Introduction

The Ni–22Cr–12Co–9Mo alloy (also known as INCOLOY 617®) is used throughout the world for high-temperature applications. It is also one of the materials proposed for some of the most ambitious EU and US projects for future generations of power plants where steam inlet temperatures might reach and even exceed 700 °C [1, 2].

The processing behavior of this type of alloy (i.e., castability, forgeability, weldability, etc.) is of primary importance for manufacturing the large-size components typical of the power plant industry. In general, severe shortcomings concerning rotor applications in high-temperature steam turbines have been noticed in all three types of superalloy, according to the accepted classification (γ′-, solid solution, γ″-strengthening classes) [3, 4]. In fact, the production of large components is hampered by the relatively high-flow strength at typical forging temperatures. On the other hand, high-standard mechanical properties (for example creep resistance) are another fundamental requirement for these alloys, for which long-term high-temperature microstructural stability is also needed.

The Ni–22Cr–12Co–9Mo alloy belongs to the group of alloys usually referred to as solid-solution strengthened (but also as carbide-hardened) materials [5, 6]. It contains molybdenum or tungsten as well as some carbon for additional carbide strengthening and is generally used in the solution-annealed condition. The mechanical properties of this alloy, like those of other superalloys, are strongly influenced by microstructural features at the normal operating temperature and may possibly be modified during service at high temperatures. Thus, a clear understanding of the role played by secondary phases is essential for a correct explanation of the long-term properties of the material [7–11].

The Ni–22Cr–12Co–9Mo alloy has been shown to exhibit a strengthening response during extended exposure to intermediate temperatures (i.e., 700–800 °C) [12]. Following heat treatment, the microstructure of the alloy is characterized by the presence of a series of precipitates both at grain boundaries and inside the grains. The effect of high-temperature exposure has been reported as enhancing the discontinuous precipitation of carbides, nitrides and carbo-nitrides, such as M23C6, TiN, and TiCN, at grain boundaries [3, 13]. Some particles also precipitate along and inside twins [3, 13–15]. By increasing the temperature from 700 to 900 °C, the amount of precipitates, both inside the grains and at grain boundaries, has been found to increase [13–15]. In fact, at ∼700 °C both TiN, TiCN, and M23C6 tend to remain small and stable because of the rather slow diffusion rate of Ti and the other heavy elements. Exposure to much higher temperatures (1,000–1,100 °C) results in a lower amount of relatively coarse precipitates [13, 15]. Among the particles observed in the alloy, CrMo(C,N) carbo-nitrides have been revealed by X-ray diffraction studies but not by TEM inspections, indicating that they are rather rare particles within the microstructure [13]. Triangular-shaped particles identified as MC carbides have also been observed while other small particles of polyhedral shape, have been identified as TiN [16, 17]. M23C6 carbides, appearing as polyhedral particles [16], characterize Ni–22Cr–12Co–9Mo alloy microstructure within the temperature range 700–1,100°C. Finally, exposure at ∼700 °C (especially during creep) has also been shown to form a network of acicular δ-phase (orthorhombic intermetallic Ni3Nb phase), bridging grain boundaries and extending across grains [18]. The latter phase has been reported as forming directly from the supersaturated solid solution during aging at relatively high temperatures. The precipitation of the δ-Ni3Mo phase initiates in the Nb and Mo enriched regions of grain and twin boundaries and matrix carbides [19–23]. Precipitation of the δ-phase in Inconel 625 has been reported as very rapid at temperatures above 800 °C [22, 23], its solvus temperature being ∼1,010 °C with a sub-solvus temperature of ∼950 °C [24]. The δ-phase is incoherent and it weakly contributes to alloy strengthening, although it is desirable for grain size control during the processing and heat treatment of the wrought alloy [25–28]. The presence of the δ-phase at grain boundaries has been reported as beneficial for enhanced ductility and toughness [21–23, 29, 30]. In fact, some authors [19, 20, 31] have shown that the presence of the δ phase among other particles (both carbides and nitrides) within the microstructure does not significantly impair the toughness of the alloy, and, in some cases, may provide a barrier to grain boundary sliding during creep by pinning the boundaries. On the other hand, the laths of δ extending across grains appear to grow to the detriment of the strengthening γ′ precipitates.

The γ′ phase is a typical strengthening phase in many Ni-base superalloys with high Ti/Al ratios. It is usually in the form Ni3Ti or Ni3(Al,Ti) [26, 27]. This phase has also been found in small amounts in the Ni–22Cr–12Co–9Mo alloy exposed to a temperature range of 650–760 °C [3, 13]. However, no trace of it has been found after prolonged exposure to 816 °C [13] or after 2 h exposure to 800 °C. Therefore, the solvus temperature of the γ′ phase could reasonably lie in the temperature range of between 760 and 800 °C.

This paper focuses on microstructural changes, and particularly on precipitation phenomena, occurring in Ni–22Cr–12Co–9Mo alloy during exposure to temperatures of 700 and 800 °C in loaded and unloaded aged specimens.

Material and experimental procedures

The investigated material was an NiCr22O12Mo alloy (UNS N006617, Wr.N. 2.4663, that can also be referred to by its trade name, INCOLOY 617®), with the chemical composition shown in Table 1. The material was supplied by Krupp VDM as 12 mm thick plate, which was hot worked, solution treated at 1,180°C for 32 min and water quenched. The material was creep tested by a consortium of Italian Laboratories at temperatures of 700 and 800 °C [12]. These temperatures were chosen as they are, respectively, the hardening and softening reference temperatures for the investigated alloy. The results of the creep tests are presented in [12]. Details of temperatures and exposure times are reported in Table 2 of this paper together with the specimen reference codes.

Table 1 Chemical composition of the investigated 617 alloy (mass %)
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Table 2 Reference codes and general information about creep tested and microstructurally characterized materials: creep test temperature (T) and duration (t), specimen elongation (EL) and reduction in area (RA); Vickers (HV) and Rockwell C (HRC) hardness measured on the load-free gripping ends of the creep specimens
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Metallographic inspections were carried out on the load-free and load-bearing regions of selected specimens (gripping ends and gauge length, respectively) using optical microscopy (OM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM) techniques in order to characterize the presence of the secondary hardening phases. Metallographic samples for OM and SEM were longitudinally and transversally sectioned from crept specimens using a diamond saw cutting device, mounted in phenolic resin and polished. Depending on the microstructural condition, specimens were etched using either a solution of 15 mL HCl, 10 mL CH3COOH, 10 mL HNO3 (Aqua regia) or 30 mL HCl and 10 mL HNO3. Energy dispersive spectrum (EDS) microanalyses were performed in an SEM by means of an Oxford Inca Energy® 200 system. Films for TEM observations were mechanically ground to 200 μm and chemically etched using a solution of 10 mL HNO3 and 90 mL methyl alcohol. Discs were punched and finally thinned by twin-jet electropolishing (Tenupol®-5) using a solution of 10 mL HCl4 and 90 mL methyl alcohol at 22 V and T = −50 °C. A Philips® CM200 microscope, equipped with a double tilt specimen holder, was used. The vast majority of the phase identifications were performed using the convergent beam (CB) selected area electron diffraction pattern (SAEDP).

Vickers micro-hardness (981 N load on the indenter) was measured on the load-free regions of crept specimens (specimens 7C and 8C), and on the as-received material (specimen B). In addition to specimens used for metallographic investigations, macro-hardness tests were performed on other selected specimens exposed to the two investigated temperatures. In order to compare results with literature data [11], Rockwell C hardness (HRC) tests were also performed on the same specimens.

Results

As-received material

The typical metallographic features of the investigated material in the as-received condition (specimen B) are shown in Fig. 1. The presence of a banded structure is clearly visible (1(a)). The bands are characterized by fine austenitic grains and by a significant presence of secondary phases.

Fig. 1
figure 1

(a) Typical metallographic features (OM) of the as-received material (longitudinal section); (b) low-magnification TEM image showing a clear grained austenitic structure with straight and parallel twins. A few intragranular carbides are also visible

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The secondary phases observed with OM and SEM images were of two types. A first type of particle, which was relatively rare and homogeneously distributed, was light orange in OM micrographs. These particles, also visible on un-etched samples, had diameters in the range of 4–10 μm. Microanalyses showed that they were rich in Ti, C and N, and are therefore probably titanium carbo-nitrides Ti(C,N). The other set of particles lay along the banded structure. These particles were blocky-shaped, situated at inter- and intra-granular sites, and with a mean size of 0.2–1 μm. EDS analyses on the particles showed that they contained Cr and Mo. No trace of other hardening phases was found by OM and SEM. Figure 2 is a plot graph summarizing the EDS analyses, where the composition of several particles is plotted against their content in Ni. Since carbon was also observed within the particles, they were probably (Cr,Mo)23C6 carbides.

Fig. 2
figure 2

Chemical composition of the precipitates found within the banded structure in specimen B. The data are plotted as a function of the Ni content

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TEM images showed a plain grain structure interrupted only by rather straight and long twins bridging the whole grains from side to side (as shown in Fig. 1b). The as-received material is characterized by a limited amount of carbides both along grain boundaries and within grain cores. Carbide identification using CB-SAEDP showed that the vast majority of these carbides were of the M23C6 type, with M being essentially Cr, Mo, Fe ((Cr,Mo,Fe)23C6). This form of carbide has a complex cubic lattice with 92 metal atoms (Cr, Mo, Fe) and 24 carbon atoms per unit cell [22, 23]. In addition to M23C6 carbides, MC ((Cr,Mo,Fe)C) and M6C ((Cr,Mo,Fe)6C) carbides were also detected in isolated cases.

In summary, the microstructural analyses revealed that, in the as-received condition, the (Cr,Mo,Fe)23C6 carbides constituted the predominant population of particles in the investigated alloy.

Material creep tested at 700 °C

The microstructure of the alloy held at 700 °C for about 3.4 × 104 h (specimen 7C) is shown in Fig. 3 (OM and SEM) and Fig. 4 (TEM). With OM the microstructure still showed a banded structure similar to the as-received material (Fig. 3a). Rare and homogeneously distributed Ti(C,N) and (Cr,Mo,Fe)23C6 carbide particles were also observed in this specimen (Fig. 3b). A feature that distinguished specimen 7C from specimen B was the presence of two additional types of particles, observed both inside and outside the fine-grained bands. The first type of particle, that appeared similar to (Cr,Mo,Fe)23C6 carbides using OM, merely decorated grain boundaries. A second type of fine and elongated particle, darker with OM, was located not only at grain boundaries but also along twin boundaries. The SE-SEM images revealed grain boundaries significantly decorated with at least two distinct phases (Fig. 3c). EDS measurements, acquired from the bright particles, revealed that they were rich in Mo.

Fig. 3
figure 3

Specimen 7C. Typical metallographic microstructure (OM): (a) SEM, secondary electron images showing grain boundary and intragranular carbides and (b) a detail of a grain boundary decorated with (Cr,Mo,Fe)23C6 carbides from which a series of almost parallel strips of plate-like alternating δ-Ni3Mo phase proceeds inwards

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Fig. 4
figure 4

Specimen 7C. TEM, bright-field images showing two twin boundaries crossing a grain: (a) γ′ phase nucleation can be seen within the twin as well as along twin boundaries, both on incoherent and on semi-coherent twin boundaries (in the figure, the horizontal segments of the twin boundary); (b) shows a micrograph with the δ-phase starting from the grain boundary and advancing inwards in straight zip-like lines which are almost parallel to each other; (c) illustrates the γ′-phase hardening effect on the microstructure, through tangled sliding dislocations impinged by the γ′ particles

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TEM inspections at a far greater magnification level on the load-bearing region of the specimen revealed another interesting feature which distinguishes specimen 7C from specimen B: the appearance of a small but significant fraction of γ′-Ni3(Al,Ti), surrounded by dislocations. The γ′ particles were spherical with a diameter of about 60 nm. In addition, a depleted zone of γ was detected at the grain boundary (Fig. 4a). The γ-matrix and γ–γ′ interfaces appeared relatively free of dislocations (Fig. 4b). TEM inspections also revealed a lath-shaped δ-Ni3Mo phase extending from grain boundaries and twin boundaries into the matrix, following a zip-like path. In some cases, the growth of this phase started from existing carbides. The identification of the δ-Ni3Mo phase was carried out using the TEM–SAEDP technique. More specifically, the δ-phase, with a plate-like morphology, is the principal cause of the sometimes observed grain boundary waviness. TEM inspection showed that the growth of lath particles occurred to the detriment of γ′, as there is a γ′-depleted region around the lath phase (Fig. 4c). Since most Ni is consumed in the process of δ-phase formation, the γ′-phase can only form at a certain distance from the δ-phase surroundings, where Ni is still available.

In addition to the presence of γ and δ phases, specimen 7C showed grain boundaries which were not only decorated with M23C6, but also with other carbides, such as MC and M6C, this second being less abundant than the MC carbide. Prolonged exposure to 700 °C led to a considerable increase in the overall carbide population, especially at the grain boundaries and within the twins, compared with specimen B. Isolated M(C,N) carbo-nitrides of the type Ti(C,N) and (Nb,Ti)C were also observed in specimen 7C (Fig. 4c). Finally, in specimen 7C, dislocations were evenly distributed in well-defined coarse cells.

With SAEDP (Fig. 5) three different patterns are easily recognized. They can be attributed to γ (a 0 = 3.55 Å), γ′ (a 0 = 3.53 Å), and M23C6 (a 0 = 10.80 Å) [25]. SAEDP clearly showed super-lattice spots coming from the plate-like, or square-shaped, γ′-Ni3(Al,Ti) with an average edge length of 60 nm, denoting a semi-coherency with the γ-matrix.

Fig. 5
figure 5

Phase identification through SAEDP and related morphologies. SAEDP showing acting crystallographic planes from matrix γ, from the γ′ phase, Ni3(Al,Ti), and from carbide M23C6, where M = Cr,Co,Ni,Mo. (a) SAEDP showing acting crystallographic planes from matrix γ, and carbide MC, where M = Cr,Mo; (b) SAEDP showing acting crystallographic planes from matrix γ and from the δ phase, Ni3Mo

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Material creep tested at 800 °C

The analyses of the material exposed to 800 °C (specimen 8C) were focused on the specimen held for about 13,000 h at this temperature. Both in the load-bearing and load-free regions of the crept specimen, the microstructure revealed by OM and SEM inspection was characterized by the presence of a great amount of intragranular elongated precipitates distributed in lines, probably related to specific crystallographic orientations. Two types of particles at grain boundaries were visible in SEM micrographs (Fig. 6a). The first type, dark in SEM micrographs, proved to be rich in Cr and C, while the second type of particle (bright, in Fig. 6a), mainly located in the intergranular region of the finer-grained bands, were rich in Mo and C.

Fig. 6
figure 6

Specimen 8C. (a) SEM image showing both intra- and intergranular particles; (b–c) TEM images showing the lower fraction of γ′ particles with some intergranular needle-like carbides; (d) TEM image showing needle-like carbides and γ′ inside and at the twin boundaries

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In this specimen TEM images again showed the presence of the γ′ phase in a reduced volume fraction compared with specimen 7C (compare Figs. 6b and 4c). Furthermore, γ′ particles in specimen 8C were relatively coarse (about 255 nm in diameter) and cubic.

Hardness measurements

The plot graph of hardness as a function of creep duration, measured at the gripping ends of the crept specimens is reported in Fig. 7. In this graph, the HRC values were plotted even when lower than 20 (the standard minimum value for valid measurements) in order to compare them with the values reported in literature for the same alloy at roughly corresponding test temperatures [32]. Data points plotted in Fig. 7 clearly show that exposure for 104 h was not enough to reach peak hardness values at the lowest investigated temperatures of 593 and 648 °C. At about 700 °C, hardness increased rapidly compared with lower temperatures. The curve corresponding to exposure to 760 °C shows a wide peak region starting from 103 to 4 × 103 h, after which material overaging was evident. The material exposed to 800 °C passed its peak value at about 103 h. All the curves in Fig. 7 show that the experimental creep rupture times occurred in the overaging stage.

Fig. 7
figure 7

HRC plotted as a function of creep duration for different specimens creep tested at 700 °C (exp. 700 °C, solid upward triangular symbols in the diagram) and 800 °C (exp. 800 °C, solid circle symbol in the diagram). The hardness value in the as-received condition (exp. as-received, solid square datapoints) is plotted at t = 100 h. Hardness measurements made on the same alloy and creep tested at 593 °C (open square symbol), 704 °C (open upward triangle datapoints), 740 °C (open diamond symbol), 815 °C (open circle symbol) by Mankins et al. [11], are also plotted for comparison

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Due to the inhomogeneous microstructure observed in the as-received (B) and aged/crept conditions (specimens 7C and 8C), micro-hardness tests were also carried out to measure this property in regions where there were different microstructural features. The micro-hardness of specimen B in the coarse grain region, when depleted of secondary phase particles, averaged at 189 HV. Hardness was higher within the bands where precipitates were observed, reaching an average value of 199 HV. Correspondingly, the micro-hardness values of specimen 7C ranged from 280 to 290 HV outside the bands, in regions where second phases were relatively rare, and increased within the bands, ranging between 300 and 310 HV, particularly when the indentation load was applied in the small grain region. Finally, the micro-hardness of specimen 8C ranged between 222 and 263 HV. In agreement with the HRC measurements, the increase in micro-hardness with respect to the as-received condition is not as marked as for material exposed to 700 °C. A slight increase in micro-hardness, when moving from the core to the external surface of crept specimen 8C, was also observed.

Discussion

Microstructure

The microstructural modification and the hardness of the Ni–22Cr–12Co–9Mo alloy, subjected to long-term creep tests at 700 and 800 °C for different periods of time were characterized. The main features identified were the grain boundary particle population (both (Cr,Mo,Fe)23C6 carbides and the δ-Ni3Mo phase) and other types of carbides, such as M6C and MC, that were also detected within the austenitic grains. Moreover, the precipitation and coalescence of the γ′-Ni3(Al,Ti) phase clearly played an important role in the creep response of the alloy (see also [33]).

Discontinuously precipitated carbides were observed at grain and twin boundaries. It is well-known that the discrete and discontinuous distribution of grain-boundary carbides improves the mechanical properties of the material, in particular creep resistance, since these particles effectively pin the grain boundaries, thereby decreasing the grain boundary sliding which the material is sensitive to at typical creep temperature exposures [29, 34–36]. Nevertheless, it should be pointed out that these carbides can adversely affect strain distribution, creep life, and material ductility when forming a continuous network along grain boundaries.

As far as the intragranular carbides are concerned, Lewis et al. [16] have suggested that most precipitates within the grain nucleate preferentially on dislocations, generated by the growth of the same nucleating particles. For this reason, carbides continue to grow as the material is exposed to creep temperature and load.

Reasons for nucleation on grain boundaries and dislocations may be two-fold: (i) segregation of the constituent atoms of M23C6 on boundaries and dislocations; (ii) reduction in the strain-energy for nucleation in regions of localized strain. The origin of this effect has generally been ascribed to the differential thermal contraction of particle and matrix. A particle growing uniformly in a matrix of smaller atomic volume would experience a similar hydrostatic compression. This particle-matrix misfit may be accommodated by the diffusion of vacancies to the growing interface. An alternative mechanism for the relaxation of strain is by generating interstitial dislocation loops at the interface. Elongated screw segments which have undergone little modification in shape are frequently observed.

For the alloy studied in this paper, the most frequently detected intergranular (Cr,Mo,Fe)23C6 carbides were formed by a discontinuous precipitation. In general, the grain boundaries appear to have one side which adheres to the particle interface, while the other is characterized by a serrated appearance. This is confirmed by electron diffraction which shows a parallel orientation between the cubic lattices of matrix and precipitate on one side of the boundary (Fig. 5). OM and TEM inspection (Figs. 3, 4a, and 6b) showed thick plate-like carbides growing on one side of the boundary. Carbide particles grow at the maximum rate in the direction of the boundary, forming plates which have edges parallel to (111) planes [7, 8]. During creep, precipitation of M23C6 is also known to occur on non-coherent twin boundaries at approximately the same time as on grain boundaries [7, 8]. The carbides formed within the grains have a preferred interfacial growth along (111)-M23C6 parallel to (111)-γ [7, 8]. Prolonged exposure to 800 °C induced a massive dissolution and coalescence of the (Cr,Mo,Fe)23C6 carbides which eventually formed a intergranular film morphology. This continuous film reduced both the effect of grain boundary pinning and the possibility of intergranular embrittlement [37–40].

Another important factor for the creep resistance of the investigated alloy is the presence of the γ′ phase. After exposure to 700 °C, the volume fraction of the γ′ phase is high enough to slightly improve the mechanical properties of the alloy. Comparing the microstructure of specimens 7C and 8C (from 700 to 800 °C), the γ′ phase particles coarsened more than four-fold, from ∼60 to 255 nm in diameter. Furthermore, at 800 °C, cubic rather than spherical particles were observed (compare Figs. 4 and 6). The differences observed in γ′ particles were clearly a consequence of the higher-temperature exposure which both induced a partial coalescence in favor of the initially bigger particles and a re-dissolution of the initially smaller particles. The coalescence of γ′-particles in both specimens 7C and 8C was faster in the load-bearing regions of the crept specimens than in the head of the creep sample. This effect could possibly be interpreted as a stress-enhanced diffusion rate or as a dislocation-driven coarsening mechanism, which is also observed in other materials like steel [5–7, 12, 28, 34]. Moreover, the cubic γ′ precipitates indicated that the γ′/γ interface might still be semi-coherent. However, the coarsened γ′ phase still provided matrix strengthening by dislocation pinning in agreement with Dahotre and co-workers [41], as shown in Fig. 8.

Fig. 8
figure 8

BF-TEM of specimen 8C. Coarsened γ′-Ni3(Al,Ti) phase with sliding dislocations pinned and bounded

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The precipitation of γ′ has been investigated in several Ni-based alloys. Recently, in an Ni–25Cr–20Co alloy with Mo, Nb, Ti, and Al addition [42, 43], precipitation has been observed at temperatures lower than 620 °C. Thermodynamic calculation on the same alloy indicated that γ′ was not an equilibrium phase at temperatures between 821 and 1091 °C [43] and it was experimentally observed that higher temperatures (greater than 650 °C) favored a very fast coarsening of γ′ that eventually transformed either into γ″ or orthorhombic δ. As a result, the alloy loses its strength, since the improvement in the yield strength is mainly due to the extensive precipitation of γ′ although the γ″ and δ-phases have similar but weaker effects [32, 42–45].

In the investigated Ni–22Cr–12Co–9Mo alloy, exposure to 700 and 760 °C induced the precipitation of small amounts of fine rounded γ′ particles. At 800 °C, these particles coarsened and took on a cubic-like shape although they were less abundant. At this temperature only a secondary improvement in the hardness of the material was observed.

Creep behavior

Some of the creep strength data available in literature for the Ni–22Cr–12Co–9Mo alloy are summarized in Fig. 9. Most of the data available refer to tests performed at 800 °C or higher temperatures. In Fig. 9 the initial applied stress is plotted versus the Larson Miller Parameter (LMP = T(log t + 20), where T is the test temperature expressed in Kelvin and t is the time to rupture, in hours). More specifically, two sets of literature creep data (datasets A and B) have been proposed by Schubert et al. [46] which refer to creep tests performed at temperatures ranging from 800 to 1,000 °C. It can easily be seen in Fig. 9 that one of the two materials (dataset A) is more creep resistant at high-LMP values. Other creep results considered in the plot graph are those reported by Mankins et al. [13]. It is clear that at low LMP the creep properties reported by Mankins for the lowest test temperatures (700 °C or lower) tend to be better that those obtained from tests at higher-test temperatures.

Fig. 9
figure 9

Comparison between experimental (ECCC data) and literature creep data in terms of the Larson Miller Parameter. Experimental creep data at 800 °C are plotted using a full circle as a symbol while open and full triangles are the symbols used to represent, respectively, interruption and predicted rupture times at 700 °C. The stress versus LMP curve used to predict the creep rupture times of the investigated material is represented as a continuous line. Creep data results from Mankins et al. [11] are represented with full squares, while the results for the two heats reported by Schubert [44] are represented with crosses and open diamond symbols

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The creep results for tests performed by a consortium of Italian Laboratories within the framework of a European Creep Collaborative Committee (ECCC), already published in [12], are also included in Fig. 9. It should be observed that creep rupture was reached only for creep tests carried out at 800 °C, while those which refer to 700 °C are only expected rupture times [12]. The same figure also shows the trend line used to define the experimental plan for ECCC tests [12], with which the rupture times at 800 °C are perfectly aligned. On the contrary, the rupture times estimated at 700 °C are about one order of magnitude longer than the predicted ones, related to a significant improvement in creep behavior at this test temperature. Therefore, creep strength data obtained at higher temperatures should be very carefully considered when extrapolating the creep strength of the alloy at 700 °C since the extrapolation could lead to a substantial underestimation of the creep life.

This can be correlated with the microstructural instability of the alloy in the investigated temperature range. As illustrated above, the microstructure of the Ni–22Cr–12Co–9Mo alloy changed at different creep temperatures with respect to the as-received condition. The main differences lie in the presence and distribution of carbides at grain boundaries, in the volume fraction and in the size both of the γ′ phase and the δ-phase. The effects of these phases on creep properties have been previously presented and could explain the improved creep behavior of the alloy at 700 °C compared with 800 °C. The fraction and geometry of the γ′-phase precipitated at 700 °C contributed to an increase in the mechanical properties of the alloy, although its partial dissolution and coarsening severely reduced the strengthening effect of γ′-phases at 800 °C. Furthermore, exposure to 800 °C induces the coalescence of the (Cr,Mo,Fe)23C6 carbides in crept material . These carbides may eventually form an intergranular film morphology, which could modify grain boundary sliding conditions and the final fracture mechanism. Finally, the effect of the δ-phase, found only in specimens crept at 700 °C, is believed to contribute to the strengthening of the alloy, as has also been reported by other authors [47–50]. The specific zip-like morphology, starting from the grain boundary region and proceeding into the grain effectively pinned the sliding dislocations. The pinning effect is competitive compared with the effect of the fine square-shaped γ′-phase particles. The γ′ pinning strength is believed to be lower due to its relatively low fraction compared with the δ-phase.

Conclusions

The microstructural evolution after different creep tests on an Ni–22Cr–12Co–9Mo alloy was characterized using TEM, SEM, and OM techniques. The microstructural investigations revealed the following main features:

  1. 1.

    In the as-received condition the material was characterized by a banded structure, which contained Cr- and Mo-carbides; this banded structure had fine grains yielding higher-hardness values compared with the surrounding coarser grained regions. More specifically, M23C6 carbides were always present, lying within the grain, on twins and along grain boundaries.

  2. 2.

    Creep exposure to 700 and 800 °C led to an extensive intragranular formation of γ′-Ni3(Al,Ti) precipitates. The size of these precipitates ranged between 60 and 255 nm depending on the test temperature. Their morphology was spherical at the lower test temperature and cubic at the higher-test temperature. This effect is consistent with the observed increase in the hardness of the material, with respect to the as-received material.

  3. 3.

    At 700 °C the δ-phase (Ni3Mo) competitively formed with the γ′-Ni3(Al,Ti). The δ-phase started to form at the grain boundaries and proceeded in a zip-like path inwards. The presence of tangled sliding dislocations all around the plate-like δ-phase indicated that they are strengthening particles for the material. At 800 °C no trace of the δ-phase was found. Other precipitates observed were carbides such as M23C6—(Cr,Co,Ni,Mo)23C6 and, less frequently, MC—(Cr,Mo)C. The amount and morphology of these carbides changed with the exposure temperature. The Cr-, Mo- carbo-nitride CrMo(C,N) was rarely detected by microstructural inspections, while no MC, M6C carbides were detected.

  4. 4.

    The microstructural instability of the alloy, in particular the large differences in size and distribution of γ′-precipitates at 700 and 800 °C, can account for the improved experimental creep resistance at 700 °C compared with expected creep life. This suggests that creep data obtained at the higher temperature should be very carefully considered when extrapolating the creep strength of the alloy at 700 °C, since they could lead to a substantial underestimation of the creep life.