A. Microstructure states of the alloys after heat treatment

The as-cast microstructure of the investigated superalloy is shown in Fig. 1. As expected, the superalloy exhibited high segregation that was confirmed by a dendritic structure. The dendritic structure exhibited a two-phase γ/γ′ microstructure with a dendrite segregation pattern and a limited amount of γ/γ′ eutectic occurring in some interdendrite regions [Figs. 1(a), 1(c), 1(e), and 1(g)]. The dendrite cores appeared white, the interdendritic regions are black, and the eutectic areas are medium gray. This microstructure caused during solidification, some elements partition to the dendrite core, while other elements tended to accumulate in the interdendritic liquid and then solidify as the interdendritic and eutectic regions. This was reflected in Yu's findings that Al, Ti, Ta, and Mo are reported to segregation to the interdendritic region, whereas, W segregation to the dendrite core.^{Reference Yu, Liu, Zhao, Zhang, Zhang and Fu24} From this figure, it also could be concluded that when the carbon level was high [Figs. 1(e) and 1(g)], the carbides can be observed in the microstructure by OM images. This problem had been much discussed in author's other literature.^{Reference Yu and Liu17}

FIG. 1.

OM images showing dendrite microstructure with carbon additions: 0.006% (a-as cast, b-heat treatment), 0.045% (c-as cast, d-heat treatment), 0.085% (e-as cast, f-heat treatment), and 0.15% (g-as cast, h-heat treatment).

Figure 2 shows the average EPMA compositional results for the alloys in Table I. From these results, it was apparent that the elements such as Co and W all tend to segregate toward the dendrite core regions during solidification, as indicated that the greater than unity. The greater the segregation coefficient, the greater the degree of segregation. A segregation coefficient was less than unity, which indicated that the element segregated to the liquid during solidification, ultimately solidifying in the interdendritic regions.

FIG. 2.

The segregation behavior of alloying elements.

Elements that segregated in an inverse manner had k′ values less than unity, such as the γ′ forming elements like Al, Ti, Mo, Ta, and Cr. These elements were present in greater concentration in the last liquid phase to solidify and result in large eutectic γ–γ′ phases to form within the interdendritic regions. The element Co that did not favorably segregate to either region has k′ value close to unity and represents the ideal segregation compositions examined here. These results were similar to the study.^{Reference Fuchs7} It was clear that W was the most severely partitioned element.

The solution temperature was chosen to ensure that γ′ phase and γ–γ′ eutectic in the as-cast microstructure are fully solutioned into matrix, and to reduce the degree of chemical segregation due to partitioning of some of the elements to the dendrite cores and the interdendritic regions, while being sufficiently below the solidus temperature to avoid incipient melting. In this study, the DSC testing of the as-cast samples indicated that the liquidus was from 1360 to 1367 °C, and the solidus was from 1326 to 1334 °C. The γ′-solvus was from 1282 to 1293 °C. The standard heat treatment was based on this DSC results. Two step aging (at 1080 and 870 °C/20 h) was applied for the investigated superalloy.

After heat treatment, all the elements turned to exhibit much more uniform distributions with only limited variations. The dendritic structure was less visible after heat treatment [Figs. 1(b), 1(d), 1(f), and 1(h)]. No obvious γ–γ′ eutectic precipitates were found either in the dendrites or in the interdendritic regions as shown in the typical microstructure micrograph.

B. γ′ morphologies of the alloys after heat treatment

Figure 3 shows the γ′ morphologies of the as-cast and heat treatment. As it can be seen γ′ phase possesses cubic morphology on the average and after heat treatment the γ′ phase became more regular than the as-cast sample. It also showed that the more regular the shape of γ′ precipitated, the more even the size of γ′ precipitated. The average size of γ′ for as-cast samples is from 245 to 350 nm, whereas, the average size of γ′ for heat treatment was from 435 to 500 nm. The equilibrium shape of the γ′ particles was probably determined by a balance between surface energy and strain energy. The morphology of the γ′ precipitates during the early stage of growth was a function of both misfit and particle size, since the transition from spheres to cubes will depend on the total matrix/precipitate strain present due to the lattice mismatch.^{Reference Pollock, Merphy, Goldman, Uyam, Antolokch, Stusrud and Mackay25} Taking into account a typical interfacial coherent energy density γ_k ≈ 10 mJ/m², spheres should transform to cubes with increasing precipitate size to minimize precipitate surface area.^{Reference Yu, Sun, Zhao, Jin, Guan and Hu14} The formation of γ′ plates aligned statistically along $\left\langle {100} \right\rangle$ directions at high temperature annealing with no external load is due to the lower surface area of plates as compared to cubes for high volume fraction of precipitates.

FIG. 3.

SEM images showing γ′ of alloys with carbon additions: 0.006% (a-as cast, b-heat treatment), 0.045% (c-as cast, d-heat treatment), 0.085% (e-as cast, f-heat treatment), and 0.15% (g-as cast, h-heat treatment).

It indicated that the heat treatment changed the morphology and size of γ′. The mismatch between the strengthening γ′ and the γ matrix was essential to determine the morphology and resulting the deformation mechanism of the alloy. The most effective precipitated microstructures had large negative misfits at the desired surface temperature. The equilibrium shape of the γ′ particles was probably determined by a balance between surface energy and strain energy, and can be a function of particle size, resulting from a change in the equilibrium matrix and γ′ composition. Furthermore, it had been shown that the shape of the precipitate varies with the degree of lattice mismatch; the shape changed from a sphere to a cuboid with increasing degree of lattice mismatch.^{Reference Yu, Sun, Zhao, Jin, Guan and Hu14}

It was well known that a distribution of immobile particles in a solid matrix tends to lower its interfacial free energy by transport of matter from the smaller to the larger members of the distribution, thereby diminishing the total particle number but increasing the average particle size. In other words, the surface energy of the precipitate was reduced by coarsening, small ones dissolve, and large ones grow at the expense of the smaller ones. This coarsening process is known as competitive growth or Ostwald ripening clearly followed kinetics which were in general dependent upon the nature of the particle dispersion, provided that the rate-limiting step was other than the transfer of matter across the particle–matrix interface.^{Reference Ges, Fornaro and Palacio26}

C. Carbide morphologies of the alloys after heat treatment

Figures 4 (a) and 4(b) shows the carbide morphologies of 0.006% C as-cast samples and heat treatment samples. It could be seen that heat treatment shows significant influence on carbide morphology. In as-cast sample, acicular carbide can be observed. After heat treatment, the carbide morphology became glomeration. Figures 4(c) and 4(d) shows the carbide morphologies of 0.045% C as-cast samples and heat treatment samples. It could be observed that most of carbides in the 0.045% C as-cast samples were blocky and sheet-like, and the size of carbide was above 5–6 µm. After heat treatment, the carbide morphologies were particle and small blocky, meanwhile, it could be observed that the size of particle carbide was above 150 nm. Figure 4 shows that the size of carbide became much smaller and the distribution of carbide became much more dispersive after heat treatment.

FIG. 4.

SEM images showing carbide morphologies of deep etched sample 0.006% (a-as cast, b-heat treatment), 0.045% (c-as cast, d-heat treatment).

Figure 5 shows that the carbide morphologies of 0.085% C as-cast samples were sheet-like, blocky, and a little acicular, however, the carbide morphologies became particle and small blocky after heat treatment. In as-cast sample, the size of sheet-like carbide was above 10 µm, and the size of blocky carbide was above 3 µm. In heat treatment sample, the size of blocky was above 2 µm, and the size of grainy carbide was above 100 nm. It also could be concluded that the size of carbide became smaller and the distribution of carbide became much more dispersive after heat treatment.

FIG. 5.

SEM images showing carbide morphologies of deep etched sample 0.085% (a, b-as cast; c, d-heat treatment).

Figure 6 shows the carbide morphologies of 0.15% C sample. It could be observed that the most of carbides in the 0.15% C as-cast sample were Chinese-script like morphology, as reported previously,^{Reference Al-Jarba and Fuchs16} which fully formed the carbide network in the interdendrite region. Stereo pair images of these scripts carbide showed rods at the center of the dendritic carbides pointing out of the image. From the center, the carbides started to form directions starting with sheets and then splitting to form rods. Figure 6(c) shows the morphology of acicular carbide after heat treatment, and it illustrated that the carbides of heat treatment sample were the decomposition of debris of as-cast samples. Figure 6(d) shows the grainy carbide of heat treatment sample, the size of which was above 100 nm.

FIG. 6.

SEM images showing carbide morphologies of deep etched sample 0.15% (a-as cast; b, c, d-heat treatment).

After heat treatment, the carbide morphology became much simpler, and the size of carbide was obviously less than the carbide of as-cast samples, meanwhile, the distribution of carbide much more dispersive than the carbide of as-cast samples. Carbides distributed mainly in the interdendritic region, and only when the carbon level was high, the carbide can be observed in the dendrite zone. After heat treatment, Chinese-script like carbide networks became smaller and more spherical [Fig. 5(c)]. These carbide spheres were observed much more frequently in the heat treatment samples [Figs. 4(d), 5(c), 5(d), and 6(d)]. The part of the Chinese script networks in 0.15% C sample retained section of their morphology after heat treatment [Fig. 6(c)].

As previous study reported^{Reference Paulo, Azevedo, Renato and Carlos9} that the carbides of as-cast sample are MC-type carbides. To investigate the chemical composition of carbide after heat treatment, representative WDS was measured, and the results were obtained as Fig. 7. Typical WDS results for small, spherical carbide after heat treatment are shown in Table II. Figure 7 and Table II show that the small, spherical carbides were found to be Ta-rich and were, therefore, considered to also be MC-type carbides.^{Reference Yu and Liu17,Reference Wasson and Fuchs27}

FIG. 7.

WDS analysis of carbide after heat treatment (0.15% C): (a) Small, spherical carbide; (b) EDS result.

TABLE II.

Typical WDS composition (wt%) results for small, round Ta-rich MC type carbide phase observed in heat treated sample (0.15% C).

A typical SEM micrograph showing grainy carbide microstructure of the heat treated sample is shown in Fig. 8(a). Together with the WDS results shown in Fig. 8(b), which can be deduced that these precipitates were chromium-rich M₂₃C₆ carbides. It seemed that the precipitation of nanoscaled M₂₃C₆ carbide was one of the most important phenomena in the heat treatment alloy. Thus, it was important to examine the formation mechanism and stability of such M₂₃C₆ carbide during heat treatment. Comparing Figs. 5(d) and 6(d), it could be found that fine M₂₃C₆ carbide precipitates with the size from 50 to 200 nm can be observed apart from the γ and γ′.

FIG. 8.

WDS analysis of carbide after heat treatment (0.15% C): (a) Grainy carbide; (b) WDS result.

In the C modification, script carbide networks that remained intact after heat treatment decomposed into smaller, spherical, or faceted MC-type carbides due to heat treatment. Secondary, Cr-rich carbides (M₂₃C₆ carbide) formed on the surfaces of some MC carbides in all three of the modifications [Figs. 4(d), 5(c), and 6(b)]. These types of carbides were observed more frequently in high C level sample. The decomposition of MC carbide appeared to occur following a reaction (1) and (2), a

(1)$${\rm{MC}} + \gamma \to {{\rm{M}}_{23}}{{\rm{C}}_6} + \gamma \prime \quad ,$$

(2)$$\left( {{\rm{Ti,Mo}}} \right){\rm{C}} + \left( {{\rm{Ni}},{\rm{Cr}},{\rm{Al}},{\rm{Ti}}} \right) \to {\rm{C}}{{\rm{r}}_{21}}{\rm{M}}{{\rm{o}}_2}{{\rm{C}}_6} + {\rm{{\rm N}}}{{\rm{i}}_3}\left( {{\rm{Al}},{\rm{Ti}}} \right)\quad .$$

reaction widely reported in literature.^{Reference Yu, Liu, Zhao, Zhang, Zhang and Fu24} Carbide reactions during the aging at 650–1080 °C caused the precipitation of the discrete M₂₃C₆ grainy carbides. On the aging with a relatively long time (20 h) at 870 °C, carbides had enough time to provide the carbon diffusion and reaction with the Cr to form M₂₃C₆ (M rich in Cr).