial area and volume of precipitate, respectively [39]. For the alloy with smaller magnitude of lattice misfit, interfacial energy predominates and spherical ␥ particles are shown due to their minimum surface area with the same volume. Inversely, for the alloy with larger magnitude of lattice misfit, elastic strain energy induced by lattice misfit predominates and cuboidal ␥ particles are exhibited due to the crystalline anisotropy and lowest elastic modulus with [0 0 1] orientation. The cuboidal degree of ␥ particles increases with increasing magnitude of lattice misfit. Furthermore, in terms of the characteristics of cuboidal ␥ precipitation, the degree of parallelism of ␥/␥ interfaces is improved with the increase of magnitude
of lattice misfit. Consequently, the distribution of ␥ particles is more uniform in alloy 3Ru. As we know, ␥ particles coherently precipitate from ( matrix in single crystal Ni-based superalloys. Therefore, the total elastic strain energy depends on the morphologies and elastic properties of ␥ and ␥ phases. It is assumed that ( matrix is elastically isotropic and the elastic modulus of ( matrix is equal to that of ␥ phase, and therefore the total elastic strain energy Gs is not related with the morphology of ␥ particles. Thus [39] Gs ∝ ı2 V
(1)
where is shear modulus of matrix, ı is lattice misfit and V is the volume of ␥ particle. It can be known that the elastic strain energy induced by lattice misfit is proportional to the square of lattice misfit and the volume of ␥ particle. Consequently, when the lattice
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Fig. 9. TEM images of the interfacial dislocation networks after 1000 h isothermal aging at 1100 ◦ C in (a) alloy 0Ru and (b) alloy 3Ru.
misfit of the alloy is the more negative, the volume of ␥ particle will be smaller, namely, the ␥ size is smaller. As mentioned previously, alloy 3Ru has a larger negative lattice misfit than alloy 0Ru. Hence the microstructure with smaller cuboidal and more uniform ␥ particles was shown in alloy 3Ru compared with alloy 0Ru. Moreover, it shows the tendency of lower ␥ volume fraction with Ru addition. It could be expected that the variation of initial microstructure caused by Ru addition should be the dominating factor for various hightemperature creep properties. That will be discussed in Section 4.4 in detail. 4.3. / lattice misfit and dislocation networks It is well understood that the magnitude of the lattice misfit controls the density of the interfacial dislocations required to relieve the misfit stresses. Assuming that the network dislocations are orientated to most efficiently relieve the misfit, the lattice misfit is inversely related to the spacing of the interfacial networks (D) such that [40,41] D=
៝ |b| |ı|
(2)
៝ is the magnitude of Burgers vector of the edge compowhere |b| nent and ı is the lattice misfit. Therefore, the average spacing of dislocations within the equilibrium interfacial networks provides a reasonably good estimate of the high-temperature lattice misfit at the temperature which they are formed. The distribution of equilibrium dislocation networks that formed at 1100 ◦ C in both alloys ៝ 0.254 nm and is shown in Fig. 9. Assuming a Burges vector of, |b|, using the dislocation spacing D, the magnitude of the lattice misfit can be estimated at 1100 ◦ C. The average dislocation spacings of the equilibrium interfacial dislocation networks after 1100 ◦ C isothermal aging and the interfacial dislocation networks after creep rupture under the condition of 1100 ◦ C/150 MPa are plotted with the corresponding estimated magnitude of lattice misfit in Fig. 10. As mentioned previously, the interfacial dislocation networks of alloy 3Ru were denser after creep rupture compared with alloy 0Ru. It is firmly believed that the denser interfacial dislocation networks would be an effective barrier to the shearing of ␥ lamellae [19]. It explains that the alloy with 3 wt% Ru addition has longer creep rupture lives from the view of dislocation movement. Additionally, it is easily known that the differences of longitudinal coordinate between the two conditions in Fig. 10 indicate the magnitude of so-called excess dislocation density during creep deformation. The multiplication of mobile dislocations occurs during the tertiary stage of creep deformation at high temperature, which means that the number of dislocations after creep rupture must exceed that of the misfit dislocations necessary to relieve the misfit. Their
differences are the so-called excess dislocations. In the present work, the most creep resistant alloy, 3Ru, had the lower density of excess dislocations. It is opposite, however, with the results obtained by Carroll et al. [41]. That is due to the fact that Ru additions promoted the formation of square dislocation networks. It has been reported that the square configurations of four dislocations of Burges vector of a/2 1 1 0 type are of edge character and thus efficient at relieving the misfit, but these segments have no shear stress resolved on them – thus they must have been formed by reactions driven by reductions in misfit and dislocation line energies [12]. As the remarkable increase of the magnitude of lattice misfit with 3 wt% Ru addition, it gave rise to a great driving force for the formation of the square dislocation networks. Therefore, it shows a lower excess dislocation density in alloy 3Ru. Furthermore, the regular square dislocation networks are more stable and act efficiently as a hindrance to interfacial advance in ␥ coarsening [42], which also can explain why the 3 wt% Ru addition improve the high-temperature creep rupture lives form the view of dislocation movement. 4.4. High-temperature creep rupture lives In terms of the creep curves illustrated in Fig. 4, the time spans of secondary and tertiary stages of each curve, as well as their percentage in the whole creep rupture lives, are listed in Table 3. It can be clearly seen that the secondary stage accounts for the greatest proportion during the whole creep rupture life under the condition
Fig. 10. Plot of the magnitude of the lattice misfit in the two experimental alloys determined from high-temperature equilibrium dislocation network spacings, including the postcrept dislocation network spacings at the corresponding lattice misfit.
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Table 3 The time spans of secondary and tertiary stages and their percentage in the whole creep rupture life of each creep curve under both conditions.
◦
1100 C/150MPa 1000 ◦ C/310MPa a
Alloy
tf a (h)
ts a (h)
tt a (h)
ts /tf
tt /tf
0Ru 3Ru 0Ru 3Ru
156.9 211.2 80.1 127.1
77.9 105.4 18.4 40.0
55.7 75.5 54.5 74.2
49.6% 49.9% 23.0% 31.5%
35.5% 35.7% 68.0% 58.4%
tf , ts and tt denote the creep rupture life, the time span of secondary and tertiary stage respectively.
Fig. 11. Schematic illustration of the ␥ evolution during the creep at high temperatures/low stresses. ␥ and ␥ phases are visualized as bright and dark regions, respectively.
of 1100 ◦ C/150 MPa, and the tertiary stage does at 1000 ◦ C/310 MPa. It indicates that the creep mechanism might be different between these two cases. Therefore, the creep properties under these two conditions should be discussed respectively. The mechanism of creep deformation in 0 0 1-orientated single crystal superalloys is very sensitive to the conditions of temperature and stress. Usually three typical creep regimes can be identified, i.e., the primary creep regime occurred at low temperatures and high stresses, the tertiary creep regime at intermediate temperatures and intermediate stresses and the rafting regime at high temperatures and low stresses, and in each a distinct mode of microstructural degradation is predominant [12]. In the rafting regime, the creep curve exhibits typically three different stages as shown in Fig. 11, and the ␥/␥ microstructure degrades rather quickly because thermally activated processes are favored strongly since the temperature is high. Therefore, it can be termed as the diffusion-controlled creep. During the short primary stage, rafting of ␥ particles occurs and dislocation networks required to relieve the misfit stress develop at the ␥/␥ interfaces. It is followed by a long plateau where a transition from rafting to topological inversion occurs, and the secondary stage is ended by the occurrence of
topological inversion. Afterwards, cutting of ␥ lamellae by paired dislocations begins with the onset of tertiary stage characterized by an acceleration of the deformation during a short time and driving to failure [4,43]. It is clear that the secondary stage predominates over the whole creep rupture life. In the present work of creep at 1100 ◦ C/150 MPa, it is noted that the creep rupture life mainly depends on the time of the occurrence of topological inversion. As mentioned previously, the addition of 3 wt% Ru decreases the ␥ volume fraction, which will prolong the secondary stage and eventually increases the creep rupture live. On the other hand, it has been reported [9] that more perfect ␥ lamellae are more resistant to thickening and thus would delay the onset of tertiary creep, thus prolonging the creep life of single crystal superalloys. Here the addition of 3 wt% Ru makes the ␥ particles more regular and further more perfect ␥ lamellae. Therefore, it can also increase the creep rupture life. Moreover, due to the fact that this process is diffusioncontrolled and the diffusion coefficient of Ru is second to Re [44,45], which further retards the occurrence of topological inversion, it is concluded that 3 wt% Ru addition could significantly promote the creep rupture life under this condition. In the tertiary regime, the creep curve exhibits an obvious tertiary stage as shown in Fig. 12(a), the density of dislocations is found to increase with increasing creep strain, while other forms of microstructural degradation (␥ rafting, creep cavitations) are absent. Usually slight ␥ coarsening occurs [46,47]. Therefore, it can be termed as a dislocation-controlled creep. In the present work, however, creep at 1000 ◦ C/310 MPa which corresponds to the condition of higher temperatures and high stresses, may differ from both conditions of high temperatures/low stresses and intermediate temperatures/intermediate stresses. As shown in Fig. 12(b), the creep curve exhibits three stages with a longer tertiary stage. From the current experimental results, topological inversion of ␥/␥ microstructure was observed. It means that there exists a transition from rafting to topological inversion. In addition to the ␥ evolution, it could be imagined that the cutting of ␥ lamellae by dislocations started from the onset of tertiary stage. Therefore, it can be termed as the dislocation-controlled creep coupled with the diffusion-controlled mode. Apparently the tertiary stage still predominates over the whole creep rupture life (as listed in Table 3). As a result that the creep rupture life will mainly depend on the duration of ␥ lamellae cutting process at 1000 ◦ C/310 MPa. The cutting resistance of ␥ lamellae is dependent on the thickness and
Fig. 12. Schematic illustration of the ␥ evolution during the creep at (a) intermediate temperatures/intermediate stresses and (b) higher temperatures/higher stresses. ␥ and ␥ phases are visualized as bright and dark regions, respectively.
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regularity of ␥ lamellae. As mentioned previously, more regular cuboidal ␥ particles formed in the initial microstructure by 3 wt% Ru addition. It is followed by more perfect ␥ lamellae after a short primary stage. It means that the number of imperfection sites at the ␥/␥ interfaces, where are preferentially cut by dislocations, will be accordingly reduced. Therefore, the more perfect ␥ lamellae formed in the alloy 3Ru are more resistant to dislocation cutting. On the other hand, the denser interfacial dislocation networks formed after the primary stage during creep in the alloy 3Ru could effectively improve the cutting resistance of ␥ lamellae. Thus, the creep rupture life was remarkably increased with 3 wt% Ru addition under this condition. To summarize, the initial microstructure after full heat treatment was noticeably influenced by the 3 wt% Ru addition, which appears to be the major reason for the various creep rupture lives eventually. In the present work, the change of lattice misfit was considered as the key role that Ru has played. From the view of ␥ evolution, lattice misfit influenced the morphology, size, volume fraction and distribution of ␥ particles, and the following regularity and perfection of ␥ lamellae during high-temperature creep. On the other hand, from the view of dislocation movement, it influenced the configuration and density of interfacial dislocation networks and the following dislocation movement. As the two aspects are equally important, it is necessary to correlate them for understanding the high-temperature creep properties. However, it is worth noting that the suppression of TCP phases by Ru addition was not highly stressed in this work. In spite of the occurrence of some small TCP phases, their influences on the creep rupture life could be negligible. 5. Conclusions The creep rupture lives of two single Ni-based superalloys with and without Ru addition were investigated in this paper and the conclusions are summarized as follows: 1. The addition of 3 wt% Ru improves the creep rupture lives of a single crystal Ni-based superalloy under both conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa. 2. At 1100 ◦ C/150 MPa, the secondary stage predominates over the whole creep rupture life. 3 wt% Ru addition retards the occurrence of the topological inversion, which prolongs the duration of the secondary stage. 3. At 1000 ◦ C/310 MPa, the tertiary stage predominates over the whole creep rupture life. 3 wt% Ru addition improves the cutting resistance of ␥ lamellae, which prolongs the duration of the tertiary stage.
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Acknowledgements This work was financially supported by the National Basic Research Program (973 Program) of China under grant No. 2010CB631200 (2010CB631206), the National Natural Science Foundation of China (NSFC) under grant No. 50931004. The authors are grateful for those supports.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of ruthenium on high-temperature creep rupture life of a single crystal nickel-based superalloy X.P. Tan a , J.L. Liu a , T. Jin a,∗ , Z.Q. Hu a , H.U. Hong b , B.G. Choi b , I.S. Kim b , C.Y. Jo b a b
Superalloy Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, People’s Republic of China High Temperature Materials Research Group, Korea Institute of Materials Science, 531 Changwondaero, Changwon, Gyeongnam 641-831, South Korea
a r t i c l e
i n f o
Article history: Received 28 March 2011 Received in revised form 17 June 2011 Accepted 28 July 2011 Available online 9 August 2011 Keywords: Nickel-based superalloy Ruthenium (Ru) Microstructural evolution Creep Misfit
a b s t r a c t The addition of 3 wt% ruthenium (Ru) has been found to improve the creep rupture lives of a single crystal Ni-based superalloy under both conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa. Creep curve analysis indicates that the creep mechanisms are different from each condition. The improvement of creep rupture lives by 3 wt% Ru addition is discussed not only from the view of dislocation movement but also the ␥ phase evolution. The change of ␥/␥ lattice misfit in the initial microstructure is believed to be the key role of Ru on the high-temperature creep deformation. The larger negative lattice misfit caused by an addition of 3 wt% Ru induces smaller and more regular ␥ particles in the initial state, as well as denser dislocation networks at the ␥/␥ interface during creep. These two aspects are crucial to the high-temperature creep rupture life. In addition, a little topologically close-packed (TCP) phases are observed after creep rupture in the two experimental alloys. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Single crystal Ni-based superalloys have become the primary candidate materials for blades and vanes in aero-engine or industrial gas turbines due to their high temperature capability, superior creep resistance and good environment resistance. Hightemperature creep resistance of single crystal Ni-based superalloys has been the most concerned point in this field [1–5]. Several microstructural features have been shown to remarkably influence the creep properties of single crystal Ni-based superalloys. The sign and magnitude of ␥/␥ lattice misfit ı, usually defined as ı = 2(a␥ − a␥ )/(a␥ + a␥ ), where a␥ and a␥ are the lattice constants of ␥ and ␥ phases respectively, can affect the rate of directional coarsening as well as the orientation of the ␥ lamellae with respect to the applied stress. It has been generally accepted that the larger negative misfit is beneficial to the high-temperature creep properties of single crystal Ni-based superalloys [6,7]. In addition, the size and morphology of the initial ␥ particles are also important. However, their influences are still controversial. It is believed that both the thickness and degree of perfection of the lamellar structure are important factors influencing the high-temperature creep properties of single crystal superalloys [8–10]. On the development of single crystal Ni-based superalloys, of note is the introduction of significant additions of Re, and the
∗ Corresponding author. Tel.: +86 24 2397 1757; fax: +86 24 23992092. E-mail address: [email protected] (T. Jin). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.07.075
reduction of Ti and Mo to low concentrations [11,12]. It is well accepted that Re plays a key role in improving the creep properties of these materials. Thus, Re has been regarded as the significant symbol to distinguish the first three generations single crystal superalloys. On the one hand, the addition of Re not only impedes the coarsening of ␥ particles, but also stabilizes the ␥ rafted structures due to its sluggish diffusion [13–16]. Moreover, the addition of Re also makes the lattice misfit more negative which is recognized to accelerate the formation of dense dislocation networks at ␥/␥ interfaces due to its preference into ␥ matrix [13,17–19]. All of these are helpful to improve the creep properties. On the other hand, the addition of Re makes the microstructural homogenization during solution heat treatment more difficult due to its low diffusivity and severe microsegregation [20,21]. The precipitation of topologically close-packed (TCP) phases caused by excessive Re additions is the most concerned. It is known that many detrimental effects can be caused by the precipitation of TCP phases, i.e., the depletion of important strengthening elements from the matrix, the cracks initiation induced by the stress concentration, the interruption of the microstructural continuity, etc., which would degrade the creep properties apparently [22–25]. It has been known that the microstructural instability caused by the TCP precipitation limits the further development of single crystal superalloys to fulfill higher requirements. In recent years, however, another significant alloying element Ru has been added in order to further improve the creep properties based on the third generation single crystal superalloys [26,27]. Thus, Ru has become the symbol element for the new generation
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Fig. 1. Initial microstructures of (a) alloy 0Ru and (b) alloy 3Ru.
Table 1 Nominal compositions of the two experimental alloys (wt%). Alloy
Ni
Al
Co
Cr + W + Mo + Ta
Re
Ru
0Ru 3Ru
Bal. Bal.
6 6
12 12
20 20
5 5
0 3
single crystal superalloys. It is generally accepted that the key roles displayed by Ru additions are the promotion of the microstructural stability and improvement of the creep resistance. So far it should say that the knowledge concerning the effects of Ru is far from enough, especially for the mechanisms of effects, and the acquired results still remain debatable [28–34]. This paper aims at elucidating the effect of Ru on the high-temperature creep rupture lives and giving some helpful guidelines for designing new generations single crystal superalloys. 2. Experimental Two single crystal Ni-based superalloys with and without Ru addition were designed to study the effect of Ru on the high-temperature creep rupture lives. The nominal chemical compositions of the alloys are listed in Table 1. According to their different Ru content, these two alloys are named as 0Ru and 3Ru respectively. The master alloys were melted by vacuum induction technique, and then directionally solidified into cylindrical bars (16 mm in diameter and 220 mm long) in an investment casting cluster mold in a Bridgman furnace with a withdrawal rate of 6 mm/min. Conventional helical starters were utilized to initiate single crystal growth. The heat treatment schemes for both alloys are given in Table 2. Specimens for creep tests with a diameter of 5 mm and a length of 25 mm parallel to the solidification direction [0 0 1] were machined from the heat treated alloys by means of electrosparking. The creep tests were carried out at 1000 ◦ C/310 MPa and 1100 ◦ C/150 MPa. Additionally, in an attempt to obtain the equilibrium dislocation networks for the estimation of unconstrained misfit at high temperature, long-term isothermal aging experiments were carried out. Both alloys were exposed at 1100 ◦ C for 1000 h to fully relieve the misfit stress and form equilibrium dislocation networks. The etchant used for microstructural observation is HCl (80 ml) + CuSO4 (20 g) + H2 O (100 ml) solution. A JMS-6301F fieldemission scanning electron microscope (SEM) was used to examine
the microstructure after full heat treatment and creep rupture tests. The fully heat treated and long-term aging specimens were cut into 3 mm discs parallel to (0 0 1) plane with 500 m in thickness and thinned down to 50 m mechanically. After creep tests, the specimens were cut into 3 mm discs of 500 m in thickness normal to the [0 0 1] direction about 3 mm away from the fracture surface and thinned down to 50 m mechanically. They were then electrochemically polished by the twin-jet method, in an solution of 8% perchloric acid and 92% ethanol at −10 ◦ C and 20–30 mA. A TECNAI 20 transmission electron microscope (TEM) equipped with energy dispersive spectrometer (EDS) was used to examine the dislocation configuration after long-term aging and creep rupture. Furthermore, the partitioning behavior of various alloying elements between ␥ matrix and ␥ precipitate in the initial microstructure was characterized by measuring the local chemical composition of ␥ and ␥ in fully heat treated specimens using an EDS in TEM. An average value for three times measurements was adopted. The partitioning ratio k or k defined as ki = Ci␥ /C␥ or ki = Ci /Ci was widely used. Therefore, it can be known that ki > 1 indicates the preferential partition to ␥ matrix for the alloying element, conversely, ki > 1 indicates the preferential partition to ␥ precipitate. 3. Results 3.1. Initial microstructure Fig. 1 shows the microstructures of both alloys after full heat treatment. The differences between alloy 0Ru and 3Ru on the morphology of the ␥ particles can be observed clearly. According to the image analysis results (Fig. 2), the ␥ size and its volume fraction
Table 2 Heat treatment schemes of the two experimental alloys. Alloy
Full heat treatment scheme
0Ru 3Ru
1325 ◦ C/16 h + 1335 ◦ C/16 h, ACa + 1150 ◦ C/4 h, AC + 870 ◦ C/24 h 1315 ◦ C/16 h + 1325 ◦ C/16 h, AC + 1150 ◦ C/4 h, AC + 870 ◦ C/24 h
a
The air cooling.
Fig. 2. Comparison of ␥ size and its volume fraction in the two experimental alloys.
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Fig. 3. Variation of the ␥/␥ partitioning ratio of each alloying element in the two experimental alloys.
were decreased, the ␥ shape more regular and the distribution of ␥ particles more uniform with 3 wt% Ru addition. Fig. 3 illustrates that all the alloying elements show the so-called “reverse partitioning” behavior with 3 wt% Ru addition. It means that Ru promotes the TCP-forming elements, such as Re, W, Mo, Cr etc., to partition to ␥ phase. It is believed that this effect may be the primary reason for the suppression of TCP phases by Ru additions [35,36]. As measured by the X-ray diffraction method, the average ␥/␥ lattice misfits of alloy 0Ru and 3Ru are −0.15% and −0.25% at room temperature, respectively. The detailed measurement can be seen in [37]. There is no doubt that the addition of 3 wt% Ru made the lattice misfit more negative. 3.2. Creep rupture life Fig. 4 shows the creep curves of both alloys under the two conditions. It shows typical high-temperature creep curves with three obvious stages under the condition of 1100 ◦ C/150 MPa; however, it differs at 1000 ◦ C/310 MPa, which shows shorter primary and secondary stages compared with its tertiary stage. It is clear that alloy 3Ru has longer creep rupture lives compared with alloy 0Ru under both conditions. 3.3. Microstructure after creep rupture Due to the short creep rupture lives less than 200 h for all the cases, only a little small TCP phases were observed after creep rupture. As shown in Fig. 5, there exist a certain amount of rectangular
Fig. 4. Creep rupture curves of the two experimental alloys at (a) 1100 ◦ C/150 MPa and (b) 1000 ◦ C/310 MPa.
pores in the cross-section microstructure near the fracture surfaces. Fig. 6 shows these pores with the shape of square on transversesection and rectangular on cross-section in detail. Therefore, they have a shape of rectangular parallelepiped. In addition, it was found that there must be some small TCP phases around every pore. TEM results for the phase identification indicated that these small TCP particles are phase. Thus, it is believed that the formation of these pores should be associated with TCP phases. However, it shows that the creep rupture lives for both the alloys were less dependent on these rectangular pores. It may be due to the fact that most of the
Fig. 5. Longitudinal microstructures near the fracture surface at1100 ◦ C/150 MPa of (a) alloy 0Ru and (b) alloy 3Ru.
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Fig. 6. Backscattered electron micrographs showing microstructures of alloy 0Ru after creep rupture at 1100 ◦ C/150 MPa of (a) transverse-section and (b) cross-section.
fatal cracks were not directly induced by these small rectangular parallelepiped pores. The rafted microstructures of both alloys after creep rupture under the conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa are shown in Fig. 7. It shows an evident “topological inversion” phenomenon [38], namely ␥ phase acts as the matrix and ␥ phase was surrounded by them. All the rafted ␥ lamellae were normal to the applied stress. It indicates that the lattice misfits of both alloys are negative at high temperature. The dislocation networks of both alloys after creep rupture under the conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa are shown in Fig. 8. It can be clearly seen that the dislocation networks were denser in alloy 3Ru than in 0Ru under both conditions. Furthermore, Ru promotes the formation of square dislocation networks. 4. Discussion 4.1. / portioning ratio and / lattice misfit When Ru caused a small amount of ␥-forming elements to partition to ␥ phase, they might replace some ␥ -forming elements in ␥
phase. Some ␥ -forming elements, in the meantime, would enter ␥ phase. Thus, it is reasonable for almost all the elements to show the behavior of reverse partitioning with Ru additions. As known, the ␥/␥ lattice misfit depends on not only the partitioning ratio but also the Vegard law [12]. Re does not have a larger Vegard coefficient [12]; however, it has a larger partitioning ratio as shown in Fig. 3. Therefore, Re acted as an effective element to make the ␥/␥ lattice misfit larger negative. It should be noted that Ru slightly preferentially partitioned to ␥ phase, which means that the addition of Ru must make the lattice misfit more negative. On the other hand, Ta has a much larger Vegard coefficient in ␥ phase than ␥ phase [12]. In the present work, Ru made slightly more Ta enter ␥ phase, which is also believed to make the lattice misfit more negative. Based on the two effects mentioned above, the lattice misfit became more negative with 3 wt% Ru addition. Hence, the influence of Ru on the lattice misfit is achieved by changing the partitioning behavior of other alloying elements. As reported in the previous study [37], there will exist obvious lattice distortion at the ␥/␥ interfaces in the initial microstructure of alloy 3Ru. Therefore, the lattice misfit mentioned here is the constrained misfit which would be highly dependent on the partitioning behavior of alloying elements.
Fig. 7. Secondary electron micrographs showing ␥/␥ microstructures near the fracture surface after creep rupture in (a) alloy 0Ru and (b) alloy 3Ru at 1100 ◦ C/150 MPa; and (c) alloy 0Ru and (d) alloy 3Ru at 1000 ◦ C/310 MPa.
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Fig. 8. TEM images of the interfacial dislocation networks after creep rupture in (a) alloy 0Ru and (b) alloy 3Ru at 1100 ◦ C/150 MPa; and (c) alloy 0Ru and (d) alloy 3Ru at 1000 ◦ C/310 MPa.
4.2. size, volume fraction and / lattice misfit The morphology of ␥ particles depends on both the interfacial energy and elastic strain energy, which are related with the interfacial area and volume of precipitate, respectively [39]. For the alloy with smaller magnitude of lattice misfit, interfacial energy predominates and spherical ␥ particles are shown due to their minimum surface area with the same volume. Inversely, for the alloy with larger magnitude of lattice misfit, elastic strain energy induced by lattice misfit predominates and cuboidal ␥ particles are exhibited due to the crystalline anisotropy and lowest elastic modulus with [0 0 1] orientation. The cuboidal degree of ␥ particles increases with increasing magnitude of lattice misfit. Furthermore, in terms of the characteristics of cuboidal ␥ precipitation, the degree of parallelism of ␥/␥ interfaces is improved with the increase of magnitude
of lattice misfit. Consequently, the distribution of ␥ particles is more uniform in alloy 3Ru. As we know, ␥ particles coherently precipitate from ( matrix in single crystal Ni-based superalloys. Therefore, the total elastic strain energy depends on the morphologies and elastic properties of ␥ and ␥ phases. It is assumed that ( matrix is elastically isotropic and the elastic modulus of ( matrix is equal to that of ␥ phase, and therefore the total elastic strain energy Gs is not related with the morphology of ␥ particles. Thus [39] Gs ∝ ı2 V
(1)
where is shear modulus of matrix, ı is lattice misfit and V is the volume of ␥ particle. It can be known that the elastic strain energy induced by lattice misfit is proportional to the square of lattice misfit and the volume of ␥ particle. Consequently, when the lattice
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Fig. 9. TEM images of the interfacial dislocation networks after 1000 h isothermal aging at 1100 ◦ C in (a) alloy 0Ru and (b) alloy 3Ru.
misfit of the alloy is the more negative, the volume of ␥ particle will be smaller, namely, the ␥ size is smaller. As mentioned previously, alloy 3Ru has a larger negative lattice misfit than alloy 0Ru. Hence the microstructure with smaller cuboidal and more uniform ␥ particles was shown in alloy 3Ru compared with alloy 0Ru. Moreover, it shows the tendency of lower ␥ volume fraction with Ru addition. It could be expected that the variation of initial microstructure caused by Ru addition should be the dominating factor for various hightemperature creep properties. That will be discussed in Section 4.4 in detail. 4.3. / lattice misfit and dislocation networks It is well understood that the magnitude of the lattice misfit controls the density of the interfacial dislocations required to relieve the misfit stresses. Assuming that the network dislocations are orientated to most efficiently relieve the misfit, the lattice misfit is inversely related to the spacing of the interfacial networks (D) such that [40,41] D=
៝ |b| |ı|
(2)
៝ is the magnitude of Burgers vector of the edge compowhere |b| nent and ı is the lattice misfit. Therefore, the average spacing of dislocations within the equilibrium interfacial networks provides a reasonably good estimate of the high-temperature lattice misfit at the temperature which they are formed. The distribution of equilibrium dislocation networks that formed at 1100 ◦ C in both alloys ៝ 0.254 nm and is shown in Fig. 9. Assuming a Burges vector of, |b|, using the dislocation spacing D, the magnitude of the lattice misfit can be estimated at 1100 ◦ C. The average dislocation spacings of the equilibrium interfacial dislocation networks after 1100 ◦ C isothermal aging and the interfacial dislocation networks after creep rupture under the condition of 1100 ◦ C/150 MPa are plotted with the corresponding estimated magnitude of lattice misfit in Fig. 10. As mentioned previously, the interfacial dislocation networks of alloy 3Ru were denser after creep rupture compared with alloy 0Ru. It is firmly believed that the denser interfacial dislocation networks would be an effective barrier to the shearing of ␥ lamellae [19]. It explains that the alloy with 3 wt% Ru addition has longer creep rupture lives from the view of dislocation movement. Additionally, it is easily known that the differences of longitudinal coordinate between the two conditions in Fig. 10 indicate the magnitude of so-called excess dislocation density during creep deformation. The multiplication of mobile dislocations occurs during the tertiary stage of creep deformation at high temperature, which means that the number of dislocations after creep rupture must exceed that of the misfit dislocations necessary to relieve the misfit. Their
differences are the so-called excess dislocations. In the present work, the most creep resistant alloy, 3Ru, had the lower density of excess dislocations. It is opposite, however, with the results obtained by Carroll et al. [41]. That is due to the fact that Ru additions promoted the formation of square dislocation networks. It has been reported that the square configurations of four dislocations of Burges vector of a/2 1 1 0 type are of edge character and thus efficient at relieving the misfit, but these segments have no shear stress resolved on them – thus they must have been formed by reactions driven by reductions in misfit and dislocation line energies [12]. As the remarkable increase of the magnitude of lattice misfit with 3 wt% Ru addition, it gave rise to a great driving force for the formation of the square dislocation networks. Therefore, it shows a lower excess dislocation density in alloy 3Ru. Furthermore, the regular square dislocation networks are more stable and act efficiently as a hindrance to interfacial advance in ␥ coarsening [42], which also can explain why the 3 wt% Ru addition improve the high-temperature creep rupture lives form the view of dislocation movement. 4.4. High-temperature creep rupture lives In terms of the creep curves illustrated in Fig. 4, the time spans of secondary and tertiary stages of each curve, as well as their percentage in the whole creep rupture lives, are listed in Table 3. It can be clearly seen that the secondary stage accounts for the greatest proportion during the whole creep rupture life under the condition
Fig. 10. Plot of the magnitude of the lattice misfit in the two experimental alloys determined from high-temperature equilibrium dislocation network spacings, including the postcrept dislocation network spacings at the corresponding lattice misfit.
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Table 3 The time spans of secondary and tertiary stages and their percentage in the whole creep rupture life of each creep curve under both conditions.
◦
1100 C/150MPa 1000 ◦ C/310MPa a
Alloy
tf a (h)
ts a (h)
tt a (h)
ts /tf
tt /tf
0Ru 3Ru 0Ru 3Ru
156.9 211.2 80.1 127.1
77.9 105.4 18.4 40.0
55.7 75.5 54.5 74.2
49.6% 49.9% 23.0% 31.5%
35.5% 35.7% 68.0% 58.4%
tf , ts and tt denote the creep rupture life, the time span of secondary and tertiary stage respectively.
Fig. 11. Schematic illustration of the ␥ evolution during the creep at high temperatures/low stresses. ␥ and ␥ phases are visualized as bright and dark regions, respectively.
of 1100 ◦ C/150 MPa, and the tertiary stage does at 1000 ◦ C/310 MPa. It indicates that the creep mechanism might be different between these two cases. Therefore, the creep properties under these two conditions should be discussed respectively. The mechanism of creep deformation in 0 0 1-orientated single crystal superalloys is very sensitive to the conditions of temperature and stress. Usually three typical creep regimes can be identified, i.e., the primary creep regime occurred at low temperatures and high stresses, the tertiary creep regime at intermediate temperatures and intermediate stresses and the rafting regime at high temperatures and low stresses, and in each a distinct mode of microstructural degradation is predominant [12]. In the rafting regime, the creep curve exhibits typically three different stages as shown in Fig. 11, and the ␥/␥ microstructure degrades rather quickly because thermally activated processes are favored strongly since the temperature is high. Therefore, it can be termed as the diffusion-controlled creep. During the short primary stage, rafting of ␥ particles occurs and dislocation networks required to relieve the misfit stress develop at the ␥/␥ interfaces. It is followed by a long plateau where a transition from rafting to topological inversion occurs, and the secondary stage is ended by the occurrence of
topological inversion. Afterwards, cutting of ␥ lamellae by paired dislocations begins with the onset of tertiary stage characterized by an acceleration of the deformation during a short time and driving to failure [4,43]. It is clear that the secondary stage predominates over the whole creep rupture life. In the present work of creep at 1100 ◦ C/150 MPa, it is noted that the creep rupture life mainly depends on the time of the occurrence of topological inversion. As mentioned previously, the addition of 3 wt% Ru decreases the ␥ volume fraction, which will prolong the secondary stage and eventually increases the creep rupture live. On the other hand, it has been reported [9] that more perfect ␥ lamellae are more resistant to thickening and thus would delay the onset of tertiary creep, thus prolonging the creep life of single crystal superalloys. Here the addition of 3 wt% Ru makes the ␥ particles more regular and further more perfect ␥ lamellae. Therefore, it can also increase the creep rupture life. Moreover, due to the fact that this process is diffusioncontrolled and the diffusion coefficient of Ru is second to Re [44,45], which further retards the occurrence of topological inversion, it is concluded that 3 wt% Ru addition could significantly promote the creep rupture life under this condition. In the tertiary regime, the creep curve exhibits an obvious tertiary stage as shown in Fig. 12(a), the density of dislocations is found to increase with increasing creep strain, while other forms of microstructural degradation (␥ rafting, creep cavitations) are absent. Usually slight ␥ coarsening occurs [46,47]. Therefore, it can be termed as a dislocation-controlled creep. In the present work, however, creep at 1000 ◦ C/310 MPa which corresponds to the condition of higher temperatures and high stresses, may differ from both conditions of high temperatures/low stresses and intermediate temperatures/intermediate stresses. As shown in Fig. 12(b), the creep curve exhibits three stages with a longer tertiary stage. From the current experimental results, topological inversion of ␥/␥ microstructure was observed. It means that there exists a transition from rafting to topological inversion. In addition to the ␥ evolution, it could be imagined that the cutting of ␥ lamellae by dislocations started from the onset of tertiary stage. Therefore, it can be termed as the dislocation-controlled creep coupled with the diffusion-controlled mode. Apparently the tertiary stage still predominates over the whole creep rupture life (as listed in Table 3). As a result that the creep rupture life will mainly depend on the duration of ␥ lamellae cutting process at 1000 ◦ C/310 MPa. The cutting resistance of ␥ lamellae is dependent on the thickness and
Fig. 12. Schematic illustration of the ␥ evolution during the creep at (a) intermediate temperatures/intermediate stresses and (b) higher temperatures/higher stresses. ␥ and ␥ phases are visualized as bright and dark regions, respectively.
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regularity of ␥ lamellae. As mentioned previously, more regular cuboidal ␥ particles formed in the initial microstructure by 3 wt% Ru addition. It is followed by more perfect ␥ lamellae after a short primary stage. It means that the number of imperfection sites at the ␥/␥ interfaces, where are preferentially cut by dislocations, will be accordingly reduced. Therefore, the more perfect ␥ lamellae formed in the alloy 3Ru are more resistant to dislocation cutting. On the other hand, the denser interfacial dislocation networks formed after the primary stage during creep in the alloy 3Ru could effectively improve the cutting resistance of ␥ lamellae. Thus, the creep rupture life was remarkably increased with 3 wt% Ru addition under this condition. To summarize, the initial microstructure after full heat treatment was noticeably influenced by the 3 wt% Ru addition, which appears to be the major reason for the various creep rupture lives eventually. In the present work, the change of lattice misfit was considered as the key role that Ru has played. From the view of ␥ evolution, lattice misfit influenced the morphology, size, volume fraction and distribution of ␥ particles, and the following regularity and perfection of ␥ lamellae during high-temperature creep. On the other hand, from the view of dislocation movement, it influenced the configuration and density of interfacial dislocation networks and the following dislocation movement. As the two aspects are equally important, it is necessary to correlate them for understanding the high-temperature creep properties. However, it is worth noting that the suppression of TCP phases by Ru addition was not highly stressed in this work. In spite of the occurrence of some small TCP phases, their influences on the creep rupture life could be negligible. 5. Conclusions The creep rupture lives of two single Ni-based superalloys with and without Ru addition were investigated in this paper and the conclusions are summarized as follows: 1. The addition of 3 wt% Ru improves the creep rupture lives of a single crystal Ni-based superalloy under both conditions of 1100 ◦ C/150 MPa and 1000 ◦ C/310 MPa. 2. At 1100 ◦ C/150 MPa, the secondary stage predominates over the whole creep rupture life. 3 wt% Ru addition retards the occurrence of the topological inversion, which prolongs the duration of the secondary stage. 3. At 1000 ◦ C/310 MPa, the tertiary stage predominates over the whole creep rupture life. 3 wt% Ru addition improves the cutting resistance of ␥ lamellae, which prolongs the duration of the tertiary stage.
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Acknowledgements This work was financially supported by the National Basic Research Program (973 Program) of China under grant No. 2010CB631200 (2010CB631206), the National Natural Science Foundation of China (NSFC) under grant No. 50931004. The authors are grateful for those supports.
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