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Novel Co-Ti-V-base superalloys reinforced by L12-ordered γ′ phase
Intermetallics 92 (2018) 126–132
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Novel Co-Ti-V-base superalloys reinforced by L12-ordered γ′ phase a
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J.J. Ruan , X.J. Liu , S.Y. Yang , W.W. Xu , T. Omori , T. Yang , B. Deng , H.X. Jiang , C.P. Wanga,∗, R. Kainumab, K. Ishidab a Department of Materials Science and Engineering, College of Materials and Fujian Provincial Key Laboratory of Materials Genome, Xiamen University, Xiamen, 361005, PR China b Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai, 980-8579, Japan
A R T I C L E I N F O
A B S T R A C T
Keywords: Cobalt-base superalloys Compression test L12 compound Phase diagram
The influences of alloying elements Al and Ni to a Co-Ti-V master alloy on microstructures, phase stabilities and high-temperature mechanical properties have been investigated. A two-phase microstructure γ/γʹ with ultrahigh γʹ volume fraction was designed in the Co-Ti-V (master alloy), Co-Ti-V-Al (Al-modified) and Co-Ti-V-Ni (Nimodified) alloys. Al shows a slight tendency to partition into γʹ, while Ni exhibits almost equal distribution between the γ and γʹ phases. Al is found to increase the volume fraction and solvus temperature of γʹ. The solvus temperatures of γʹ in the master alloy, Al-modified and Ni-modified alloys exceed that of Co-9Al-9W alloy (1000 °C) by 91, 112 and 82 °C, respectively. The flow stresses of the designed alloys exhibit an anomalous positive dependence on temperature rising from 600 to 750 °C. The strengths of the master alloy and Ni-modified alloys are not only higher than that of traditional Co-base superalloy MAR-M 302 at the temperature ranging from 700 to 1000 °C, but also comparable with that of the commercial Ni-base superalloy IN-939 between 850 and 1000 °C. In conclusion, the present Co-Ti-V-base alloys reinforced by γʹ are suggested as potential candidates for high-temperature utilizations.
1. Introduction The superalloys, which normally serve in severe industrial environments facing high temperature and pressure, can be generally divided into three classes named as nickel-, cobalt- and iron-base superalloys [1,2]. The melting point of nickel (∼1455 °C) is lower compared with that of cobalt (∼1495 °C) and iron (∼1538 °C) [3,4], and it partly suggests the higher possibility for cobalt- and iron-base superalloys to be utilized at evaluated temperature than Ni-base superalloys. However, the fact is that the Ni-base superalloys reinforced by coherent fcc-ordered Ni3Al (L12 structure) phase [5–8] are more widely utilized in nowadays industries than the other two types. It is believed that the lack of a coherent geometrically close-packed (GCP) phase in cobaltand iron-base superalloys is the main limitation for their applications [2]. It needs to point out that the cobalt- and iron-base superalloys are still playing a main role in some industrial fields including petrochemical industries, nuclear power plants and gas turbine industries etc. due to some reasons such as the less cost of iron-base superalloys and good resistance to high-temperature hot corrosion (HTHC) of cobalt-base superalloys [2,9]. In recent years, GCP phases were discovered in some Co-base ternary alloys: Co-Al-W [10], Co-Ge-W [11] and Co-Ga-W [12]. The
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most attractive aspects are high melting point (Tm) and good mechanical properties shown in Co-Al-W-base alloys [10]. Thus, many studies relating to the Co-Al-W-base alloys were also conducted, for instance, the influence of alloying elements, Ta, Re, Cr, Ti, Nb, Mo, W, Mn, Fe, Ni, V and B, on mechanical properties, volume fraction (Vγʹ) and solvus temperature (Tsolvus-γʹ) of γʹ of Co-Al-W-base alloy were investigated [10,13–19]. Their results suggest a great possibility of producing novel Co-base superalloys reinforced through coherent precipitate to satisfy the increasing severe demands in industries. However, the γʹ phase in Co-Al-W ternary was confirmed to be metastable [20]. Large amounts of nickel need to be incorporated to improve its thermal stability [19]. We, thus, paid attention to other Cobase system, meanwhile noted that apart from the Co-W [21,22] and Co-Al [23,24] systems, the L12-ordered phases were also observed in other Co-base systems, such as Co-V [25–27], Co-Ti [28,29] and Co-TaCr [30,31] systems, respectively. Among these Co-base systems, only the γʹ phase in Co-Ti system is stable, and exists until its melting point (Tm) [32]. Furthermore, the γʹ-Co3Ti not only shows a low density, but also displays attractive ductility [33–36] which is one of the important factors for practical usages. Disappointedly, the large mismatch between γ and γʹ-Co3Ti and low Tm restrict the applications of the alloys reinforced by γʹ-Co3Ti [29,37]. In our previous work [38], V was found
Corresponding author. E-mail address: [email protected] (C.P. Wang).
http://dx.doi.org/10.1016/j.intermet.2017.09.015 Received 7 July 2017; Received in revised form 22 September 2017; Accepted 23 September 2017 0966-9795/ © 2017 Elsevier Ltd. All rights reserved.
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to increase the Tm of the alloy containing γʹ-Co3Ti, and V-modified CoTi-V ternary alloy also shows the moderate strength. Meanwhile, according to other researchers' work, V was also confirmed to raise the γʹCo3Ti's yield stress corresponding to TP [39] (Tp: the upper critical temperature relating to positive temperature dependence of strength) and to solid-solute into γʹ-Co3Ti phase with a large solubility (20 at. %) [40]. Therefore, the Co-Ti-V ternary system is suggested as the potential sub-system for developing novel Co-base superalloys. The primary objectives of the present work focusing on the Co-Ti-Vbase alloys are: (1) to study the microstructure evolution; (2) to investigate the alloying behavior of elements Al and Ni; (3) to examine the phase stabilities of γ and γʹ; (4) to exploit the high-temperature mechanical properties. The obtained results will provide important information for the development of the novel Co-base superalloys.
2. Experimental procedures High purity cobalt (99.9 wt %), vanadium (99.7 wt %), titanium (99.9 wt %), aluminum (99.9 wt %) and nickel (99.9 wt %) were utilized for preparing the cast ingots (Co-5Ti-15V (2X) (in at.%) where 2X is added as outer percentage of master alloy (Co-5Ti-15V), (X = Al and Ni)) which are respectively defined as MA, 2Al and 2Ni, as listed in Table 1 where the details of the heat treatments on these alloys are also provided. The alloys annealed at 800 °C for different times t (t = 1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h are designated as MAHt, 2AlHt and 2NiHt (t indicates the annealing time with the unit of hour), respectively. Some steps for preparing the samples investigated are summarized as follows: (1) the grinding machine was used to remove the oxide layer of cobalt, while the vanadium was cleaned by dilute HF solution; (2) the refractory element was melted with a certain amount of cobalt by arc melting under high purity argon atmosphere to form the alloys with lower Tm; (3) the vacuum induction furnace was employed to prepare the cast ingots under high purity argon atmosphere; (4) the ingots, weighted around 400 g, were cut into the shape with dimension of ϕ 6 mm × 9 mm by wire-cutting machine, and cylindrical ingots were slightly polished, sealed in quartz capsules filled with argon gas before heat treatment; (5) the samples were quickly quenched into ice water after the heat treatment. The transmission electron microscope (TEM) was employed to identify the crystal structure of the γʹ phase, and the foil specimens for TEM analyses were prepared through twin jet electropolishing in a solution of HClO4 (8 vol %) + CH3COOH (72 vol %) + CH3CH2OH (vol. 12%) + HOCH2CH2OH (vol. 8%). A field emission scanning electron microscope (FE-SEM) and electron probe microanalyzer (EPMA) were introduced to observe the microstructures of the Co-Ti-Vbase alloys, and the alloys for FE-SEM observations were mechanically polished, then etched in a solution of HCl (vol. 50%) + HNO3 (vol. 50%) for a few seconds. Since the distribution of the alloying elements strongly affects the high-temperature strength, creep resistance, and γʹ stability in superalloys, the partitioning behaviors of Al and Ni are thus desired to be evaluated. Normally used cold-rolling [11,12,15,19] was conducted on MAH94, 2AlH94 and 2NiH94 alloys, then, the cold-rolled alloys were annealed at 800 °C for 744, 406 and 704 h respectively to obtain the alloys with coarsened precipitates for composition analyses. The compositions of coarsened γ and γʹ phases were measured by a
Fig. 1. Phase equilibria in the Co-rich corner of the Co-Ti-V ternary system at 800 and 1100 °C [38].
field-emission electron probe microanalyzer (FE-EPMA). Besides, the phase transition temperatures were examined by differential scanning calorimetry (DSC) under argon atmosphere at a heating rate of 10 °C/ min. The compression tests with the strain rate of 1.0 × 10−4 s−1 at the temperature ranging from room temperature to 1000 °C were performed through a Thermecmastor Z machine. Hardnesses of the alloys were measured using a hardness testing machine (MVK-H1, Akashi) under the load with 0.5 kg. 3. Results and discussions 3.1. Composition design and microstructure control According to our previous reported results [38], 800 and 1100 °C were selected for annealing and homogenization respectively in the present work. Fig. 1 shows the γ-(αCo) + γʹ-Co3Ti phase equilibrium at 800 °C (solid lines) and 1100 °C (dotted lines). It is seen that the solubility of V in γʹ-Co3Ti and γ-(αCo), and Ti in γ-(αCo) increase with the temperature rising from 800 °C to 1100 °C. In order to obtain the Co-TiV ternary alloy with γ/γʹ two-phase microstructure, as well as to achieve a high volume fraction of the strengthening phase γʹ-Co3Ti, the alloy with a composition of Co-5Ti-15V (at. %) was, thus, designed using the information indicated in Fig. 1, and defined as master alloy (MA). Fig. 2(a) shows the morphology of single γ phase of the MA alloy homogenized at 1100 °C for 24 h, while Fig. 2(b), the bright-field TEM image (BFI), displays the uniform two-phase γ/γʹ microstructure of the MA alloy annealed at 800 °C for 184 h after high-temperature homogenization (MAH184). These results agree well with the phase relationship indicated in Fig. 1. Fig. 2(c)-(d) show a selected area diffraction pattern (SADP) and a dark-field TEM image (DFI) obtained from the MAH184 alloy, respectively. The homogeneously distributed irregular cuboidal precipitates are proved to be with L12-ordered structure using the superlattice reflections, 110 and 010, seen in
Table 1 The designation of the alloys investigated and the relating heating processes. Designation
Heat-treatment process
MA, 2Al, 2Ni MAH1,6,9 … 184 2AlH1,6,9 … 184 2NiH1,6,9 … 184
As-cast Obtained from the MA alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h Obtained from the 2Al alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h Obtained from the 2Ni alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h
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Fig. 2. Back-scattered SEM image obtained from the MA alloy homogenized at 1100 °C for 24 h (a); TEM images obtained from the alloy MAH184: (b) bright-field TEM image (BFI); (c) selected area diffraction pattern (SADP); (d) dark-field TEM image (DFI).
Fig. 5(a) indicate that: (1) aluminum promotes the growth of γʹ, while nickel hardly changes the growth rate of γʹ in early annealing stage (before ∼6 h) compared with that of the master alloy; (2) aluminum increases the Vγʹ significantly by 10%, while nickel decreases the Vγʹ slightly compared with that of MAH (79%) after long-time annealing (the Vγʹ after long-time heat treatment is represented by the mean value of Vγʹ obtained from alloys annealed for 94, 136, 160 and 184 h); (3) the Vγʹ is almost the same after 94 h, and this means that the annealing time with 94 h is adequate for developing the microstructure with high Vγʹ. It should be noted that the high Vγʹ is one of the essential aspects for improving the creep strength of the superalloys during its service period. Thus, the alloys for mechanical property tests in the present work were annealed at 800 °C for 94 h, and designated as MAH94, 2AlH94 and 2NiH94, as listed in Table 1. Besides, the Vickers-hardnesses of MAH94, 2AlH94 and 2NiH94 were measured to be about 416HV4.9, 297HV4.9 and 430HV4.9, respectively, and the Vickershardnesses of MAH94 and 2NiH94 alloys are even comparable with that of Co-9Al-9W (∼400) [10] and Waspaloy (∼390) [10]. The compositions of coarsened γ were measured as Co-1.9Ti-10.5V, Co-3.3Ti-11.9V-1.6Al and Co-2.3Ti-11.6V-1.9Ni (at. %), while the compositions of coarsened γʹ were measured as Co-5.5Ti-17.8V, Co5.2Ti-16.0V-2.1Al and Co-5.2Ti-17.3V-2.1Ni (at. %), as listed in Table 2. Fig. 5(b) shows the partitioning behavior of aluminum in CoTi-V-Al, and nickel in Co-Ti-V-Ni alloys. The partitioning parameter Kx is described as Kx = Cγ ′− x / Cγ − x where Cγʹ-x and Cγ-x indicate the concentration of element X in γʹ and γ phases, respectively. The element X tends to partition into γʹ when Kx > 1, while prefers to distribute into γ when Kx > 1. Both aluminum (KAl ∼ 1.3) and nickel (KNi ∼ 1.1) are confirmed as γʹ-former in the present work. The partitioning behavior
Fig. 2(c). Alloying elements play a significant role in improving the materials' properties. For example, Al is important for the oxidation resistance of the superalloy [2], and Ni is in favor of stabilizing the γʹ phase in Co-AlW alloy [19]. In the present work, the alloying elements, Al and Ni, were incorporated into the master alloy with 2% (at. %) to investigate their influences on Vγʹ, Tsolvus-γʹ, Tm and mechanical properties. Al- and Ni-modified alloys show the two-phase γ/γʹ microstructure after annealing at 800 °C for different times, as seen in Fig. 3(a)-(d) where the microstructures of 2AlH184 and 2NiH184 are displayed. It is seen that there is no other precipitate phase at the grain boundary, as indicated in Fig. 3(a) and (c), and both of them exhibit the ultrahigh fraction of irregular cuboidal γʹ particles precipitated in the matrix γ (Fig. 3(b) and (d)). The Co-Ti-V, Co-Ti-V-Al and Co-Ti-V-Ni alloys were respectively annealed at 800 °C for different times after homogenization at 1100 °C for 24 h to see their microstructure evolution during heating, then to find a proper annealing time for obtaining high Vγʹ. Fig. 4(a1)-(a4), (b1)-(b4) and (c1)-(c4) with same scale exhibit the microstructures of MAHt, 2AlHt and 2NiHt alloys annealed at 800 °C for 1, 6, 94 and 160 h, respectively. The area fraction of the γʹ was calculated on the obtained SEM images, and taken as the volume fraction. It is worth noting that the different observation direction on the sample will affect the value of the area fraction of the γʹ in the matrix. However, due to the high area fraction of the γʹ in the alloys investigated in the present work, the influence of observation direction on the area fraction of the γʹ is negligible. The γʹ particles in these alloys gradually grow in the first ∼94 h, and its morphology changes from vermicular into irregular cuboidal shape. The curves, Vγʹ versus annealing time, plotted in
Fig. 3. Back-scattered ((a) and (c)) and secondary electron ((b) and (d)) images respectively obtained from the alloys annealed at 800 °C: 2AlH184 ((a)–(b)) and 2NiH184 ((c)–(d)).
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Fig. 4. Secondary electron images obtained from: (a1-a4) MAHt; (b1-b4) 2AlHt; (c1-c4) 2NiHt.
of aluminum is consistent with the result indicated in Fig. 5(a) that aluminum increases the Vγʹ. Additionally, Fig. 5(a) also shows that the addition of nickel decreases the Vγʹ compared with the master alloy annealed, although nickel is proved to be a γʹ-former (Fig. 5(b)). Normally, alloying with γʹ-former element will increase Vγʹ in the alloy, however, it should be noted that the densities of the γ and γʹ phases also affect Vγʹ. The experimental results obtained by Zenk [41] shows that chromium is γ-former in the Co-11Ti-15Cr alloy but increase the Vγʹ in the alloy. Therefore, the present measured partitioning behavior of nickel has no conflict with the measured Vγʹ in Co-Ti-V-Ni alloy.
3.2. Phase transformations in Co-Ti-V-X (X: Al, Ni) alloys Since the Tm and Tsolvus-γʹ are two important properties for the practical usage of superalloys, the phase transformations in Co-Ti-V-X (X: Al, Ni) alloys were thus measured. The DSC heating curves respectively obtained from MAH94, 2AlH94 and 2NiH94 alloys are provided in Fig. 6(a)-(c), and relating phase transition points are summarized in Table 2 where the transition temperatures of potential novel cobalt-base superalloy Co-9Al-9W [13], traditional cobalt-base superalloy MAR-M302 [42] and commercial nickel-base superalloy IN-939 [43] are also provided. The peaks respectively located at 1091, 1112 and 1082 °C are related to the phase reaction γ+ γ′ → γ , while the peaks respectively located at 1292, 1273 and 1300 °C are referred to the phase reaction γ+ γ→ L in Fig. 6(a)-(c). Aluminum is found to increase the Tsolvus-γʹ by 21 °C, while nickel is discovered to increase the Tm by 8 °C compared with that of the master alloy. The data obtained from DSC measurements, together with the phase transition points of Co-9Al-9W [13], MAR-M-302 [42] and IN-939 [43] are plotted in Fig. 6(d). It is seen that: (1) the Tsolvus-γʹ of the MAH94, 2AlH94 and 2NiH94 are higher than that of Co-9Al-9W (1000 °C) by 91, 112, and 82 °C, respectively, while the Tsolvus-γʹ of Al-modified alloy is higher than the transition temperature of γʹ (1100 °C) in Ni-base superalloy IN-939 by 12 °C; (2) the Tm of the MAH94, 2AlH94 and 2NiH94 are higher than the incipient melting temperature of Co-base superalloy MAR-M302 (1238 °C) by 54, 35 and 62 °C, and higher than the incipient melting temperature of Ni-base superalloy IN-939 (1150 °C) by 142, 123 and 150 °C.
Fig. 5. Annealing time dependence of Vγ′ in MAHt, 2AlHt and 2NiHt at 800 °C (a), and the partitioning behaviors of Al and Ni are also plotted in (b).
Table 2 Compositions and phase transition points of the γ and γʹ phases in the present Co-Ti-Vbase alloys, and the reported phase transition points of the Co-9Al-9W [13], MAR-M302 [42] and IN-939 [43] are also provided. Alloys
Master alloy Al-modified alloy Ni-modified alloy Co-9Al-9W MAR-M302 IN-939
Composition (at. %) γ
γʹ
Co-1.9Ti-10.5V Co-3.3Ti-11.9V1.6Al Co-2.3Ti-11.6V1.9Ni – – –
Co-5.5Ti-17.8V Co-5.2Ti-16.0V2.1Al Co-5.2Ti-17.3V2.1Ni – – –
Tm (°C)
Tsolvus-γʹ (°C)
1292 1273
1091 1112
1300
1082
1458 1238 1150
1000 – 1100
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Fig. 6. DSC heating curves respectively obtained from: (a) MAH94; (b) 2AlH94; (c) 2NiH94, and the transition temperatures of the Co-9Al-9W [13], MAR-M302 [42] and IN-939 [43] are also plotted with the present DSC data, as displayed in (d).
dislocations, is widely accepted as one of the explanations for this abnormal phenomenon. In the third stage, the strength exhibits the negative temperature dependence above the Tp. It is accepted that this phenomenon is attributed to the thermally activated slip on the cube plane [51]. However, in the work performed by Suzuki et al. [13,16,52], the slip on octahedral plane has been observed above the Tp, and they claimed that this abnormal phenomenon could be attributed to the transformation from γʹ to γ before the slip on cube plane. Suzuki et al. [13,16] investigated the Co-9Al-10W-2Ta alloy slightly deformed at 1173 and 1243 K which are lower than the solvus temperature (1352 K) of the γʹ particle in this alloy. A large amount of 1/ 3 < 112 > partial dislocations can be seen in these alloys. Besides, the density of the superlattice intrinsic stacking fault (SISF) in alloy deformed at 1243 K seems lower than that in the alloy deformed at 1173 K. Meanwhile, their investigations [16] on the Co-9Al-11W alloy slightly deformed at 1173 and 1239 K shows that the major defects in these alloys are 1/2 < 110 > unpaired dislocations bypassing γʹ particle rather than SISF. A small amount of SISF can be seen in the alloy deformed at 1173 K. Suzuki et al. [16] also mentioned that the γʹ particles were supposed to be partially dissolved due to the high compression temperature (1239 K) which is close to the solvus temperature (1306 K) of the γʹ particle in this alloy. Therefore, it is reasonable to speculate that the rapid decrement of the strength above Tp in the present studied alloys is attributed to the activation of 1/ 3 < 112 > partial dislocations at lower temperature and the transformation in interaction mechanism between dislocation and γʹ particle from shearing to bypassing at higher temperature. The same interaction mechanism between dislocation and γʹ particle were also reported in Ni-base superalloy PWA1480 [53,54]. Besides, considering the higher Tsolvus-γʹ of the present studied alloys compared with that of Co-Al-Wbase alloys [13,16], the slip on cube plane above Tp is supposed to happen if the deformation temperature is high enough. The 0.2% flow stresses relating to Tp are 522, 355 and 544 MPa in MAH94, 2AlH94 and 2NiH94 alloys, respectively. The 0.2% flow stresses of the MAH94, 2AlH94 and 2NiH94 are found as 416, 332 and 364 MPa at 900 °C, while decrease to 197, 163 and 221 MPa at 1000 °C, respectively. The 0.2% flow stresses of the MAH94 and 2NiH94 are higher than that of Co3Ti [45], Co-12Ti [46] and Co-9Al-9W [13] in the temperatures investigated. The 0.2% flow stress of the 2AlH94 exceeds the strength of the Co-12Ti above 850 °C, while almost equals the strength of the Co-12Ti below 850 °C [46]. The 0.2% flow stresses of the MAH94 and 2NiH94 are higher than that of Co-11Ti-15Cr [41] below 750 °C, and the difference between them increases as the temperature gradually decreases to room temperature. Above 900 °C, the 0.2% flow stresses of the MAH94 and 2NiH94 are close to that of Co-11Ti-15Cr alloy [41]. The Tps (∼750 °C) of the alloys compressed in the present
3.3. Mechanical properties of the Co-Ti-V-X (X: Al, Ni) alloys The temperature dependences of the 0.2% flow stresses of the MAH94, 2AlH94 and 2NiH94 heat treated alloys are plotted in Fig. 7 where the high-temperature strengths of IN-939 (tensile tests) [44], MAR-M-302 (tensile tests) [44], Co-9Al-9W (compression tests) [13], Co3Ti (compression tests) [45], Co-12Ti (compression tests) [46] and Co-11Ti-15Cr (compression tests) [41] are also provided. For the alloys compressed, the compression curves can be divided into three stages. The first stage is the normal degeneration of the strength with increasing temperature below ∼600 °C. The second stage displays the abnormal phenomenon that the yield stress increases rather than decreases as the temperature rises from ∼600 to 750 °C. The same phenomenon is normally found in Ni3Al [45,47] which is a main strengthening phase in Ni-base superalloys, and believed to favor the high-temperature utilizations of the superalloy reinforced by Ni3Al. The Kear-Wilsdorf mechanism [48–50], demonstrating the locking of the
Fig. 7. The temperature dependences of the 0.2% flow stress for: (a) MAH94; (b) 2AlH94; (c) 2NiH94, the tensile yield stresses of the nickel-base superalloy IN-939 [44], cobaltbase superalloy MAR-M302 [44] and the 0.2% flow stresses of the Co-9Al-9W [13], Co3Ti [45], Co-12Ti [46] and Co-11Ti-15Cr [41] are also provided.
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work are higher than that of Co-9Al-9W [13] by ∼ 50 °C, while almost equal to that of Co3Ti [45]. The strengths of the MAH94 and 2NiH94 alloys exceed that of the MAR-M-302 [44] at 700–1000 °C. The strengths of MAH94 and 2NiH94 alloys are higher than that of MAR-M302 by about 70 and 92 MPa at 750 °C, respectively, while the differences increase to 120 and 144 MPa at 850 °C. The 0.2% flow stresses of MAH94 and 2NiH94 alloys are comparable with that of the IN-939 [44] at high-temperature range (850–1000 °C). The Vγʹ of the 2AlD94 alloy is the highest among the alloys examined, as indicated in Figs. 4 and 5(a). Generally, the high Vγʹ suggests the high strength of the alloys. However, the 0.2% flow stress (Fig. 7) and hardness (297HV4.9) of the alloy containing Al are the poorest. Suzuki et al. [16] speculated the decrement of superlattice intrinsic stacking fault energy (ΔESISF ) of γʹ by adding Ta into Co-Al-W alloy through the observation of a large amount of SISF within γʹ in the CoAl-W-Ta alloy deformed at the temperature above Tp. The presence of SISF was thought to sustain the strength of Co-Al-W-Ta alloy to higher temperature. However, according to the first-principles study on Co-AlW-Ta performed by Mottura et al. [55], Ta was calculated to increase rather than decrease the ΔESISF of γʹ, that means the maintenance of strength at higher temperature is not merely due to the presence of SISF within γʹ. Thus, Mottura et al. [55] indicated that the higher Vγʹ combined with higher shear stress for dislocation initially penetrate into γʹ caused by high ΔESISF are expected to be the main reasons for the strengthening of Co-Al-W-Ta alloy. Meanwhile, it should be noted that the high Vγʹ is the significant premise for the explanation proposed by Mottura et al. [55]. Thus, it is reasonable to speculate that the poor performance of Al-containing alloys in the present work is attributed to the decrement of ΔESISF caused by alloying with Al, the low ΔESISF will in favor of initial penetration of dislocation into γʹ.
Acknowledgements This work was supported by the Ministry of Science and Technology of China (Grant No. 2014DFA53040), the National Natural Science Foundation of China (Grant Nos. 51471138 and 51571168) and the National Key Research and Development Program (2017YFB0702901). References [1] M.J. Donachie, S.J. Donachie, Superalloys: A Technical Guide, ASM International, USA, 2002, pp. 2–3. [2] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, New York, 1987, pp. 97–188. [3] T. Nishizawa, K. Ishida, The Co-Ni (Cobalt-Nickel) system, J. Phase Equilibria 4 (4) (1983) 390–395. [4] L. Swartzendruber, V. Itkin, C. Alcock, The Fe-Ni (iron-nickel) system, J. phase Equil. 12 (3) (1991) 288–312. [5] W. Huang, Y.A. Chang, A thermodynamic analysis of the Ni-Al system, Intermetallics 6 (6) (1998) 487–498. [6] Y. Du, N. Clavaguera, Thermodynamic assessment of the Al-Ni system, J. Alloys Compd. 237 (1–2) (1996) 20–32. [7] I. Ansara, N. Dupin, H.L. Lukas, B. 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Inui, Effects of quaternary alloying elements on the γ′ solvus temperature of Co-Al-W based alloys with fcc/L12 two-phase microstructures, J. Alloys Compd. 508 (1) (2010) 71–78. [19] K. Shinagawa, T. Omori, J. Sato, K. Oikawa, I. Ohnuma, R. Kainuma, K. Ishida, Phase equilibria and microstructure on. GAMMA.'Phase in Co-Ni-Al-W system, Mater. Trans. 49 (6) (2008) 1474–1479. [20] S. Kobayashi, Y. Tsukamoto, T. Takasugi, H. Chinen, T. Omori, K. Ishida, S. Zaefferer, Determination of phase equilibria in the Co-rich Co-Al-W ternary system with a diffusion-couple technique, Intermetallics 17 (12) (2009) 1085–1089. [21] J. Dutkiewicz, G. Kostorz, Strengthening of cobalt-tungsten alloys upon discontinuous precipitation, Acta Metall. Mater. 38 (11) (1990) 2283–2286. [22] J. Dutkiewicz, G. Kostorz, Structure of martensite in Co-W alloys, Mater. Sci. Eng. A 132 (1991) 267–272. [23] A. Bradley, G. Seager, An X-ray investigation of cobalt-aluminium alloys, J. Inst. Met. 64 (1939) 81–88. [24] T. Omori, Y. Sutou, K. Oikawa, R. Kainuma, K. Ishida, Shape memory and magnetic properties of Co-Al ferromagnetic shape memory alloys, Mater. Sci. Eng. A 438 (2006) 1045–1049. [25] E.T. Peters, L.E. Tanner, A new high-temperature form of the intermetallic compound Co3V, AIME Met. Soc. Trans. 233 (12) (1965) 2126–2128. [26] Y. Aoki, K. Asami, M. Yamamoto, Transformation temperatures and magnetic properties of the ordered hexagonal VCo3 compound, Phys. Status Solidi (a) 23 (2) (1974) K167–K169. [27] Y. Aoki, J. Echigoya, Phase separation in an 18.2at. % V-Co alloy annealed at 1073 K, Scr. Metall. 19 (5) (1985) 639–642. [28] J.M. Blaise, P. Viatour, J.M. Drapier, On the stability and precipitation of the Co3Ti phase in Co-Ti alloys, Cobalt 49 (1970) 192–195. [29] P. Viatour, J.M. Drapier, D. Coutsouradis, Stability of the gamma prime Co3Ti compound in simple and complex Co alloys, Cobalt 3 (1973) 67–74. [30] J.M. Drapier, J.L. De Brouwer, D. 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4. Summary The microstructure, partitioning behavior, phase stability and mechanical property have been investigated on the Co-Ti-V-base alloys. Importantly, the ultrahigh Vγʹ, high Tsolvus-γʹ, mild Tm and the favorable high-temperature strength of the present designed Co-Ti-V-base alloys suggest them as the potential candidates for being utilized at elevated temperature. The main results are summarized as follows: (1) The Co-Ti-V-X (X: Al and Ni) alloys were investigated, and the twophase microstructure γ/γʹ with high Vγʹ were designed successfully in Co-Ti-V, Co-Ti-V-Al and Co-Ti-V-Ni systems. (2) The microstructure evolutions of Co-Ti-V, Co-Ti-V-Al and Co-Ti-VNi alloys at 800 °C were studied. Al is found to accelerate the growth of γʹ, and to increase the Vγʹ. The appropriate annealing time for reaching the high Vγʹ in these alloys was confirmed as 94 h. The Vγʹs of the present designed alloys are higher than 75%. (3) The compositions of the γʹ were measured to be about Co-5.5Ti17.8V, Co-5.2Ti-16.0V-2.1Al and Co-5.2Ti-17.3V-2.1Ni (at. %) in the coarsened microstructure, respectively. Al is confirmed as γʹformer while Ni equally distributes into γ and γʹ phases. The Tsolvusγʹs were determined as 1091, 1112 and 1082 °C, while the Tms are 1292, 1273 and 1300 °C in the MAH94, 2AlH94 and 2NiH94, respectively. (4) The mechanical properties of the Co-Ti-V, Co-Ti-V-Al and Co-Ti-VNi alloys were studied, and the flow stresses of these alloys exhibit an anomalous positive dependence on temperature ranging from 600 to 750 °C. The strengths of Co-Ti-V and Co-Ti-V-Ni alloys are higher than the strengths of Co3Ti and Co-9Al-9W at all temperatures investigated, and exceeds the strength of traditional Co-base superalloy MAR-M 302 at the temperature ranging from 700 to 1000 °C, and also comparable with the strength of commercial Nibase superalloy IN-939 between 850 and 1000 °C.
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[33] [34] [35]
[36] [37] [38]
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[44] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, New York, 1987, pp. 584–585. [45] D.M. Wee, O. Noguchi, Y. Oya, T. Suzuki, New Ll2 ordered alloys having the positive temperature dependence of strength, Trans. Jpn. Inst. Metals 21 (4) (1980) 237–247. [46] C.H. Zenk, S. Neumeier, H. Stone, M. Göken, Mechanical properties and lattice misfit of γ/γ′ strengthened Co-base superalloys in the Co-W-Al-Ti quaternary system, Intermetallics 55 (2014) 28–39. [47] R.W. Guard, J.H. Westbrook, Alloying behavior of NI3AL (GAMMA-PHASE), Trans. Am. Inst. Min. Metall. Eng. 215 (5) (1959) 807–814. [48] B.H. Kear, H.G.F. Wilsdorf, Dislocation configurations in plastically deformed polycrystalline Cu3Au alloys, Trans. Metall. Soc. AIME 224 (1962) 382–386. [49] S. Takeuchi, E. Kuramoto, Temperature and orientation dependence of the yield stress in Ni3Ga single crystals, Acta metall. 21 (4) (1973) 415–425. [50] V. Paidar, D.P. Pope, V. Vitek, A theory of the anomalous yield behavior in L12 ordered alloys, Acta Metall. 32 (3) (1984) 435–448. [51] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge university press, New York, 2006, pp. 33–120. [52] A. Suzuki, H. Inui, T.M. Pollock, L12-strengthened cobalt-base superalloys, Annu. Rev. Mater. Res. 45 (2015) 345–368. [53] W.W. Milligan, S.D. Antolovich, Yielding and deformation behavior of the single crystal superalloy PWA 1480, Metall. Trans. A 18 (1) (1987) 85–95. [54] W.W. Milligan, S.D. Antolovich, The mechanisms and temperature dependence of superlattice stacking fault formation in the single-crystal superalloy PWA 1480, Metall. Mater. Trans. A 22 (10) (1991) 2309–2318. [55] A. Mottura, A. Janotti, T.M. Pollock, A first-principles study of the effect of Ta on the superlattice intrinsic stacking fault energy of L12-Co3(Al, W), Intermetallics 28 (2012) 138–143.
J.E. Blendell, A.J. Shapiro, Determination of the CoTi congruent melting point and thermodynamic reassessment of the Co-Ti system, Metall. Mater. Trans. A 32 (9) (2001) 2175–2186. T. Takasugi, O. Izumi, High temperature strength and ductility of polycrystalline Co3Ti, Acta Metall. 33 (1) (1985) 39–48. T. Takasugi, O. Izumi, Recrystallization and grain growth of Co3Ti, Acta Metall. 33 (1) (1985) 49–58. T. Takasugi, O. Izumi, Electronic and structural studies of grain boundary strength and fracture in L12 ordered alloys-I. On binary A3B alloys, Acta Metall. 33 (7) (1985) 1247–1258. H.Y. Yasuda, S. Yamamura, Y. Umakoshi, Recrystallization texture in cold-rolled Co3Ti polycrystals, Scr. Mater. 44 (5) (2001) 765–769. C. Rogister, D. Coutsouradis, L. Habraken, Improvement of heat-resisting cobaltbase alloys by precipitation hardening, Cobalt 34 (1967) 3–9. J.J. Ruan, C.P. Wang, C.C. Zhao, S.Y. Yang, T. Yang, X.J. Liu, Experimental investigation of phase equilibria and microstructure in the Co-Ti-V ternary system, Intermetallics 49 (2014) 121–131. Y. Liu, T. Takasugi, O. Izumi, H. Suenaga, Mechanical properties of Co3Ti polycrystals alloyed with various additions, J. Mater. Sci. 24 (12) (1989) 4458–4466. Y. Liu, T. Takasugi, O. Izumi, Alloying behavior of Co3Ti, Metall. Trans. A 17 (8) (1986) 1433–1439. C.H. Zenk, I. Povstugar, R. Li, F. Rinaldi, S. Neumeier, D. Raabe, M. Göken, A novel type of Co-Ti-Cr-base γ/γ′ superalloys with low mass density, Acta Mater. 135 (2017) 244–251. J.C. Freche, Nickel-base alloys for aerospace applications, Aerosp. Struct. Mater. NASA SP-227 (1970) 71–89. M.R. Jahangiri, S.M.A. Boutorabi, H. Arabi, Study on incipient melting in cast Ni base IN939 superalloy during solution annealing and its effect on hot workability, Mater. Sci. Technol. 28 (12) (2012) 1402–1413.
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Novel Co-Ti-V-base superalloys reinforced by L12-ordered γ′ phase a
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J.J. Ruan , X.J. Liu , S.Y. Yang , W.W. Xu , T. Omori , T. Yang , B. Deng , H.X. Jiang , C.P. Wanga,∗, R. Kainumab, K. Ishidab a Department of Materials Science and Engineering, College of Materials and Fujian Provincial Key Laboratory of Materials Genome, Xiamen University, Xiamen, 361005, PR China b Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai, 980-8579, Japan
A R T I C L E I N F O
A B S T R A C T
Keywords: Cobalt-base superalloys Compression test L12 compound Phase diagram
The influences of alloying elements Al and Ni to a Co-Ti-V master alloy on microstructures, phase stabilities and high-temperature mechanical properties have been investigated. A two-phase microstructure γ/γʹ with ultrahigh γʹ volume fraction was designed in the Co-Ti-V (master alloy), Co-Ti-V-Al (Al-modified) and Co-Ti-V-Ni (Nimodified) alloys. Al shows a slight tendency to partition into γʹ, while Ni exhibits almost equal distribution between the γ and γʹ phases. Al is found to increase the volume fraction and solvus temperature of γʹ. The solvus temperatures of γʹ in the master alloy, Al-modified and Ni-modified alloys exceed that of Co-9Al-9W alloy (1000 °C) by 91, 112 and 82 °C, respectively. The flow stresses of the designed alloys exhibit an anomalous positive dependence on temperature rising from 600 to 750 °C. The strengths of the master alloy and Ni-modified alloys are not only higher than that of traditional Co-base superalloy MAR-M 302 at the temperature ranging from 700 to 1000 °C, but also comparable with that of the commercial Ni-base superalloy IN-939 between 850 and 1000 °C. In conclusion, the present Co-Ti-V-base alloys reinforced by γʹ are suggested as potential candidates for high-temperature utilizations.
1. Introduction The superalloys, which normally serve in severe industrial environments facing high temperature and pressure, can be generally divided into three classes named as nickel-, cobalt- and iron-base superalloys [1,2]. The melting point of nickel (∼1455 °C) is lower compared with that of cobalt (∼1495 °C) and iron (∼1538 °C) [3,4], and it partly suggests the higher possibility for cobalt- and iron-base superalloys to be utilized at evaluated temperature than Ni-base superalloys. However, the fact is that the Ni-base superalloys reinforced by coherent fcc-ordered Ni3Al (L12 structure) phase [5–8] are more widely utilized in nowadays industries than the other two types. It is believed that the lack of a coherent geometrically close-packed (GCP) phase in cobaltand iron-base superalloys is the main limitation for their applications [2]. It needs to point out that the cobalt- and iron-base superalloys are still playing a main role in some industrial fields including petrochemical industries, nuclear power plants and gas turbine industries etc. due to some reasons such as the less cost of iron-base superalloys and good resistance to high-temperature hot corrosion (HTHC) of cobalt-base superalloys [2,9]. In recent years, GCP phases were discovered in some Co-base ternary alloys: Co-Al-W [10], Co-Ge-W [11] and Co-Ga-W [12]. The
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most attractive aspects are high melting point (Tm) and good mechanical properties shown in Co-Al-W-base alloys [10]. Thus, many studies relating to the Co-Al-W-base alloys were also conducted, for instance, the influence of alloying elements, Ta, Re, Cr, Ti, Nb, Mo, W, Mn, Fe, Ni, V and B, on mechanical properties, volume fraction (Vγʹ) and solvus temperature (Tsolvus-γʹ) of γʹ of Co-Al-W-base alloy were investigated [10,13–19]. Their results suggest a great possibility of producing novel Co-base superalloys reinforced through coherent precipitate to satisfy the increasing severe demands in industries. However, the γʹ phase in Co-Al-W ternary was confirmed to be metastable [20]. Large amounts of nickel need to be incorporated to improve its thermal stability [19]. We, thus, paid attention to other Cobase system, meanwhile noted that apart from the Co-W [21,22] and Co-Al [23,24] systems, the L12-ordered phases were also observed in other Co-base systems, such as Co-V [25–27], Co-Ti [28,29] and Co-TaCr [30,31] systems, respectively. Among these Co-base systems, only the γʹ phase in Co-Ti system is stable, and exists until its melting point (Tm) [32]. Furthermore, the γʹ-Co3Ti not only shows a low density, but also displays attractive ductility [33–36] which is one of the important factors for practical usages. Disappointedly, the large mismatch between γ and γʹ-Co3Ti and low Tm restrict the applications of the alloys reinforced by γʹ-Co3Ti [29,37]. In our previous work [38], V was found
Corresponding author. E-mail address: [email protected] (C.P. Wang).
http://dx.doi.org/10.1016/j.intermet.2017.09.015 Received 7 July 2017; Received in revised form 22 September 2017; Accepted 23 September 2017 0966-9795/ © 2017 Elsevier Ltd. All rights reserved.
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to increase the Tm of the alloy containing γʹ-Co3Ti, and V-modified CoTi-V ternary alloy also shows the moderate strength. Meanwhile, according to other researchers' work, V was also confirmed to raise the γʹCo3Ti's yield stress corresponding to TP [39] (Tp: the upper critical temperature relating to positive temperature dependence of strength) and to solid-solute into γʹ-Co3Ti phase with a large solubility (20 at. %) [40]. Therefore, the Co-Ti-V ternary system is suggested as the potential sub-system for developing novel Co-base superalloys. The primary objectives of the present work focusing on the Co-Ti-Vbase alloys are: (1) to study the microstructure evolution; (2) to investigate the alloying behavior of elements Al and Ni; (3) to examine the phase stabilities of γ and γʹ; (4) to exploit the high-temperature mechanical properties. The obtained results will provide important information for the development of the novel Co-base superalloys.
2. Experimental procedures High purity cobalt (99.9 wt %), vanadium (99.7 wt %), titanium (99.9 wt %), aluminum (99.9 wt %) and nickel (99.9 wt %) were utilized for preparing the cast ingots (Co-5Ti-15V (2X) (in at.%) where 2X is added as outer percentage of master alloy (Co-5Ti-15V), (X = Al and Ni)) which are respectively defined as MA, 2Al and 2Ni, as listed in Table 1 where the details of the heat treatments on these alloys are also provided. The alloys annealed at 800 °C for different times t (t = 1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h are designated as MAHt, 2AlHt and 2NiHt (t indicates the annealing time with the unit of hour), respectively. Some steps for preparing the samples investigated are summarized as follows: (1) the grinding machine was used to remove the oxide layer of cobalt, while the vanadium was cleaned by dilute HF solution; (2) the refractory element was melted with a certain amount of cobalt by arc melting under high purity argon atmosphere to form the alloys with lower Tm; (3) the vacuum induction furnace was employed to prepare the cast ingots under high purity argon atmosphere; (4) the ingots, weighted around 400 g, were cut into the shape with dimension of ϕ 6 mm × 9 mm by wire-cutting machine, and cylindrical ingots were slightly polished, sealed in quartz capsules filled with argon gas before heat treatment; (5) the samples were quickly quenched into ice water after the heat treatment. The transmission electron microscope (TEM) was employed to identify the crystal structure of the γʹ phase, and the foil specimens for TEM analyses were prepared through twin jet electropolishing in a solution of HClO4 (8 vol %) + CH3COOH (72 vol %) + CH3CH2OH (vol. 12%) + HOCH2CH2OH (vol. 8%). A field emission scanning electron microscope (FE-SEM) and electron probe microanalyzer (EPMA) were introduced to observe the microstructures of the Co-Ti-Vbase alloys, and the alloys for FE-SEM observations were mechanically polished, then etched in a solution of HCl (vol. 50%) + HNO3 (vol. 50%) for a few seconds. Since the distribution of the alloying elements strongly affects the high-temperature strength, creep resistance, and γʹ stability in superalloys, the partitioning behaviors of Al and Ni are thus desired to be evaluated. Normally used cold-rolling [11,12,15,19] was conducted on MAH94, 2AlH94 and 2NiH94 alloys, then, the cold-rolled alloys were annealed at 800 °C for 744, 406 and 704 h respectively to obtain the alloys with coarsened precipitates for composition analyses. The compositions of coarsened γ and γʹ phases were measured by a
Fig. 1. Phase equilibria in the Co-rich corner of the Co-Ti-V ternary system at 800 and 1100 °C [38].
field-emission electron probe microanalyzer (FE-EPMA). Besides, the phase transition temperatures were examined by differential scanning calorimetry (DSC) under argon atmosphere at a heating rate of 10 °C/ min. The compression tests with the strain rate of 1.0 × 10−4 s−1 at the temperature ranging from room temperature to 1000 °C were performed through a Thermecmastor Z machine. Hardnesses of the alloys were measured using a hardness testing machine (MVK-H1, Akashi) under the load with 0.5 kg. 3. Results and discussions 3.1. Composition design and microstructure control According to our previous reported results [38], 800 and 1100 °C were selected for annealing and homogenization respectively in the present work. Fig. 1 shows the γ-(αCo) + γʹ-Co3Ti phase equilibrium at 800 °C (solid lines) and 1100 °C (dotted lines). It is seen that the solubility of V in γʹ-Co3Ti and γ-(αCo), and Ti in γ-(αCo) increase with the temperature rising from 800 °C to 1100 °C. In order to obtain the Co-TiV ternary alloy with γ/γʹ two-phase microstructure, as well as to achieve a high volume fraction of the strengthening phase γʹ-Co3Ti, the alloy with a composition of Co-5Ti-15V (at. %) was, thus, designed using the information indicated in Fig. 1, and defined as master alloy (MA). Fig. 2(a) shows the morphology of single γ phase of the MA alloy homogenized at 1100 °C for 24 h, while Fig. 2(b), the bright-field TEM image (BFI), displays the uniform two-phase γ/γʹ microstructure of the MA alloy annealed at 800 °C for 184 h after high-temperature homogenization (MAH184). These results agree well with the phase relationship indicated in Fig. 1. Fig. 2(c)-(d) show a selected area diffraction pattern (SADP) and a dark-field TEM image (DFI) obtained from the MAH184 alloy, respectively. The homogeneously distributed irregular cuboidal precipitates are proved to be with L12-ordered structure using the superlattice reflections, 110 and 010, seen in
Table 1 The designation of the alloys investigated and the relating heating processes. Designation
Heat-treatment process
MA, 2Al, 2Ni MAH1,6,9 … 184 2AlH1,6,9 … 184 2NiH1,6,9 … 184
As-cast Obtained from the MA alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h Obtained from the 2Al alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h Obtained from the 2Ni alloy annealed at 800 °C for different times (1, 6, 9, 24, 47, 69, 94, 136, 160, 184 h) after homogenization at 1100 °C for 24 h
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Fig. 2. Back-scattered SEM image obtained from the MA alloy homogenized at 1100 °C for 24 h (a); TEM images obtained from the alloy MAH184: (b) bright-field TEM image (BFI); (c) selected area diffraction pattern (SADP); (d) dark-field TEM image (DFI).
Fig. 5(a) indicate that: (1) aluminum promotes the growth of γʹ, while nickel hardly changes the growth rate of γʹ in early annealing stage (before ∼6 h) compared with that of the master alloy; (2) aluminum increases the Vγʹ significantly by 10%, while nickel decreases the Vγʹ slightly compared with that of MAH (79%) after long-time annealing (the Vγʹ after long-time heat treatment is represented by the mean value of Vγʹ obtained from alloys annealed for 94, 136, 160 and 184 h); (3) the Vγʹ is almost the same after 94 h, and this means that the annealing time with 94 h is adequate for developing the microstructure with high Vγʹ. It should be noted that the high Vγʹ is one of the essential aspects for improving the creep strength of the superalloys during its service period. Thus, the alloys for mechanical property tests in the present work were annealed at 800 °C for 94 h, and designated as MAH94, 2AlH94 and 2NiH94, as listed in Table 1. Besides, the Vickers-hardnesses of MAH94, 2AlH94 and 2NiH94 were measured to be about 416HV4.9, 297HV4.9 and 430HV4.9, respectively, and the Vickershardnesses of MAH94 and 2NiH94 alloys are even comparable with that of Co-9Al-9W (∼400) [10] and Waspaloy (∼390) [10]. The compositions of coarsened γ were measured as Co-1.9Ti-10.5V, Co-3.3Ti-11.9V-1.6Al and Co-2.3Ti-11.6V-1.9Ni (at. %), while the compositions of coarsened γʹ were measured as Co-5.5Ti-17.8V, Co5.2Ti-16.0V-2.1Al and Co-5.2Ti-17.3V-2.1Ni (at. %), as listed in Table 2. Fig. 5(b) shows the partitioning behavior of aluminum in CoTi-V-Al, and nickel in Co-Ti-V-Ni alloys. The partitioning parameter Kx is described as Kx = Cγ ′− x / Cγ − x where Cγʹ-x and Cγ-x indicate the concentration of element X in γʹ and γ phases, respectively. The element X tends to partition into γʹ when Kx > 1, while prefers to distribute into γ when Kx > 1. Both aluminum (KAl ∼ 1.3) and nickel (KNi ∼ 1.1) are confirmed as γʹ-former in the present work. The partitioning behavior
Fig. 2(c). Alloying elements play a significant role in improving the materials' properties. For example, Al is important for the oxidation resistance of the superalloy [2], and Ni is in favor of stabilizing the γʹ phase in Co-AlW alloy [19]. In the present work, the alloying elements, Al and Ni, were incorporated into the master alloy with 2% (at. %) to investigate their influences on Vγʹ, Tsolvus-γʹ, Tm and mechanical properties. Al- and Ni-modified alloys show the two-phase γ/γʹ microstructure after annealing at 800 °C for different times, as seen in Fig. 3(a)-(d) where the microstructures of 2AlH184 and 2NiH184 are displayed. It is seen that there is no other precipitate phase at the grain boundary, as indicated in Fig. 3(a) and (c), and both of them exhibit the ultrahigh fraction of irregular cuboidal γʹ particles precipitated in the matrix γ (Fig. 3(b) and (d)). The Co-Ti-V, Co-Ti-V-Al and Co-Ti-V-Ni alloys were respectively annealed at 800 °C for different times after homogenization at 1100 °C for 24 h to see their microstructure evolution during heating, then to find a proper annealing time for obtaining high Vγʹ. Fig. 4(a1)-(a4), (b1)-(b4) and (c1)-(c4) with same scale exhibit the microstructures of MAHt, 2AlHt and 2NiHt alloys annealed at 800 °C for 1, 6, 94 and 160 h, respectively. The area fraction of the γʹ was calculated on the obtained SEM images, and taken as the volume fraction. It is worth noting that the different observation direction on the sample will affect the value of the area fraction of the γʹ in the matrix. However, due to the high area fraction of the γʹ in the alloys investigated in the present work, the influence of observation direction on the area fraction of the γʹ is negligible. The γʹ particles in these alloys gradually grow in the first ∼94 h, and its morphology changes from vermicular into irregular cuboidal shape. The curves, Vγʹ versus annealing time, plotted in
Fig. 3. Back-scattered ((a) and (c)) and secondary electron ((b) and (d)) images respectively obtained from the alloys annealed at 800 °C: 2AlH184 ((a)–(b)) and 2NiH184 ((c)–(d)).
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Fig. 4. Secondary electron images obtained from: (a1-a4) MAHt; (b1-b4) 2AlHt; (c1-c4) 2NiHt.
of aluminum is consistent with the result indicated in Fig. 5(a) that aluminum increases the Vγʹ. Additionally, Fig. 5(a) also shows that the addition of nickel decreases the Vγʹ compared with the master alloy annealed, although nickel is proved to be a γʹ-former (Fig. 5(b)). Normally, alloying with γʹ-former element will increase Vγʹ in the alloy, however, it should be noted that the densities of the γ and γʹ phases also affect Vγʹ. The experimental results obtained by Zenk [41] shows that chromium is γ-former in the Co-11Ti-15Cr alloy but increase the Vγʹ in the alloy. Therefore, the present measured partitioning behavior of nickel has no conflict with the measured Vγʹ in Co-Ti-V-Ni alloy.
3.2. Phase transformations in Co-Ti-V-X (X: Al, Ni) alloys Since the Tm and Tsolvus-γʹ are two important properties for the practical usage of superalloys, the phase transformations in Co-Ti-V-X (X: Al, Ni) alloys were thus measured. The DSC heating curves respectively obtained from MAH94, 2AlH94 and 2NiH94 alloys are provided in Fig. 6(a)-(c), and relating phase transition points are summarized in Table 2 where the transition temperatures of potential novel cobalt-base superalloy Co-9Al-9W [13], traditional cobalt-base superalloy MAR-M302 [42] and commercial nickel-base superalloy IN-939 [43] are also provided. The peaks respectively located at 1091, 1112 and 1082 °C are related to the phase reaction γ+ γ′ → γ , while the peaks respectively located at 1292, 1273 and 1300 °C are referred to the phase reaction γ+ γ→ L in Fig. 6(a)-(c). Aluminum is found to increase the Tsolvus-γʹ by 21 °C, while nickel is discovered to increase the Tm by 8 °C compared with that of the master alloy. The data obtained from DSC measurements, together with the phase transition points of Co-9Al-9W [13], MAR-M-302 [42] and IN-939 [43] are plotted in Fig. 6(d). It is seen that: (1) the Tsolvus-γʹ of the MAH94, 2AlH94 and 2NiH94 are higher than that of Co-9Al-9W (1000 °C) by 91, 112, and 82 °C, respectively, while the Tsolvus-γʹ of Al-modified alloy is higher than the transition temperature of γʹ (1100 °C) in Ni-base superalloy IN-939 by 12 °C; (2) the Tm of the MAH94, 2AlH94 and 2NiH94 are higher than the incipient melting temperature of Co-base superalloy MAR-M302 (1238 °C) by 54, 35 and 62 °C, and higher than the incipient melting temperature of Ni-base superalloy IN-939 (1150 °C) by 142, 123 and 150 °C.
Fig. 5. Annealing time dependence of Vγ′ in MAHt, 2AlHt and 2NiHt at 800 °C (a), and the partitioning behaviors of Al and Ni are also plotted in (b).
Table 2 Compositions and phase transition points of the γ and γʹ phases in the present Co-Ti-Vbase alloys, and the reported phase transition points of the Co-9Al-9W [13], MAR-M302 [42] and IN-939 [43] are also provided. Alloys
Master alloy Al-modified alloy Ni-modified alloy Co-9Al-9W MAR-M302 IN-939
Composition (at. %) γ
γʹ
Co-1.9Ti-10.5V Co-3.3Ti-11.9V1.6Al Co-2.3Ti-11.6V1.9Ni – – –
Co-5.5Ti-17.8V Co-5.2Ti-16.0V2.1Al Co-5.2Ti-17.3V2.1Ni – – –
Tm (°C)
Tsolvus-γʹ (°C)
1292 1273
1091 1112
1300
1082
1458 1238 1150
1000 – 1100
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Fig. 6. DSC heating curves respectively obtained from: (a) MAH94; (b) 2AlH94; (c) 2NiH94, and the transition temperatures of the Co-9Al-9W [13], MAR-M302 [42] and IN-939 [43] are also plotted with the present DSC data, as displayed in (d).
dislocations, is widely accepted as one of the explanations for this abnormal phenomenon. In the third stage, the strength exhibits the negative temperature dependence above the Tp. It is accepted that this phenomenon is attributed to the thermally activated slip on the cube plane [51]. However, in the work performed by Suzuki et al. [13,16,52], the slip on octahedral plane has been observed above the Tp, and they claimed that this abnormal phenomenon could be attributed to the transformation from γʹ to γ before the slip on cube plane. Suzuki et al. [13,16] investigated the Co-9Al-10W-2Ta alloy slightly deformed at 1173 and 1243 K which are lower than the solvus temperature (1352 K) of the γʹ particle in this alloy. A large amount of 1/ 3 < 112 > partial dislocations can be seen in these alloys. Besides, the density of the superlattice intrinsic stacking fault (SISF) in alloy deformed at 1243 K seems lower than that in the alloy deformed at 1173 K. Meanwhile, their investigations [16] on the Co-9Al-11W alloy slightly deformed at 1173 and 1239 K shows that the major defects in these alloys are 1/2 < 110 > unpaired dislocations bypassing γʹ particle rather than SISF. A small amount of SISF can be seen in the alloy deformed at 1173 K. Suzuki et al. [16] also mentioned that the γʹ particles were supposed to be partially dissolved due to the high compression temperature (1239 K) which is close to the solvus temperature (1306 K) of the γʹ particle in this alloy. Therefore, it is reasonable to speculate that the rapid decrement of the strength above Tp in the present studied alloys is attributed to the activation of 1/ 3 < 112 > partial dislocations at lower temperature and the transformation in interaction mechanism between dislocation and γʹ particle from shearing to bypassing at higher temperature. The same interaction mechanism between dislocation and γʹ particle were also reported in Ni-base superalloy PWA1480 [53,54]. Besides, considering the higher Tsolvus-γʹ of the present studied alloys compared with that of Co-Al-Wbase alloys [13,16], the slip on cube plane above Tp is supposed to happen if the deformation temperature is high enough. The 0.2% flow stresses relating to Tp are 522, 355 and 544 MPa in MAH94, 2AlH94 and 2NiH94 alloys, respectively. The 0.2% flow stresses of the MAH94, 2AlH94 and 2NiH94 are found as 416, 332 and 364 MPa at 900 °C, while decrease to 197, 163 and 221 MPa at 1000 °C, respectively. The 0.2% flow stresses of the MAH94 and 2NiH94 are higher than that of Co3Ti [45], Co-12Ti [46] and Co-9Al-9W [13] in the temperatures investigated. The 0.2% flow stress of the 2AlH94 exceeds the strength of the Co-12Ti above 850 °C, while almost equals the strength of the Co-12Ti below 850 °C [46]. The 0.2% flow stresses of the MAH94 and 2NiH94 are higher than that of Co-11Ti-15Cr [41] below 750 °C, and the difference between them increases as the temperature gradually decreases to room temperature. Above 900 °C, the 0.2% flow stresses of the MAH94 and 2NiH94 are close to that of Co-11Ti-15Cr alloy [41]. The Tps (∼750 °C) of the alloys compressed in the present
3.3. Mechanical properties of the Co-Ti-V-X (X: Al, Ni) alloys The temperature dependences of the 0.2% flow stresses of the MAH94, 2AlH94 and 2NiH94 heat treated alloys are plotted in Fig. 7 where the high-temperature strengths of IN-939 (tensile tests) [44], MAR-M-302 (tensile tests) [44], Co-9Al-9W (compression tests) [13], Co3Ti (compression tests) [45], Co-12Ti (compression tests) [46] and Co-11Ti-15Cr (compression tests) [41] are also provided. For the alloys compressed, the compression curves can be divided into three stages. The first stage is the normal degeneration of the strength with increasing temperature below ∼600 °C. The second stage displays the abnormal phenomenon that the yield stress increases rather than decreases as the temperature rises from ∼600 to 750 °C. The same phenomenon is normally found in Ni3Al [45,47] which is a main strengthening phase in Ni-base superalloys, and believed to favor the high-temperature utilizations of the superalloy reinforced by Ni3Al. The Kear-Wilsdorf mechanism [48–50], demonstrating the locking of the
Fig. 7. The temperature dependences of the 0.2% flow stress for: (a) MAH94; (b) 2AlH94; (c) 2NiH94, the tensile yield stresses of the nickel-base superalloy IN-939 [44], cobaltbase superalloy MAR-M302 [44] and the 0.2% flow stresses of the Co-9Al-9W [13], Co3Ti [45], Co-12Ti [46] and Co-11Ti-15Cr [41] are also provided.
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work are higher than that of Co-9Al-9W [13] by ∼ 50 °C, while almost equal to that of Co3Ti [45]. The strengths of the MAH94 and 2NiH94 alloys exceed that of the MAR-M-302 [44] at 700–1000 °C. The strengths of MAH94 and 2NiH94 alloys are higher than that of MAR-M302 by about 70 and 92 MPa at 750 °C, respectively, while the differences increase to 120 and 144 MPa at 850 °C. The 0.2% flow stresses of MAH94 and 2NiH94 alloys are comparable with that of the IN-939 [44] at high-temperature range (850–1000 °C). The Vγʹ of the 2AlD94 alloy is the highest among the alloys examined, as indicated in Figs. 4 and 5(a). Generally, the high Vγʹ suggests the high strength of the alloys. However, the 0.2% flow stress (Fig. 7) and hardness (297HV4.9) of the alloy containing Al are the poorest. Suzuki et al. [16] speculated the decrement of superlattice intrinsic stacking fault energy (ΔESISF ) of γʹ by adding Ta into Co-Al-W alloy through the observation of a large amount of SISF within γʹ in the CoAl-W-Ta alloy deformed at the temperature above Tp. The presence of SISF was thought to sustain the strength of Co-Al-W-Ta alloy to higher temperature. However, according to the first-principles study on Co-AlW-Ta performed by Mottura et al. [55], Ta was calculated to increase rather than decrease the ΔESISF of γʹ, that means the maintenance of strength at higher temperature is not merely due to the presence of SISF within γʹ. Thus, Mottura et al. [55] indicated that the higher Vγʹ combined with higher shear stress for dislocation initially penetrate into γʹ caused by high ΔESISF are expected to be the main reasons for the strengthening of Co-Al-W-Ta alloy. Meanwhile, it should be noted that the high Vγʹ is the significant premise for the explanation proposed by Mottura et al. [55]. Thus, it is reasonable to speculate that the poor performance of Al-containing alloys in the present work is attributed to the decrement of ΔESISF caused by alloying with Al, the low ΔESISF will in favor of initial penetration of dislocation into γʹ.
Acknowledgements This work was supported by the Ministry of Science and Technology of China (Grant No. 2014DFA53040), the National Natural Science Foundation of China (Grant Nos. 51471138 and 51571168) and the National Key Research and Development Program (2017YFB0702901). References [1] M.J. Donachie, S.J. Donachie, Superalloys: A Technical Guide, ASM International, USA, 2002, pp. 2–3. [2] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, New York, 1987, pp. 97–188. [3] T. Nishizawa, K. Ishida, The Co-Ni (Cobalt-Nickel) system, J. Phase Equilibria 4 (4) (1983) 390–395. [4] L. Swartzendruber, V. Itkin, C. Alcock, The Fe-Ni (iron-nickel) system, J. phase Equil. 12 (3) (1991) 288–312. [5] W. Huang, Y.A. Chang, A thermodynamic analysis of the Ni-Al system, Intermetallics 6 (6) (1998) 487–498. [6] Y. Du, N. Clavaguera, Thermodynamic assessment of the Al-Ni system, J. Alloys Compd. 237 (1–2) (1996) 20–32. [7] I. Ansara, N. Dupin, H.L. Lukas, B. 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4. Summary The microstructure, partitioning behavior, phase stability and mechanical property have been investigated on the Co-Ti-V-base alloys. Importantly, the ultrahigh Vγʹ, high Tsolvus-γʹ, mild Tm and the favorable high-temperature strength of the present designed Co-Ti-V-base alloys suggest them as the potential candidates for being utilized at elevated temperature. The main results are summarized as follows: (1) The Co-Ti-V-X (X: Al and Ni) alloys were investigated, and the twophase microstructure γ/γʹ with high Vγʹ were designed successfully in Co-Ti-V, Co-Ti-V-Al and Co-Ti-V-Ni systems. (2) The microstructure evolutions of Co-Ti-V, Co-Ti-V-Al and Co-Ti-VNi alloys at 800 °C were studied. Al is found to accelerate the growth of γʹ, and to increase the Vγʹ. The appropriate annealing time for reaching the high Vγʹ in these alloys was confirmed as 94 h. The Vγʹs of the present designed alloys are higher than 75%. (3) The compositions of the γʹ were measured to be about Co-5.5Ti17.8V, Co-5.2Ti-16.0V-2.1Al and Co-5.2Ti-17.3V-2.1Ni (at. %) in the coarsened microstructure, respectively. Al is confirmed as γʹformer while Ni equally distributes into γ and γʹ phases. The Tsolvusγʹs were determined as 1091, 1112 and 1082 °C, while the Tms are 1292, 1273 and 1300 °C in the MAH94, 2AlH94 and 2NiH94, respectively. (4) The mechanical properties of the Co-Ti-V, Co-Ti-V-Al and Co-Ti-VNi alloys were studied, and the flow stresses of these alloys exhibit an anomalous positive dependence on temperature ranging from 600 to 750 °C. The strengths of Co-Ti-V and Co-Ti-V-Ni alloys are higher than the strengths of Co3Ti and Co-9Al-9W at all temperatures investigated, and exceeds the strength of traditional Co-base superalloy MAR-M 302 at the temperature ranging from 700 to 1000 °C, and also comparable with the strength of commercial Nibase superalloy IN-939 between 850 and 1000 °C.
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