Metallography 251

The Microstructure of Superalloys

P. S. KOIVXL

Utlim Carhide Corporation, Materials Systetm Dicisim, Speedway Lahoratory, Indianapolis, Iudiarza

Some recent thin-foil transtmssion electron microscopy studies that contribute to an Increased understanding of the microstructure of nickel- and cobalt-base superalloys are discussed. Emphasis has been placed on the interaction between imperfections and precipitates that cannot be studied hy x-ray and replica methods. The effects of matrix stacking fault energy on the dislocation suhstructurc and the role of different dislocation substructures in the nucleation of precipitatcs are summarized. Attention has been given to the intragranular precipitation of MC carbides in association with stacking faults in nickel-base alloys.

Recent data on thc behavior of y-phase in simpler nickel-base alloy systems are discussed with respect to the possible application of these results in understanding the hehavior of commercial y-hearing superalloys. Finally, brief consideration is givcn to recent studies of order-disorder reactions in nickel-base alloys.

1. Introduction

In the last twenty vears, the metallurgy of nickel- and cobalt-base alloys has _ . received great attention becausc of the high-temperature applications of these materials. A great deal is known and has been written about structures in these superalloys, and correlation of structure with properties has been extensive. However, the bulk of this development has involved empiricism. It would seem that future advances in this area of technology wil1 be very dependent on the degree to which precipitation of phases (both strengthening and deleterious) is understood and controlled in these systems. Understanding of precipitation in superallovs in turn requires that al1 available metallographic techniques be employed. In other areas of physical metallurgy, the ability to observe imper- fections directly in the electron microscope has led to an increased appreciation of their role not only in deformation but also in nucleation of precipitates. In superalloys, however, this type of approach has not often been used, mainly because metallographic techniques such as thin-foil transmission electron microscopy have been only partially exploited. Manv articles have reviewed

Metallography, 1 (1969) 251-285

Copyright ((2 1969 by hmerican Elsevier Publishing Company, Inc.

252 P. S. Kotval

the state of knowledge concerning structures in superalloys. XIost of these papers have concentrated on discussing the type, distribution, and morphology of phases in superalloys and some parameters controlling their occurrence but not on the interaction of phases and lattice defects.

In the above context and by way of introduction it is relevant to consider briefly the composition of these alloys and what is known, generally, about the microstructure of wrought nikkel- and cobalt-base superalloys from studies using optical, replication, and x-ray techniques. The main commercial alloys contain 1.5 to 20% chromium for oxidation resistance, varying additions of iron, molybdenum, and tungsten for solid solution strength, and lesser additions of titanium, aluminum, and, in some cases, columbium as precipitation hardening constituents. Table 1 shows the nomina1 compositions (in weight percent) of a few of the main commercial wrought superalloys.

The principle microstructural differente between commercial alloys and the simpler Ni-Cr-Ti matrices is the formation of a series of carbide phases in the commercial compositions. These carbide phases are accompanied, in the majority of commercial nickel-base alloys, by a strengthening phase such as the ordered f.c.c. y phase. Essentially, the precipitation of these phases is, in addition to the solid solution strengthened f.c.c. matrix, the main strengthening mechanism in superalloys. In cobalt-base alloys, strengthening by y precipitation has only recently been utilized. The role of carbides as strengthening agents is twofold. Primary carbides, retained from the as-tast condition, are usually large and fulfill only a minor dispersion strengthening role. Carbides precipi- tated during an aging reaction have a more direct role. Sims has suggested that precipitation at the grain boundaries, which usually involves a globular M,,C, type of carbide, inhibits grain boundary sliding and helps in retarding the onset of fracture processes. In some cases, however, discontinuous precipitation of M,,C, carbides at the grain boundaries has also been observed. Intragranular precipitation of carbides has a far more significant role in strengthening, but it is only recently, with the use of thin-foil techniques that permit a study of the interaction between precipitates and matrix imperfections, that this has been made clear.

In the present article, emphasis is placed on the relationship of structural imperfections, such as dislocations and stacking faults, to precipitates in super- alloy matrices, as revealed in recent studies by modern metallographic techniques. In many of the observations, both from first-hand studies and from those of other workers, simplified versions of commercial alloy compositions have been utilized to illustrate a particular argument. In no way is this article a com- prehensive survey of precipitation in superalloys; rather it is intended to be a discussion of recent information concerning the microstructure in nickel- and cobalt-base alloys which increases our understanding of the behavior of super- alloys and, it is hoped, provides a pointer to future developments.

The Nicrostructure of Superalloys 253

254 P. S. Kotval

TABLE 11

SOIIE TYPICAL ETCHANTS AND ELECTROLYTES USED IN THE METALLOGRAPHY OF SUPERALLOYS

Tempe- Composition Voltage rature Remarks

Electropolishing (1) Methanol-10 4, sulfuric 40 Room Useful for most nickel- solutions acid and cobalt-base

alloys. Stainless-steel or graphite cathodes.

(2) Ethanol-10 70 perchloric 20-22 32F Nickel- and cobalt-base

Chemical etchants

Electrolytic etchants

Extractive electrolytes

acid

(3) Acetic anhydride-10 Cr perchloric acid

(4) Marbles reagent: Cupric sulfate 40g Hydrochloric acid 40 ml Water 40 ml

(5) Lactic acid 30 ml Hydrochloric acid 20 ml Njtric acid IOml

(6) Aqua regia

(7) Methanol-10 1 sulfuric acid

(8) Water-10 ,) oxalic acid

(9) Hydrochloric acid-10 0, hydrogen peroxide

20

alloys. Slower elec- tropolishing action than electrolyte in (1). Suitable for electrothinning.

32F Suitable for electro- thinning, particularly samples with large second-phase parti- cles.

- Room Useful for revealing genera1 microstruc- ture in both optica1 and replica techni- ques.

- Room For y-containing alloys.

- Room Immersion etching of cobalt-base alloys.

6 Room Useful for most nickel- and cobalt-base alloys.

4 Room Suitable for y-bearing alloys.

2-3 Room Stainless-steel catho- des. Suitable for cobalt-base alloys.

(10) Water-10 yu hydrochloric acid

(11) Methanol-10 9,) hydrochloric acid

4-1/2 Room For extracting y and intermetallic phases for x-ray diffraction.

4 Room For extraction of car- bides for x-ray dif- fraction. Tantalum cathode.

lhe MLcrostructure of Superalloys 255

2. Metallographic Techniques and Their Limitations

The bulk of the work on the microstructure of superalloys has involved extensive use of optica1 microscopy and a combination of the electron microscopy of replicas plus x-ray or electron diffraction of extracted phases. Table 11 gives a brief list of the etchants and electrolytes used in these methods and the experimental conditions. Detailed information on these metallographic tech- niques is available elsewhere.3J

In the majority of investigations, the method adopted has been to determine the shape, distribution, and size of precipitated phases as revealed by a replica (usually plastic) taken from the polished and etched surface of a sample. Residues extracted in various media from the same sample are used for x-ray diffraction, and a correlation is then attempted between the microstructural observations from the replication techniques and the structural data from x-ray diffraction. This method of investigation has provided the bulk of the present-day under- standing of precipitation in superalloys, and it appears that many of the restric- tions in understanding and controlling the microstructure of superalloys are a direct function of the limitations of these techniques, which may be summarized as follows:

1. The microstructure revealed by the replication technique is controlled by the type of etchant used, and hence the reproducibility and reliability of results become questionable.

2. In the case of the same phase being precipitated in several modes in a given sample, the mere confirmation of its existente by x-ray diffraction methods stil1 does not permit a proper identification of each of the morphologies of that phase in a replica photomicrograph.

3. Very little information regarding precipitate-matrix or precipitate- dislocation interactions is obtained.

4. Different extractive media have to be iteratively used for different phases, and hence the absente of a phase in an x-ray diffraction pattern does not necessarily imply its absente in the original sample. This ambiguity due to extractive media applies in addition to the obvious limitation of the x-ray tech- nique when the volume fractions of precipitate are small.

5. In a complex multiphase superalloy, x-ray diffraction patterns often have lines that can be ascribed to more than one phase, because of the marked isomorphous substitution and variation in stoichiometry and hence in lattice spacing that is possible for most phases in superalloys.

The technique of thin-foil transmission electron microscopy obviates many of the limitations enumerated above. The technique allows a study of the role of imperfections in the nucleation and growth of precipitated phases to be made.

P. S. Kotval

Equally important is the capability of in situ identification of a particular phase by selected area diffraction so that the ambiguities of the other (indirect) tech- niques are eliminated. The results presented throughout the text make both these advantages self-evident.

3. Stacking Fault Energies and Dislocation Substructure in Superalloys

Dislocation substructures in the matrices of nickel- and cobalt-base alloys vary widely with composition, primarily as a function of stacking fault energy (SFE). Considerable work has been done on binary alloy systems of nickel- and cobalt-base alloys5-s to study the effect of alloying additions on stacking fault energy. In most cases, these stacking fault energy measurements are confined to metals and alloys of little commercial value. Nevertheless, knowledge of the effect of a given alloying addition on stacking fault energy, and hence on dis- location substructure, permits some understanding of the dislocation sub- structure in compositions of commercial importante.

As is wel1 known, a simple lattice dislocation of the a/2 (110) type in the face-centered-cubic lattice can dissociate in the { 1 1 1} glide plane as follows:

42 [lO] = a/6 [121] + 46 [211]

This dissociation lowers the energy of the resultant extended dislocation. The two partial dislocations that bound the extended dislocation wil1 repel each other and provide an intervening region of stacking fault. The finite energy of the stacking fault supplies a constant attractive force, and when this balances the elastic repulsive forces the total energy is minimized at some equilibrium separation. In a simplified farm the equilibrium separation, d, is given by

Ga . k a d = 4% . y

where G = shear modulus. a = cubic cel1 edge.

K = constant for a given material which is dependent on b, the Burgers vector.

y = stacking fault energy (SFE).

It is thus clear that the stacking fault energy is the principle parameter that influences the extent of dissociation of dislocations in an f.c.c. matrix. Methods for measuring stacking fault energy are described in detail by Christian and Swann.s

The Microstuuctuue of Superalloys 257

It must be mentioned that the various techniques by which the stacking fault energy can be measured in nickel and its alloys have application only within certain ranges of stacking fault energy.* High values of stacking fault energy cannot be measured by the Node method,g whereas in materials of lower stacking fault energy this method can be utilized. The converse restriction applies to Seegers method, l based on the stress for cross-slip where the method is most applicable to higher stacking fault energy matrices. Thus caution has to be exercised in comparing values of stacking fault energy measured by different methods. Further, in the case of stacking fault energy values measured by the Node method, the corrections applied significantly alter the values, and this has to be borne in mind when values from different investigations are compared.

By contrast, Smallmans methodll of estimating stacking fault energies from the rolling texture developed in a given material has the advantage of being apphcable to a whole range of alloys and stacking fault energies.

Figure 1 shows data from recent work by Beeston and France;12 y/Gb values determined from the texture-analysis techniquel have been plotted against atomic percent solute additions of Fe, Cr, Al, and Ti in pure Ni. Each of these solute elements is found to lower the stacking fault energy of nickel. Table 111

L

14-

m 0 8- - x _

+J 6-

4-

2-

?

0 1 1 1 1 1 1 1 10 20 30 40

ATOMIC /o SOLUTE 4

FIG. 1. Values of yGb versus composition (atomic percent soiute) for nickel-base binary alloys. Data from Beeston and France.?

258 P. S. Kotval

TABLE 111

SUMMARY OF BEESTON AND FRANCES DATA~

Solute element Decrease in y/Gb per atomic percent

solute added to nickel

CO 0.14 Cr 0.25 Ti 0.72 Al 0.23 Fe 0.08

summarizes Beeston and Frances results on the decrease in y/Gb per atomic percent addition of a particular solute element. It is noticeable that Co, which is widely accepted as the element most capable of lowering the stacking fault energy of nickel alloys, is less effective than Ti, Al, or Cr. Beeston and Francel have estimated percent reductions in r/Gb for certain Nimonic alloy compositions on the basis of their data on binary alloys and found a correlation between the estimated r/Gb values and the shear strength of a series of alloys.

However, measurements of stacking fault energy by Douglass et uL.,~ using dislocation node-radius measurements, indicate that the effects of a solute element in a binary system are not the same when a ternary system is involved. In a nickel-20% chromium alloy the stacking fault energy is lowered to -16 ergs/cm2. However, in a Ni-20Cr40Fe alloy the stacking fault energy has a much higher value of -60 ergs/cm 2; that is, the stacking fault energy increases with addition of iron to a nickel-20% chromium alloy. Thus, estimation of stacking fault energy for commercial alloy compositions based on data from binary systems has to be approached with caution.

NO systematic data on stacking fault energy variations in alloys of commercial composition are available. However the range of stacking fault energy within which a given superalloy matrix falls is manifested by the type of dislocation substructure observed in lightly deformed thin foils.

Figures 2,3,4, and 5 show the variation of substructure in thin foils of (1) pure nickel, (2) Ni-20Cr-5Fe-8Mo-3.5Cb-O.O5C (nomina1 composition of Inconel- 625 alloy), (3) Co-13Ni-20Cr-2A1-2Ti-O.O2C, and (4) Coo14W-20Cr-3Fe-O.05C after 5 % strain at room temperature. In Fig. 2 a cellular dislocation substructure is seen that typifies the case of a high stacking fault energy matrix. Cross-slip in this type of matrix is easy,2J3J4 and there is a tendency for cel1 wal1 formation rather than for the formation of coplanar groupings. The cel1 size decreases with decreasing stacking fault energy.2 Figure 3 shows the dislocation substructure of a medium stacking fault energy matrix. Cel1 structure has been replaced by coplanar groupings of dislocations. In this type of matrix cross-slip is more difficult than in the high stacking fault energy matrix, and it becomes pro-

The Microstructure of Superalloys 259

FIG. 2. Thin foil showing dislocation substructure of pure ni at roon 1 trmpera ture. From Kotval.14

ckel. Def< wd 4 0

FIG. 3. Thin foil showing dislocation substructure of Inconel compo: Gtion: Ni- -2OCr-8Mo-5Fe-3.5Cb-O.05C) solution heat-treatel quench ed, and dc zformed 50: at room temperature. From Kotval.

allo! 1 at

i 625 2280.

lomir \vate

ia1 r-

260 P. S. Kotval

FIG. 4. Thin foil showing dislocation substructure of Co-13Ni-20Cr allo?_, solution heat-treated at 22OOF, water-quenched, and deformet tempel -ature. From Kotval.*

-2 :X1-2li d _5,, at

4 1.02C -oom

FIG. 5. Thin foil showing dislocation substructure of a Co-I4W-2 alloy, I solution heat-treated at 2150F, water-quenched, and deforme temper ature. From Kotval.14

!O( Zr-3Fe- d 40: at

.OSC oom

The Micvostvucture of Superalloys 261

Ni-16Cr

Ni-ZOEr-40Fe

Ni-20Cr-17Fe -8Mo-.05C

Ni-2OCr-5Fe-8Mo -3.5Cb-.05C

Co-14W-20Cr

0 IncreaSing stacking fault energy

FIG. 6. Correlation between stacking fault energy and type of dislocation substructure in several nickel- and cobalt-base alloys. From Kotval.13

gressively more difficult as stacking fault energy decreases with alloying additions. Dislocations tend to be confined to their own slip planes. This type of matrix typifies the structure of most commercial nickel-base alloys. With further decreases in stacking fault energy, the substructure shows widely estended dislocations with marked dissociation as shown in Figs. 4 and 5. In these 10~ stacking fault energy cobalt-base alloys, cross-slip is almost totally inhibited and a high stacking fault density is observed. Alloys with the highest stacking fault density usually have a tendency toward the f.c.c. --f h.c.p. transformation in the matrix.

The correlation between the three (high, medium, or 10~ stacking fault energy) types of dislocation substructure described above and alloy com- positions of varying stacking fault energy is represented in Fig. 6. The precip- itation reactions in a given alloy and the type of precipitate-dislocation inter- action vary significantly with the stacking fault energy of the matrix. In the following sections, precipitation in superalloys is discussed with reference to the influence of the type of dislocation substructure in different types of matrix.

4. Precipitation on Dislocations

Grain boundaries have long been established as the preferred regions of precipitation in nickel- and cobalt-base superalloys. By thin-foil studies using transmission electron microscopy, it has also been widely observed that precip-

3082-2

262 P. S. Kotval

itates nucleate on dislocations in preferente to other sites within a grain. The ease of nucleation on dislocations varies with different precipitates, with the temperature at which precipitation is taking place, and with the supersaturation of solute in the matrix. The type of matrix, based on stacking fault energy, also has an important effect in governing the role of dislocations in nucleating precipitation.

Figure 7 shows a thin-foil micrograph of HASTELLOY alloy X (Ni-20Cr- SMo-17Fe-0.05C alloy) in which a pre-aging strain has produced undissociated dislocations of the medium stacking fault energy type. Aging the material at 1300F has caused the precipitation of Ma& carbides on the dislocations. A similar effect can be seen in Fig. 8, where precipitation of Ma& carbide occurs on dislocations generated around primary M,C carbide particles during the quenching treatment.

In a 10~ stacking fault energy Co-14W-20Cr-3Fe-0.05C alloy, precip- itation of M,,Cs occurs primarily at grain boundaries, but no appreciable precipitation on dislocations is observed. In this case, the 10~ stacking fault energy causes substantial dissociation of dislocations (Fig. 5) and it appears

FIG. 7. Thin foil of HASTELLOY alloy X (nomina1 composition: Ni-ZOCr-SMo- 17Fe-O.OSC), solution heat-treated at 2150F for 1 hou, water-quenched, deformed 20,, at room temperature, and aged for 144 hours at 1300F. From Kotval.14

The ~lJicrostr.ucture of Superalloys 263

FIG. 8. Thin foil of HASTELLOY alloy X, solution heat-treated at 2150F for 1 hour, water-quenched. and aged for 500 hours at 1300F. M,,C, carbide precipitation on dis- locations generated around primary M,C carbide particles can be seen. From Kotval.

that the a/6 (112) Shockley partial dislocations do not act as preferential nucleation sites for the Ma&, carbide.

Precipitation on partial dislocations has, however, been demonstrated (A. Fourdeux, unpublished work) in the case of a Co-lOCr-1.3Be alloy, which also has a 10~ stacking fault energy. The alloy undergoes a f.c.c. - h.c.p. transformation in the matrix. The hexagonal regions are not perfect but heavily faulted. Figure 9 shows the aged structure of this alloy. In cubic regions of the matrix, beryllium-rich Guinier-Preston zones have formed, as is revealed by the streaking effects in the diffraction pattern shown in Fig. 10. At the end partials of the faults in the hexagonal region, however, precipitation of beryllide phase has occurred. Thus, there is a different nucleating role of the partial dislocations vis--vis the Ma&, carbide precipitate (in the Co-W-Cr-Fe alloy) and the beryllide precipitate (in the Co-Cr-Be alloy) in matrices of comparably low stacking fault energy (see Section 6). Discussion on the role of dislocations in nucleating y precipitates is postponed until later.

264 P. S. Kotval

FIG. 9. Thin foil of a Co-lOCr-1.3Be alloy. Aged for 96 hours at 932F. The cubic regions of the matrix exhibit beryllium-rich Guinier-Preston zones (see Fig. lO), whereas in the faulted hexagonal regions the partial dislocations have nucleated a beryllide precipitate. From Fourdeux (unpublished work).

FIG. 10. Diffraction pattern from a thin foil of Co-lOCr-1.3Be alloy indicating the presence of Guinier-Preston zones in the cubic regions of the matrix (see Fig. 9). From Fourdeux (unpublished werk).

The Microstructure of Superalloys 265

5. The Gamma-Prime Phase

The precipitation of y (a precipitate of the Ni,X type with Ll, crystal structure) is by far the most important strengthening mechanism in commercial nickel-base superalloys. In these compositions the classica1 NiaAl y frequently contains numerous dissolved elements, and complications in aging behavior are the result. In cobalt-base alloys, y precipitation has also been utilized, but this is a relatively recent development. Proper appreciation of the factors affecting y precipitation has, as in the case of other phenomena discussed previously, largely been a function of the techniques employed to study the precipitation reaction, and it is relevant to review the most recent results. Studies on binary and ternary alloys are very important here, as they serve to clarify the more complex situation in commercial superalloys. Factors to which attention has to be addressed in considering y precipitation are: (1) nucleation, (2) mismatch with the lattice and coherency, and (3) coarsening behavior of the y phase.

Manenc15,t6 and Williamsl used x-ray techniques and electron microscopy of replicas in studying the precipitation of y in binary Ni-Al alloys and con- cluded that a multistage process is involved. Manenc suggested a first stage of pre-precipitation, followed by precipitation of a coherent intermediate phase (on (100) planes) which was slightly tetragonal owing to coherency strains, followed by a visible stable precipitate possessing the equilibrium f.c.c. structure. The pre-precipitation stage has been interpreted variously by different authors. Bagariatskii and Tiapkinls concluded that the initial precipitation structure consists of isolated Guinier-Preston zones. Ben Israel and Finen have followed a suggestion of Cahn2 and postulated that y in NiiTi alloys is caused by spinodal decomposition. Working with a Ni-12Cr-37Fe-6Mo-3.22Ti0.26A1-0.05C alloy, Hammond and Ansell 21 have concluded that y particles have random crystallographic orientations and no coherency with the matrix. They suggest that y is heterogeneously nucleated on collapsed vacancy discs. However, this result is not supported by any other investigations, and evidente of the proposed collapsed vacancy discs was not found by Hammond and Ansellzl in their thin-foil studies.

Recent thin-foil studies by Ardell and Nicholsonz2 and by Phillipsz3 have helped to clarify much of the confusion of the previous results from x-ray diffraction and electron microscopy of replicas. Gamma-prime precipitation is frequently found to have a modulated structure both in the binary alloys of Ni-Al studied by Ardell and Nicholsonz2 and by Phillips23 and also in commercial alloys. Figure 11 shows an example of this modulated y structure in Udimet-700 alloy (nomina1 composition: Ni-15Cr-l5Coo5Mo4.3A1-3.5Ti-O.O5C). The dark-field micrograph in Fig. 11 was taken by using the superlattice reflections from the coherent precipitates; it is preferable to a bright-field image, since the coherency strains tend to obscure the particle-matrix interface in the latter case.

266 P. S. Kotval

FIG. ll. Dark-field thin-foil micrograph of Udimet-700 alloy showing modulated structure of the y precipitate. Courtesy of D. S. Acuncius.

Ardell and Nicholsonz2 have shown in a Ni-Al alloy that the type of modulated structure observed in Fig. 11 develops from an initially random precipitate distribution which gradually changes to one where cubic precipitates are aligned along (100) directions. Clearly, if the initial formation of y were due to spinodal decomposition, the modulation would be present from the very earliest stages of the aging sequence. Thus, for the smal1 supersaturations of solute in Ardell and Nicholsons alloy, spinodal decomposition is shown not to play a role in the nucleation of y. Phillips 23 has investigated a similar alloy and has again shown that spinodal decomposition is not responsible for the initial precipitation. He shows that precipitation of Ni-Al y is not a multistage process but is a normal nucleation and growth process. At high supersaturations coherent y precipitates by homogeneous nucleation and overages to a modulated structure. At lower supersaturations (higher aging temperatures) Phillips has shown that localized precipitation of y on dislocations occurs (Fig. 12). The morphology of these heterogeneously nucleated y precipitates is quite different from the homogeneously nucleated y, which again emphasizes the important effects that

The Microstructure qf Superalloys 267

268 P. S. Kotval

dislocations have on nucleation of precipitates. (See previous section.) Phillips23 has suggested that aluminum atoms segregate to the tension side of the edge dislocations as a Cottrell atmosphere, thus facilitating nucleation and growth in one direction. The growth of the precipitate edges parallel to (001) directions is implied by the fact that the particle edges (as at D) in Fig. 12a are at 45 to the direction of the dislocation-that is, (110). At a later stage in the aging, the individual platelets at a dislocation, such as EF in Fig. 126, apparently grow together, giving one big platelet that has thickened sufficiently to grow past the dislocation. Figure 12~ is a slightly tilted view of the same area as that shown in Fig. 12b, and the almost total absente of mass thickness contrast at EF indicates that the precipitate is quite thin. This is supported by the absente of fringes at the particle EF in both Figs. 12b and 12~. The close orientation of y precipitates with the matrix is also obvious, since no particle boundaries can be seen in Fig. 12~.

In commercial superalloys, the shape of individual y precipitates varies from perfect spheres to perfect tubes, although most particles have some inter- mediate shape. Hagel and Beattie 24 have suggested that spheres are formed when the lattice misfit between y and the matrix is less than -0.5 %, and tubes are formed in the range of 0.5 to 1.0%. When the disregistry is greater than

FIG. 13. Thin foil of Waspaloy, showing spherical y precipitates. Solution heat- treated for 1 hour at 1950F, water-quenched, and aged at 1292F for 1000 hours. From Kotval (unpublished work).

The ~UGrostructure of Superalloys 269

FIG. 14. Thin foil of Waspaloy showing cuboidal y precipital treated at 1950F for 1 hour, water-quenched, and aged at 1598F for Kotval (unpublished work).

zes. Solution heat- 1000 hours. From

FIG. 15. Same area as in Fig. 14, but with only one g operating sh no contrast perpendicular to g. From Kotval (unpublished work).

honing the lim of

270 P. S. Kotval

-1%, y precipitation is inhibited and only discontinuous precipitation occurs. However, Heslop2j has shown that in Nimonic alloy 115 (Ni-15Cr-I5Co-3.5Mo- 4Ti-5Al-O. 15C) cubical y phase occurs and the misfit is less than 0.176. This and other evidente discussed later indicate that the Hagel-Beattie suggestion is only partially valid. Recent thin-foil studies have shed considerable light on the coarsening behavior of y. In thin foils of Waspaloy (nomina1 composition: Ni-l9Cr-3.5Mo-2.8Ti-l.2Al-O.O96C), Kotval (unpublished work) has observed that particles that are spherical when smal1 (Fig. 13) often become cubical in shape with increasing size (Fig. 14). Kelly and Nicholsot? have suggested that this type of effect may be associated with the loss of coherency of the particle at an intermediate stage in the aging process. However, analysis of diffraction effects from thin-foil studies27 on a binary Ni-Al alloy indicates that such 10s~ of coherency is not necessarily implied when y particles cease to be spherical.

For smal1 spherical particles, indications of coherency can be obtained from a thin-foil micrograph taken under conditions such that only one reciprocal lattice reflection operates. With only one g (the reciprocal lattice vector) oper- ating, the Ashby-Brown 2s theory predicts that a line of no contrast wil1 occur perpendicular to g (Fig. 15) for smal1 coherent particles. This type of contrast is due to the strain field around the coherent particles.

For larger particles, the Ashby-Brown2s criterion no longer holds. Arde112 has shown that the strain contrast for larger particles takes the form of pseudo- fringes. The fact that the large particles are stil1 coherent is borne out by the existente of the fringes (Fig. 16).

Arde112 has shown that the width of the pseudo-fringe increases with in- creasing g. Further, irrespective of reflecting conditions, dislocation networks at the particle-matrix interfaces are not observed, which is consistent with the suggestion that the particles are coherent. Observations of this type employing thin-foil transmission electron microscopy thus provide a less confusing picture of coarsening behavior of y than that provided by observations from surface replication techniques. Most of the observations discussed above involved transmission techniques using lOO-kV electrons. Figure 16 shows the excellent resolution that can be obtained with higher accelerating voltages. This 500-kV micrograph (due to Ardell) illustrates the delta fringe contrast at the interfaces between large coherent y particles and the matrix in a Ni-6.71 wt. y aluminum alloy aged 450 hours at 1382F.

It was stated earlier that in a complex commercial superalloy the y phase frequently contains dissolved elements that lead to the substitution of solute atoms in both the Ni and Al portions of the ideal Ni,Al stoichiometric precipitate. The most common replacement for Al is Ti. The Al + Ti content of a commercial superalloy and the Al/T i ratio play important roles in governing the kinetics of y precipitation. Nordheim and Grant,2g Beattie and Hageh30

The Microstvucture of Superalloys 271

272 P. S. Kotval

and more recently Hughes31 have discussed the role of the Al/Ti ratio at length, and to avoid repetition, further discussion is not undertaken here. However, an example of the effect of substitution of a third solute element resulting in a Ni,(Al, Ti, X) type of phase is worth considering.

An example of the typical aged structure of a y-bearing superalloy (Udimet- 700) was shown earlier in Fig. 11. Udimet-700 alloy has a total Al + Ti content of approximately 7.8 wt. %, and hence the degree of supersaturation required to form Ni,(Al, Ti) y precipitate exists. By contrast, in an alloy such as Inconel- 718 (nomina1 composition: 53Ni-19Cr-19Fe-3Mo-5Cb-O.8Ti&0.6A1-0.05C) the Al + Ti content is -1.4 wt. 7. Earlier work32 on this alloy using replication techniques and x-ray diffraction of extracted residues had identified the fine strengthing phase as f.c.c.-ordered y with a composition of the type Ni,(Al, Ti, Cb). It was suggested 33 that the y phase was a metastable precipitate that overaged to the orthorhombic Ni,Cb phase found in this alloy. However, recent studies34,3S using transmission electron microscopy of thin foils of Inconel-718 have shown the strengthening phase to be body-centered-tetragonal in structure with a composition of the Ni,(Cb, Al, Ti) type. In a 55Nii15Cr- 25Fe-5Cb alloy, Kirman and Warrington 36 have shown that the Ni,Cb phase is body-centered-tetragonal and have proposed the structure shown in Fig. 17. Kotval has found that, in certain unique orientations, thin foils of Inconel-718 show unambiguously that the strengthening phase is body-centered-tetragonal.

0 Cb t C o Ni

I I 0 0 t 0

R b -- -.__ __ >

/' 0

Q

FIG. 17. Proposed structure of Ni,Cb. Xfter Kirman and Warrington.3U

The A!licrostructure of Superalloys 273

FIG. 18. Thin-foil electron micrograph of Inconel alloy 718 aged for 1118 hours at

1200F. From Kotval (unpublished werk).

I

0 . . . 0

0 .D .o ZO Schematic Reproduction

Of Diffraction Pattern FIG. 19. Selected area electron diffraction pattern of Inconel alloy 718 aged for 1118

hours at 1200F. FIG. 20. Schematic representation of the diffraction pattern in

Fig. 19.

274 P. S. Kotval

0 0 400m 442m

0 0 200m 242m

0 OOOm

0

042m

0 0 200m 242m

21 MATRIX: (012) f.c.c.

ai 004p

CD

I I

002p I

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0 . . . 0

0 . * . 0 Soep 211P 522p

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0 . . 0 2oop 211p 222p 233~ 244~

22 ZONE 1: (074) b.c.tet

8

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134P

8 ozop

ooop

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12ip

I Q

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0 253p

8

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a 004p

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0

23 ZONE 2: (210) b.c.tet 24 ZONE 3: (Tol) b.c.tet

I;I(:. 21. Matrix zone represented in thr difkaction pattern:(OT2) f.c.c. FICS. 22, 23, 24. Three body-centered-tetragonal zones, (OT4), (2TO), and (Tol), respectively, representing three mutually perpendicular orientations of the platelets of the strengthening phase. Figures 19-24 from Kotval.3

The Microstructure of Supwalloys 275

Figure 18 is a thin-foil micrograph of Inconel-718 showing the platelets of the strengthening phase.

A selected area diffraction pattern taken from a different area of the same specimen (aged for 1118 hours at 1200F) is shown in Fig. 19 and schematically reproduced in Fig. 20. Figure 21 shows the indexing of the f.c.c. matrix zone in the diffraction pattern as (012) f. C.C. The precipitate spots are composed of three body-centered-tetragonal zones, (O4), (2O), and (Tol) (Figs. 22, 23, and 24, respectively). The streaking of certain precipitate spots in (100) matrix directions is consistent with the c-axis of the precipitate being normal to the platelet of the body-centered-tetragonal phase.

The identification of the strengthening phase in Inconel-718 (with Al + Ti content approximately 1.4 wt. %) as a body-centered-tetragonal phase provides a good illustration of the effect of substitution of columbium for Al + Ti in the Ni,(Al, Ti) y precipitate. In alloys with high Al + Ti contents the phase retains the ordered f.c.c. structure of the type shown in Figs. 11, 13, and 14. In Inconel- 718, sufficient columbium is dissolved in the phase to alter its structure to body-centered-tetragonal.

It would appear that subtle changes in the structure of strengthening phases are often caused by variations in chemistry which are always possible in a complex commercial superalloy, and here again we find that such changes are beginning to be appreciated with the application of more direct metallographic techniques.

6. Precipitation in Association with Stacking Faults

An interesting aspect of precipitation of MC carbides in certain superalloys is the tendency to form broad bands of very fine precipitate, as studied by KotvaP4 in Inconel alloy 625 (nomina1 composition: Ni-20Cr-5Fe-8Mo-3.5Cb- O.OSC). This type of structure was observed only if the alloy was solution heat-treated at a very high temperature-for example, 2282F-quenched, and then tempered around 1200F. Figure 25 shows the type of structure observed in a plastic replica. Bands of the intragranular precipitate can be seen. If this structure were precipitation on dislocations, then the dislocation density would be anomalously high for a solution heat-treated and aged material. Thin-foil studies14 indicated that the structure shown in Fig. 25 was in fact fine precipitate of CbC associated with thin sheets of stacking faults (Fig. 26) that did not exist in the as-quenched alloy (compare Fig. 3). With increasing aging times and temperatures, the stacking faults with associated particles of carbide grew in length. Figure 27 shows the thin-foil structure after 144 hours at 1200F. On al1 the octahedral planes in the matrix, stacking faults with characteristic fringe contrast can be seen. In certain areas of the faults the fringe contrast has been

P. S. Kotval

FIG. 25. Plastic replica of Inconel alloy 625. Solution heat-treated for 1 hour 2282F, water-quenched, and aged for 1000 hours at 1200F. From Kotval.12

FIG. 26. Thin-foil electron micrograph of Inconel alloy 625. Solution heat-treated for 1 hour at 2282F and aged for 93 hours at 1200F. From Kotval.14

The Microstuucture of Superalloys 277

FIG. 27. Thin-foil electron micrograph of Inconel alloy 625. Solution heat-treated at 2282F, water-quenched, and aged for 144 hours at 1200F. From Kotval.14

eliminated and there is clear evidente of precipitation of discrete particles. Selected area diffraction showed the particles to be CbC; this fact was confirmed by dark-field micrographs taken with a CbC reflection.14 In both the case of Inconel-62S4 and the work of Honeycombe and his co-workers37-3s on austenitic stainless steels exhibiting this mode of MC carbide precipitation, the stacking faults were found to be extrinsic. A detailed model explaining this precipitation reaction has been proposed by Silcock and TunstalL4 Essentially, Silcock and Tunstall have concluded that the following steps are involved in stacking fault precipitation:

1. A dislocation $ [llO] lying, at least in part, in either the (111) or (1 li) plane rather than in its slip planes (lil) or (711) dissociates according to the reaction

g [llO] -$[lll] +&[llz]

2. A Frank partial 4 [ 11 l] dislocation created by the above dislocation reaction acts as nucleating site for the MC carbides. The metal atoms (columbium in the case of the Inconel-625) initially segregate to the expanded side of the Frank dislocation and eventually nucleate particles of the MC carbide.

3. As the particles grow, the misfit created can be accommodated by vacanties emitted by the Frank moving forward (by climb) between the particles. Even- tually, the particle wil1 be surrounded by interface dislocation and the bowed-out

3082-3

278 P. S. Kotval

part of the Frank once again straightens into a configuration where it is possible for further metal atoms to segregate and nucleate MC particles. As the Frank climbs, an extrinsic stacking fault is created behind it.

Clearly, the reaction proposed by Silcock and Tunstal140 in (1) is possible only in medium stacking fault energy matrices as defined earlier in the text. For example, in a 10~ stacking fault energy matrix (see Figs. 4 and 5) dissociation of lattice dislocations into two Shockley partials precludes the possibility of the dissociation given in (l), and no stacking fault precipitation of the type observed in Figs. 26 and 27 can occur. This mode of precipitation illustrates the close relationship between the nature of the dislocation substructure in a superalloy matrix and precipitation reactions that occur. These results also exemplify the increased understanding that can result from thin-foil studies of a precipitation reaction as compared to the knowledge obtained from surface replica observations.

It has been observed (Fig. 26) that regions adjacent to grain boundaries are devoid of any stacking fault precipitation; that is, near an effective vacancy sink, no stacking fault precipitation is observed. Froes et a1.3g have conducted exper- iments with varying quenching rates, and it has been shown that the nucleation of stacking fault precipitation is dependent upon the matrix vacancy concen- tration even though the subsequent growth in the manner proposed by Silcock and Tunstal140 is independent of this parameter. Thus, the observation of a zone free of stacking fault precipitation near grain boundaries can be explained on the basis of the difficulty in nucleating stacking fault precipitation in regions where the matrix vacancy concentration is low. The dislocations in regions adjacent to grain boundaries do not receive a sufficient supply of vacanties during the quench to permit them to achieve an orientation in which they can dissociate according to the equation given in (1) above. It can be argued that the zone free of stacking fault precipitation near grain boundaries is due to the depletion of solute in these regions (because of the formation of grain boundary carbides). However, this argument is disproved by the observation of zones free of stacking fault precipitation around primary MC carbide particles. Clearly, the solute concentration around such particles is sufficiently high (since the particles are dissolved during solution heat-treatment) for stacking fault precipitation, but again the matrix-primary-particle interface can act as a vacancy sink and the vacancy concentration necessary for the nucleation of stacking fault precipitate cannot be achieved.

7. Order-Disorder Reactions

As has already been discussed, y precipitation is an important strengthening mechanism in superalloys. The strengthening is particularly effective because

The Microstructure of Superalloys 279

the precipitates are ordered. Dislocations tend to shear ordered precipitates in pairs where the leading dislocation creates an antiphase boundary which is destroyed by the trailing dislocation.41 It has been demonstrated42 that the dislocations also bow in pairs and that this tends to raise the critical Orowan stress over that required to bow a single dislocation. Recently, for nickel-base superalloys, Copley and Keaf13 have discussed the effects of ordered precipitates

FIG. 28. Thin-foil micrograph of the fine scale structure of antiphase boundaries in Ni,Mo. (A) Bright-field image; (B) dark-field image produced by a double diffraction reflection; (C) dark-field image using a reflection belonging to one tetragonal lattice; and (D) dark-field image using a reflection belonging to the other tetragonal lattice. Courtesy of Ruedl et ~1.~

280 P. S. Kotval

and put forward a theory of precipitation strengthening based on thin-foil metallographic observations, and to dwell on this subject here would be repetitive. Electron microscopic studies of the antiphase boundary structure of AB and AB, alloys have been reviewed by Beeler.44

In certain nickel-base superalloys, such as HASTELLOY alloy B (nomina1 composition: Ni-26Mo-2.5Co-1.OCr-1.OMn-l.OSi4.O5C) and HASTELLOY alloy C (nomina1 composition: Ni-16Mo-l6Cr4.5Fe-2.5Co-O.lC), the high molybdenum content results in the formation of nickel-molybdenum inter- metallic compounds. The intermetallic ordered compound Ni,Mo has recently been studied by thin-foil transmission electron microscopy45 and by field ion microscopy.46 It is interesting to consider these data and the increased under- standing that results from these techniques. These results also extend the studies44 on AB and AB, alloys to AB, compounds.

The structure of the ordered Ni,Mo compound has been determined by Harker and by Guthrie and Stansbury48 using x-ray diffraction techniques. The unit cel1 is body-centered-tetragonal with a = 5.727 A, c = 3.566 8, and contains eight nickel and two molybdenum atoms.47,48 It can be considered as a tetragonally distorted face-centered-cubic structure. The c-axis of the body- centered-tetragonal unit cel1 is oriented parallel to the (001) directions of the face-centered-cubic matrix (compare the orientation of the body-centered- tetragonal strengthening phase in Inconel alloy 718) and, on ordering, it can adopt three different orientations vis--vis the face-centered-cubic matrix.

Ruedl, Delavignette, and Amelinckx4j have studied the nature of the ordering reaction in Ni,Mo. In the bright-field thin-foil micrograph shown in Fig. 28A, a system of antiphase boundaries is visible. In some regions the visibility of antiphase boundaries is better than in others. The diffraction pattern of such a region is shown in Fig. 29a. This pattern is in fact the superposition of two (001) patterns produced by the tetragonal superlattice. This is shown by the diagrams in Fig. 29. Certain very intense spots are common to tetragonal lattices as wel1 as to the diffraction pattern of the basic f.c.c. lattice. The spots of medium intensity belong to one or the other of the tetragonal patterns. The very weak spots are due to double diffraction; they fill in the gaps so as to produce a complete square pattern again. The orientation differente between the two tetragonal patterns is also the angle enclosed by the a-axis in regions that have antiparallel c-axes. Ruedl et al. 45 therefore conclude that the region responsible for the diffraction pattern must contain domains with the c-axis parallel and others with the c-axis antiparallel to the incident electron beam.

In the dark-field images of Fig. 28C and Fig. 280, superlattice reflections from the two tetragonal patterns were used; it is clear that complementary images are obtained for the background intensities, showing that, for the dark regions in Fig. 280, the c-axis is up, whereas it is down for the bright region. Figure 28B shows a dark-field image made with a double diffraction spot;

The Microstructuve of Supe-ralloys 281

. .

020 ??.

L

. . 310

200 mafrir (C)

fwin

282 P. S. Kotval

2 hours which completely transforms the structure to Ni,Mo. In Fig. 30, planar defects can be observed (see the bright and dark lines indicated by arrows). By comparison, in Fig. 31, a large number of boundaries (indicated by arrows) are visible.

These boundaries lie in between domains whose c-axes are nonparallel, and hence the (rotational) boundary has the appearance of a grain boundary in the field ion micrograph. These observations provide complementary data to the thin-foil observations (Fig. 28) on ordering in Ni,Mo.

Certainly, orderdisorder reactions in nickel-base matrices have received the least attention of al1 the precipitation reactions that occur and can be employed usefully in superalloys. The above results provide two examples of the increased understanding that can result from the applications of modern metallographic techniques.

FIG. 30. Field ion micrograph of a Ni-21% Mo alloy (14 at. y0 MO) quenched from 1562F. Courtesy of Newman and HremaG

The Microstructure of Superalloys 283

FIG. 31. Field ion micrograph of Ni,Mo fully ordered at 1382F. Negative was taken by contact with fiber optic window. Courtesy of Newman and Hren.4

8. Conclusions

In this paper an attempt has been made to deal with some aspects of the microstructure of superalloys and to present a few key points where experimental techniques are catching up with theory and hence lead to an increased appre- ciation of the phenomena by which superalloys are strengthened. The fact that dislocation-precipitate interactions can be studied in the electron microscope has provided a stimulus for studying superalloys on the basis of their stacking fault energy. This has led to a new dimension in studying the established strengthening mechanisms in superalloys and provides clues for the development of new mechanisms. Studies of this type wil1 ensure the possibility of synthesis of different strengthening mechanisms in a given matrix, once the precipitate- dislocation reactions are understood and controlled.

284 P. S. Kotval

Acknowledgments

1 should like to thank Mlle. Fourdeux for providing some of her unpublished data and those workers who provided some of the micrographs that are individually acknowledged in the text. 1 am grateful to many of my colleagues at Union Carbide for helpful discussions.

References

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Ohio, Report 217 (1965).

The Microstructure of Superalloys 285

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(1964) 1. 38. H. Harding and R. W. K. Honeycombe, J. Iran Steel Inst. (Londen), 204 (1966) 259. 39. F. H. Froes, R. W. K. Honeycombe, and D. H. Warrington, Acta Met., IS (1967) 157. 40. Jeanne Silcock, and W. T. Tunstall, Phil. Mag., 10 (1964) 361. 41. Paul A. Flinn, Trans. Met. Sm. AZME, 218 (1960) 145. 42. A. J. Ardell, Hardening of Alloys Containing Bi-Modal Distributions of Preci-

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ORNL-3078 (1961).

Accepted Septemberll, 1968