ISSN 0031918X, The Physics of Metals and Metallography, 2014, Vol. 115, No. 1, pp. 21–29. © Pleiades Publishing, Ltd., 2014. Original Russian Text © A.I. Epishin, T. Link, G. Noltze, I.L. Svetlov, B.S. Bokshtein, A.O. Rodin, R. SalivanNeumann, G. Oder, 2014, published in Fizika Metallov i Metallovedenie, 2014, Vol. 115, No. 1, pp. 23–31.
STRUCTURE, PHASE TRANSFORMATIONS, AND DIFFUSION
Diffusion Processes in Multicomponent NickelBase Superalloy–Nickel System A. I. Epishina, T. Linka, G. Noltzeb, I. L. Svetlovc, B. S. Bokshteind, A. O. Rodind, R. SalivanNeumannb, and G. Oderb aBerlin
Technical University, Ernst Reuter Platz 1, Berlin, 10587 Germany Federal Institute for Materials Research and Testing (BAM), Unter den Eichen 87, Berlin, 12205 Germany cAllRussia Research Institute of Aviation Materials, ul. Radio 17, 105005 Moscow d National University of Science and Technology “MISiS”, Leninskii pr. 4, Moscow, 119049 Russia email:
[email protected]
b
Received March 5, 2013; in final form May 14, 2013
Abstract—Optical and scanning electron microscopy, as well as electron microprobe analysis and electron backscatter diffraction, have been used to study diffusion processes that occur in a diffusion pair that consists of a singlecrystal CMSX10 nickelbase superalloy and polycrystalline nickel, at temperatures of 1050– 1250°C. It has been found that, in this system, the distributions of γstabilizing elements (Cr, Co, W, and Re) are described by the Boltzmann solution for diffusion between two semiinfinite plates of a binary alloy. The pro cessing of these distributions has shown that the diffusion coefficients of Cr, Co, W, and Re in the multicompo nent system are close to those in binary alloys of these elements with Ni. The diffusion redistribution of the ele ments leads to the dissolution of the γ' phase in the nickelbase superalloy, growth of nickel grains toward the superalloy constituent of the diffusion pair, and the formation of porosity on both sides of the migrating inter face, which is determined from a crystal misorientation of the alloy single crystal and nickel grains. Keywords: nickelbase superalloys, interface, diffusion, phase transformations, porosity DOI: 10.1134/S0031918X14010050
INTRODUCTION
process. These fluxes cause undesirable phase trans formations in the vicinity of the interface and, in some cases, the formation of porosity due to the Kirken
The nickelbase superalloys (NSAs) are widely used in aviation motor building to produce parts of hot sec tions of gas turbines, in particular for casting single crystal turbine blades [1, 2]. These alloys have a com plex multicomponent chemical composition and a two phase γ/γ' microstructure. In the NSAs, the strengthen ing effect is achieved due both to the solidsolution hardening of the nickel γ matrix by the elements, such as Re, W, Co, etc., and to the precipitation hardening of an alloy by microscopic particles of the γ' phase based on the alloyed intermetallic compound Ni3Al. Impor tant factors that determine the hightemperature strength of NSAs are the phase and structural thermal stability, which are directly related to diffusion pro cesses at elevated temperatures. Therefore, the primary aim of this work was to study diffusion processes that occur in a multicomponent nickelbased system. The technology of the production of gasturbine blades involves not only their casting to form a single crystal structure, but also the deposition of corrosion and heatresistant coatings, as well as soldering and welding processes. The material adjacent to the vari ous interfaces formed has a lower thermal stability than that of the NSA itself because of the existing con centration gradients and diffusion fluxes of the alloy ing elements that arise in the course of the production
1
dall–Frenkel mechanism , i.e., the appearance of defects that initiate crack nucleation and premature fracture, especially under cyclic loading. A similar poreformation process occurs in the case of the homogenization of single crystals of NSAs at the inter face between the γ matrix and nonequilibrium eutectic precipitates (γ/γ') [5]. Therefore, the second purpose of this work was to study the interrelation of the diffu sion processes with the phase transformations and pore formation in an NSA–Ni model system. EXPERIMENTAL The objects of the study were CMSX10 NSA–Ni diffusion pairs. The CMSX10 alloy was developed by the Cannon Muskegon Corp. for the casting of uncooled singlecrystal blades of the highpressure turbine of the Trent 800 engine [6]. The chemical composition of the CMSX10 alloy (Table 1) is char acterized by high concentrations of refractory ele 1A
review concerning the Kirkendall effect is presented in [3]. The Frenkel effect, i.e., the formation of porosity when Kirken dall effect is observed, is considered in [4].
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EPISHIN et al.
Table 1. Chemical composition of CMSX10 nickelbase superalloy [6] Element Wt % At %
Al
Ti
Cr
Co
Ni
Nb
Mo
Hf
Ta
W
Re
5.7 13.2
0.2 0.3
2 2.4
3 3.2
Base Base
0.1 0.06
0.4 0.3
0.03 0.01
8 2.8
5 1.7
6 2.0
ments. In particular, this alloy contains 6 wt % Re and, according to the international classification, belongs to the third generation of singlecrystal NSAs. The singlecrystal ingots of the CMSX10 alloy were pro duced by the Bridgman–Stockbarger method in a fur nace used for the commercial production of single crystal blades. The singlecrystals solidified under these conditions have a block crystal structure [7] and are chemically inhomogeneous because of dendritic segregation. The ingots were subjected to usual heat treatment, which included tenstep homogenization in the temperature range of 1315–1365°C for 45 h [8], as well as aging at 1150°C for 4 h, at 870°C for 24 h, and at 760°C for 30 h. After the complete cycle of heat treatment, the alloy has a cuboidal γ/γ' microstructure in which γ' particles about 0.5 μm in size are separated by thin γphase interlayers ≈0.05 μm thick. According to [9], the fraction of the γ' phase at room temperature is approximately 80%. The temperature dependence of the dissolution of the γ' phase in the CMSX10 alloy was measured by the resistivity method [10]. The mea surements were carried out using a TER2000 setup (Ulvac, Sinku Riko); the results are presented in Fig. 1. According to these results, the solvus tempera ture of the γ' phase in the CMSX10 alloy is 1342 ± 2°С. As the nickel component of the diffusion pairs, highpurity Ni produced by zone melting was used. 1350
γ' solvus, 1342°С 31%
Temperature, °С
1250
55%
1150
69%
1050
γ'
950
γ
850 0
20 40 60 80 Volume fraction of γ' phase, %
Fig. 1. Temperature dependence of volume fraction of γ' phase in CMSX10 alloy.
100
The specimens had a columnar grain structure with a transverse grain size of a few millimeters and low intra granular mosaicity. The diffusion pairs were made of disks (~2.5 mm thick) cut out of ingots of the CMSX10 alloy and highpurity Ni perpendicularly to the ingot axes. Both sides of the disk were initially ground using abrasive Emery papers; then, one side was polished using dia mond pastes with a grain size of down to 1 μm. The polished surfaces of the disks were joined by diffusion welding in an MTS810 testing machine using induc tion heating in a vacuum of about 10–6 mbar at a tem perature of 1050°С under a compressive stress of 10 MPa for 1 h. After welding, the specimens were cut and the quality of the weld was inspected using optical microscope and scanning electron microscope (SEM). Three diffusion pairs were annealed in a vac uum at 1050°C for 3972 h (128 days), at 1150°C for 768 h (32 days), and at 1250°C for 192 h (8 days). These annealing conditions were selected for the fol lowing reasons. First, the width of the diffusion zone of the least mobile element (Re) should be acceptable for measuring concentration profiles using electron microprobe analysis (EMA), i.e., no less than 100 μm. Second, the total volume of diffusion mass transfer at various temperatures should be approximately the same. Note that, at these annealing temperatures, the CMSX10 alloy has a twophase γ/γ' microstructure with a fraction of the γ' phase of ≈69, ≈55, and ≈31 vol % at 1050, 1150, and 1250°C, respectively (Fig. 1). After annealing, the specimens were cut across the interface and prepared for study by optical metallography. Microstructure was studied using a Zeiss Axioscope optical electron microscope and a LEO GEMINI 1530 VP SEM equipped with the following Bruker analytical systems: an EDS XFlash 5030 energy dis persive detector of characteristic Xray radiation; and an eFlash HR detector of electron backscatter diffraction (EBSD) with ESPRIT and CrystAlign software. The concentration profiles were recorded by a JEOL JXA 8900 wavelengthdispersive electron microprobe analyzer. The electronmicroprobe mea surements were performed in the direction perpendic ular to the interface with a step of 5 μm. To reduce local fluctuations of the measured concentrations, the irradiated pointlike area was extended by defocusing the electron beam to a 5 μm diameter. The concentration profiles were approximated using Boltzmann’s solution for the diffusion between two semiinfinite plates of a binary alloy as follows: x − x0 ⎞ C − C1 (1) C(x, t ) = C1 + 2 erfc ⎛⎜ , ⎟⎠ 2 ⎝ 2 Dt
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DIFFUSION PROCESSES IN MULTICOMPONENT … SYSTEM Weld interface Ni
γsolid solution
γ/γ'microstructure
10 MPa
Ni
CMSX10
CMSX10
23
Primary porosity 10 MPa OI Al Ni
4 μm 200 μm
(a) Fig. 2. Structure of CMSX10–Ni pair near original inter face (Kirkendall plane) after diffusion welding for 1 h at 1050°С at a pressure of 10 MPa (EBSD image). In region where outflow of light Al occurs, the brightness of the image increases due to the growth in the average atomic number.
Ni
CMSX10
where C1 and C2 are the initial concentrations of one of the elements to the left and to the right of the interface, D is the interdiffusion coefficient, x is the coordinate in the direction perpendicular to the interface, x0 is the coordinate of the interface, and t is the annealing time. RESULTS AND DISCUSSION Structure of the Diffusion Pairs after Diffusion Welding and Long Annealings Figure 2 shows the structure of a diffusion pair after diffusion welding. It can be seen that, even in the course of welding, a slight diffusion interaction occurred between the NSA and nickel: the region to the left of the interface is a typical twophase γ/γ' microstructure of the NSA; the region to the right of the interface is a nickel grain. These regions are sepa rated by a bright singlephase region of the γ solid solution that appeared in the NSA due to the dissolu tion of γ'phase particles (dark cuboidal particles) and the rapid diffusion outflow of aluminum into nickel. After welding for 1 h at 1050°C, the width Lγ of the γ region is about 4 μm. Since the aluminum flux is not compensated for by the slower opposite nickel flux, a chain of primary pores less than 1 μm in size is formed at the interface due to the Kirkendall–Frenkel effect. Below, we will use this characteristic configuration of pores as a marker to identify the initial coordinate of the interface, which is usually called the Kirkendall plane. Figure 3 shows the optical images of the structure of the pairs under study after diffusion annealings per formed at temperatures of 1050, 1150, and 1250°C of THE PHYSICS OF METALS AND METALLOGRAPHY
OI
200 μm
(b)
Ni
CMSX10
OI
20 μm
(c)
Fig. 3. Porosity near interface in CMSX10–Ni pair after diffusion annealing under various conditions: (a) 1050°C, 3072 h; (b) 1150°C, 768 h; and (c) 1250°C, 192 h. OI is the original interface. Vol. 115
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EPISHIN et al. (a)
20 μm (b) Single crystal CMSX10
Ni grain
[111]
HAB
OI CMSX10 Ni
ΔLNi
200 μm
[001]
[011]
Fig. 4. Structure of the CMSX10–Ni diffusion pairs after annealing for 768 h at 1150°C: (a) pores at the front of the dissolution of the γ' phase in the CMSX10 alloy; (b) EBSD, crystallattice orientation map. Orientations of CMSX10 alloy and Ni are shown in the righthand part of the figure in the stereographic triangle. HAB is highangle boundary; OI is original interface.
various durations. It can be seen that, after annealing, the size of the primary pores increased to 10–50 μm. In addition, a large number of chaotically arranged secondary pores of a similar size appeared. An SEM examination has shown that these pores are located in the γ region, which expanded many times during the diffusion annealings. The width Lγ of the γ region depends on the annealing conditions; it is equal to 215, 300, and 500 μm for the annealings carried out at 1050°C for 3072 h, 1150°C for 768 h, and 1250°C for 192 h, respectively. It has been found that large sec ondary pores are formed directly at the front of disso lution of the γ' phase (Fig. 4a).
It has been shown using EBSD that the crystallo graphic orientation of the NSA single crystal in the direction perpendicular to the interface is close to [001], while the orientations of nickel columnar grains deviate by 10°–20° from this direction. Thus, the alloysinglecrystal–nickelgrain interface is a high angle boundary. Fig. 4b shows an orientation map of the considered pair subjected to annealing at 1150°C for 768 h and the stereographic triangle with the orien tations of the CMSX10 alloy and nickel. It can be seen from the difference in the orientationcontrast levels that, in the course of diffusion annealing, the highangle boundary shifts to the left of its initial posi tion due to nickelgrain growth. The shift ΔLNi is equal
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Flow, 10–6, mol/(m2 s)
3
СMSX10
25 Ni
JAl
2 Jv
1 0 –1
JNi γ/γ'
–2
γregion
Nickel grain
Jv –3
Concentration, at %
OI
HAB
4
CNi/25
CCo
3
2
CRe CAl/5
1 Plateau 0 –6
–5
–4 –3 –2 –1 Distance from original interface x, 100 μm
0
1
2
3
Fig. 5. Concentrationdistribution profiles for Al, Co, Ni, and Re (at the bottom) and calculated fluxes of Al, Ni, and vacancies (at the top). To present profiles for Al and Ni, as well as for Co and Re, in one plot, the concentrations of Al and Ni were decreased by dividing by 5 and 25, respectively. HAB is highangle boundary; OI is original interface.
to ≈50, ≈63, and ≈72 μm for the annealings carried out at 1050°C for 3072 h, 1150°C for 768 h, and 1250°C for 192 h, respectively. Taking into account these results, we can draw the following conclusion on pore formation: the primary pores have nucleated at the original interface, but their subsequent growth occurred in the bulk of a nickel grain. Shape of Concentration Profiles and Diffusion Parameters The concentration profiles of Al, Co, Ni, and Re in the diffusion pair subjected to annealing at 1150°C for 768°C are shown at the bottom of Fig. 5. The origin of coordinates (x = 0) was assumed to coincide with the position of the initial interface. THE PHYSICS OF METALS AND METALLOGRAPHY
It can be seen from Fig. 5 that the shape of the con centration profiles depends on the type of the alloying element. The shapes of the concentration profiles of the γstabilizing elements (Cr, Co, W, and Re) are close to that of the profile that results from Boltz mann’s solution (1) (see, e.g., the experimental pro files of Co and Re (the fluctuating curve), as well as their approximation by Boltzmann’s solution (the smooth curve)). The processing of the profiles of the γ stabilizing elements using solution (1) has allowed us to determine the effective interdiffusion coefficients D of these elements and nickel, as well as to calculate the corresponding preexponential factors D0 and activa tion energies Q from the Arrhenius equation D (T ) = D0exp(–Q/RT), where T and R are the absolute tem perature and the universal gas constant, respectively. Vol. 115
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Table 2. Effective interdiffusion coefficients D (m2/s), preexponential factors D0 (m2/s), and activation energies Q (kJ/mol) cal culated for Co, Cr, W, and Re using Boltzmann’s solution Parameter
Cr
Co
W
Re
D (1050°C)
3.3 × 10–15
1.4 × 10–15
1.5 × 10–16
4.6 × 10–17
D (1150°C)
1.6 × 10–14
6.7 × 10–15
9.8 × 10–16
2.9 × 10–16
D (1250°C) D0
7.2 × 10–14
3.4 × 10–14
6.1 × 10–15
2.1 × 10–15
3.4 × 10–5
4.8 × 10–5
2.5 × 10–4
1.7 × 10–4
250 ± 5
270 ± 5
310 ± 5
320 ± 5
Q
The results are presented in Table 2. The relative error of measuring the interdiffusion coefficients is no more than 5%. The obtained values D0 ∼ 10–5–10–4 m2/s are typical of substitutional alloying elements in fcc met als [11]. An increase in the values of Q in the Cr, Co, W, and Re sequence (250, 270, 310, and 320 kJ/mol, respectively) also seems reasonable. The graphs of the obtained D = f(T) dependences are presented in Fig. 6. The points and the solid lines show our results for the diffusion pairs under study; the fainter lines show the results of studying Ni–X binary alloys. It can be seen from Fig. 6 that the D = f(T) dependences for the diffusion of Cr, Co, W, and Re in the multicompo nent system under consideration are similar to those for binary alloys of these elements with Ni. The values of the diffusion parameters of Re (which plays an important role in enhancing the hightemperature
Interdiffusion coefficient, m2/s
1E–13
Cr
1E–14 Co
W
1E–15
Re 1E–16
6.6
6.8 7.0 7.2 7.4 7.6 Reciprocal temperature 104/T, K–1
Fig. 6. Temperature dependence of interdiffusion coeffi cients D for Cr, Co, W, and Re. Points and solid lines show our results; fainter lines show results for Ni–X binary alloys: Ni–Re (other our work); Ni–Cr, Ni–Co, and Ni⎯W [12].
strength of modern NSAs [13]) obtained in this work are almost the same for the multicomponent system and the Ni–Re binary alloy, i.e., D0 = 1.66 × 10–4 and 1.16 × 10–4 m2/s and Q = 320 and 317 kJ/mol, respec tively. It should be noted that the observed diffusion of the γ stabilizing elements occurs in the γ solid solution, since the rate of the expansion of the γ region, which is controlled by the outflow of movable Al from the NSA, exceeds the rate of the outflow of the slower γstabilizing elements. The most complex shape of the concentration pro file is characteristic of the main γ'forming element, i.e., Al (see Fig. 5). In the twophase γ/γ' region (shown by dark gray in Fig. 5), the concentration pro file of Al fluctuates noticeably, since the size of the γ' phase particles, which has become coarsened during annealing, is comparable with the electronbeam diameter. In this region, the concentration of alumi num CAl corresponds to that in the CMSX10 alloy, i.e., ≈13–14 at %. In the singlephase γ region (shown by light gray in Fig. 5), the concentration profile is smoother. At the interface between the singlephase and twophase regions, where the intense dissolution of the γ' phase occurs, the concentration of aluminum corresponds to its solubility limit in the alloy with a γ ≈ 10.3 at %) at 1150°С. given local composition (C Al Similar to a decrease in the solubility of Al in γNi in the Ni–Al binary system [14], the limiting value of the γ aluminum concentration C Al decreases with decreas ing annealing temperature; it is ≈12.1, ≈10.3, and ≈8.4 at % at 1250, 1150, and 1050°С, respectively. In the singlephase region,СAl decreases monotonically. Going from the NSA single crystal into the growing nickel grain (shown by white in Fig. 5), the rate of the decrease in the concentration of Al drops sharply and a plateau appears, the width of which is equal to the increment in the nickel grain size ΔLNi. Outside this plateau, the concentration of Al continues to decrease monotonically. The concentration profile of nickel is more monotonic than that of aluminum, but it has two inflection points, one near the boundary of the grow ing nickel grain and one at the front of the dissolution of the γ' phase.
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Interrelation between Diffusion and Pore Formation
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10
8 1150°C
1050°C
6
ΔC
Al
1250°C p
=
4
V
Porosity, vol %
To explain the processes of pore formation in this system, we analyzed the diffusion fluxes of Ni and Al, which have the largest atomic concentration and high diffusion mobility. To determine these fluxes, the experimental concentration profiles were smoothed and differentiated, and the calculated partial deriva tives ∂ C ∂ x were substituted into the expression for the first Fick’s law J = −D ∂C ∂x . The diffusion coeffi cients DNi and DAl were calculated using the tempera ture dependences of the impurity concentrations for the Ni–Al binary alloy [15]; to represent the fluxes in mol/(m2 s), the data were normalized to the molar vol ume of Ni (6.59 × 10–6 m3/mol). Note that this calcu lation is only valid within the γ region of the NSA and in γNi, but not in the region of the γ/γ' microstruc ture, since the diffusion coefficients of Al and Ni in the γ' phase are much lower than those in the γ phase [16]. Following Bardeen and Herring [17], we assume that, in local regions, the difference between the fluxes of Ni and Al is compensated for by the inflow of vacan cies Jν = − ( J Ni + J Al ) . The results of calculating JNi, JAl, and Jν for annealing carried out at 1150°C for 768 h are presented in the upper part of Fig. 5. It can be seen that, at the front of the dissolution of the γ' phase (x ≈ –360 µm), the rapid outflow of Al toward Ni occurs. In this region, the flux JAl exceeds by almost an order of magnitude the counterflow JNi, which causes a powerful flux of vacancies Jν from the γ region toward the front of the dissolution of the γ' phase, i.e., from the right to the left, according to the negative sign of Jν. The inflow of the vacancies and their condensa tion at the front of the dissolution of the γ' phase lead to an intense increase in the secondary porosity (Fig. 4a). However, with an increase in the distance from the front of the dissolution of the γ' phase, the flux of vacancies Jν decreases and the rate of the pore growth diminishes. The existence of an interrelation between the pore formation in the NSA and the diffu sion outflow of Al is also confirmed by the existence of a direct dependence of the volume fracture of second ary pores Vp on the decrease in the average atomic concentration of Al in this region (ΔCAl) in compari son with CAl in the original CMSX10 alloy (Vp ≈ ΔCAl, see Fig. 7). It should be noted that a similar process of pore formation occurs in single crystals of NSA during their homogenization. The authors of [5] suggested that the homogenizationrelated porosity resulted from the existence of unbalanced fluxes of atoms. Later, quantitative models of pore formation were developed [18–20] based on the assumption that all generated vacancies condensed on the surface of the pores. Returning back to the upper part of Fig. 5, note that, in the region between the highangle and original boundaries, the inversion of the intensities of the dif fusion fluxes of Al and Ni takes place. In this region,
27
2
0
2 4 6 8 Decrease in concentration of Al, at %
10
Fig. 7. Dependence of volume fraction of secondary porosity on decrease in average atomic concentration of Al in γ region for CMSX10 alloy.
the flux of Ni has the maximum value. In the region of the concentration plateau, the flux of Al is close to zero. The disbalance of the diffusion fluxes of nickel and aluminum JNi + JAl is compensated for by the flux of vacancies. Because of the positive sign of Jν, this flux is directed from the left to the right, and these vacan cies are apparently generated by the migrating high angle boundary. The generated vacancies diffuse toward the primary pores and condense on their sur face, thus increasing the volume of the pores. Activation Energies of Processes of NickelGrain Growth and γRegion Expansion To estimate the activation energies of the processes of growth of nickel grains and of the expansion of the γ region in the NSA, we analyzed their temperature dependences. Since these processes are diffusioncon trolled, one can expect linear dependences for the functions ΔL2Ni t = f (1 T ) and L2γ t = f (1 T ), where t is the annealing time. It can be seen from Fig. 8 that these dependences are indeed close to linear. It follows from the slope of the ΔL2Ni t = f (1 T ) graph that the activation energy of the process of nickelgrain growth is equal to 294 kJ/mol, which is close both to the acti vation energy of the selfdiffusion of Ni and to the activation energy of the diffusion of Al in Ni, which are equal to 287 and 284 kJ/mol, respectively [15]. This result seems to be reasonable. On the one hand, the growth of nickel grains is controlled by the diffu sion inflow of nickel to the migrating highangle boundary. On the other hand, this growth process also Vol. 115
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EPISHIN et al. 1E–12
Q= 372 Exp kJ/ ans mo ion l of γ Lγ (t = 1 h) reg ion in N SA Q= 2 93 kJ/m ol Gro wth of n icke l gra in
L2/t, m2/s
1E–13
1E–14
1E–15
1E–16 6.6
7.2 6.8 7.0 7.4 Reciprocal temperature 104/T, K–1
Fig. 8. Temperature dependences of kinetic parameter L2/t for processes of nickelgrain growth (L = ΔLNi) and expansion of γ region in NSA (L = Lγ). Solid point in 2 graph of Lγ t vs. 1/T graph corresponds to the γ region formed upon diffusion welding (Fig. 2).
requires an inflow of aluminum, since the growth of nickel grains is accompanied by the pickup of a defi nite concentration of this element. It follows from the slope of the curve of L2γ t vs. 1/T that the activation energy of the process of the expansion of the γ region in the NSA is equal to 372 kJ/mol, which is substan tially higher than the activation energy of the diffusion of Al in Ni. This difference can be explained by the fact that the dependence of L2γ t on 1/T is due to not only the temperature dependence of the diffusion coeffi cient of Al, but also to a decrease in the volume frac tion of the γ' phase with increasing temperature (Fig. 1). At a lower volume fraction of the γ' phase, its complete dissolution requires a weaker outflow of Al and, hence, it develops more rapidly, which increases the kinetic factor L2γ t and the slope of the curve of L2γ t vs. 1/T. CONCLUSIONS Methods of optical and scanning electron micros copy, as well as electron microprobe analysis and elec tron backscatter diffraction, have been used to study the diffusion processes that occur in the CMSX10 nickelbase superalloy–Ni pair at temperatures of 1050–1250°C. It has been found that the following main processes develop under these conditions: (1) the diffusion of the alloying elements from the NSA into Ni and the dissolution of the γ' phase;
(2) the growth of nickel grains into the NSA; (3) the formation of pores on both sides of the migrating interface according to the Kirkendall–Fren kel effect due to the dominant flux of Al on the NSA side and due to the dominant flux of Ni on the Ni side. An analysis of the concentration profiles has shown that the kinetics of the diffusion of the γstabilizing elements, such as Cr, Co, W, and Re, in the multicom ponent system considered is similar to that observed in the binary alloys of these elements with Ni. ACKNOWLEDGMENTS This work was supported in part by the Russian Foundation for Basic Research (project no. 0903 91332NNIO_a) and by the German Research Foun dation (projects nos. PO 405/121 and LI 494/51). We are grateful to Yu.V. Loshchinin for measuring the temperature dependence of the dissolution of the γ' phase by the electricalresistance method. REFERENCES 1. R. E. Shalin, I. L. Svetlov, E. B. Kachanov, V. N. Tolo raiya, and O. S. Gavrilin, Single Crystals of Nickel Superalloys (Mashinostroenie, Moscow, 1997) [in Rus sian]. 2. R. C. Reed, The Superalloys: Fundamentals and Appli cations (Cambridge University, Cambridge, 2006). 3. H. Nakajima, “The discovery and acceptance of the Kirkendall effect: The results of a short research career,” JOM 49 (1), 15–19 (1997). 4. Ya. I. Frenkel’, Introduction to the Physics of Metals (Fizmatgiz, Moscow, 1958) [In Russian]. 5. D. L. Anton and A. F. Giamei, “Porosity distribution and growth during homogenization in single crystals of a nickelbase superalloy,” Mater. Sci. Eng. 76 (12), 173–180 (1985). 6. G. L. Erickson, “The development and application of CMSX10,” in Superalloys 1996, Ed. by R. D. Kiss inger, D. J. Deye, D. L. Anton, A. D. Cetel, M. V. Nathal, T. M. Pollock, and D. A. Woodford (TMS, Warrendale, 1996), pp. 35–44. 7. U. Brückner, A. Epishin, and T. Link, “Local Xray dif fraction analysis of the structure of dendrites in single crystal nickelbase superalloys,” Acta Mater. 45, 5223– 5231 (1997). 8. G. E. Fuchs, “Solution heat treatment response of a third generation single crystal Nibase superalloy,” Mater. Sci. Eng. A 300, 52–60 (2001). 9. C. Schulze and M. FellerKniepmeier, “Transmission electron microscopy of phase composition and lattice misfit in the Recontaining nickelbase superalloy CMSX10,” Mater. Sci. Eng. A 281, 204–212 (2000). 10. B. Roebuck, D. Cox, and R. Reed, “The temperature dependence of γ' volume fraction in a Nibased single crystal superalloy from resistivity measurements,” Scr. Mater. 44, 1917–1921 (2000). 11. B. S. Bokshtein, Diffusion in Metals (Metallurgiya, Moscow, 1978), [in Russian].
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Translated by D. Tkachuk
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