heat treatment. Increasing the aluminum/titanium ratio improves high-temperature properties. The volume fraction, size, and spacing of γ' are important parameters to control. Alloys with low amounts of γ' require greater attention to γ' spacing than alloys with high amounts. Other factors, such as coherency strain due to the lattice mismatch between γ and γ', appear to be important in certain alloys, such as Waspaloy. Grain size is an important microstructural parameter. Fine grain sizes normally provide superior room-temperature properties, such as toughness, strength, and fatigue resistance. Coarse grain sizes generally yield better creep resistance at elevated temperatures, although properties under other types of loading may suffer. Duplex grain structures generally are undesirable. Grain size also affects carbide precipitation at the grain boundaries. Coarse grain sizes have less grain- boundary surface area; therefore, carbide precipitation will be more continuous and thicker, thus impairing properties. Due to these problems, a uniform, intermediate grain size is generally preferred. Cobalt-Base Alloys. Cobalt-base superalloys are strengthened by solid-solution alloying and carbide precipitation. The grain-boundary carbides inhibit grain-boundary sliding. Unlike the iron-nickel- and nickel-base alloys, no intermetallic phase has been found that will strengthen cobalt-base alloys to the same degree that γ' or γ'' strengthens the other superalloys. Gamma is not stable at high temperatures in cobalt-base alloys. The carbides in cobalt-base superalloys are the same as those in the other systems. References cited in this section 41. C. Razim, Metallography of Heat Resistant and High Temperature Alloys, Pract. Metallog., Vol 5, May 1968, p 225-241; June 1968, p 299-309 42. G.P. Sabol and R. Stickler, Microstructure of Nickel-Base Superalloys, Phys. Stat. Sol., Vol 35 (No. 1), 1968, p 11-52 43. P.S. Kotval, The Microstructure of Superalloys, Metallography, Vol 1, 1969, p 251-285 44. F.R. Morral et al., Microstructure of Cobalt-Base High-Temperature Alloys, ASM Met. Eng. Q., Vol 9, May 1969, p 1-16 45. J.F. Radavich and W.H. Couts, Metallography of the Superalloys, Rev. High-Temp. Mater., Vol 1, Aug 1971, p 55-96 46. C.P. Sullivan and M.J. Donachie, Some Effects of Microstructure on the Mechanical Properties of Nickel- Base Superalloys, ASM Met. Eng. Q., Vol 7, Feb 1967, p 36-45 47. R.F. Decker, Strengthening Mechanisms in Nickel-Base Superalloys, in Symposium: Steel- Strengthening Mechanisms, 5-6 May 1969, Zurich, 1970, p 147-170 48. C.T. Sims and W.C. Hagel, Ed., The Superalloys, John Wiley & Sons, 1972 49. C.P. Sullivan et al., Relationship of Properties to Microstructure in Cobalt-Base Superalloys, ASM Met. Eng. Q., Vol 9, May 1969, p 16-29 50. C.P. Sullivan and M.J. Donachie, Microstructures and Mechanical Properties of Iron-Base (- Containing) Superalloys, ASM Met. Eng. Q., Vol 11, Nov 1971, p 1-11 51. D.R. Muzyka, Controlling Microstructure and Properties of Superalloys Via Use of Precipitated Phases, ASM Met. Eng. Q., Vol II, Nov 1971, p 12-20 52. D.R. Muzyka and G.N. Maniar, Microstructure Approach to Property Optimization in Wrought Superalloys, in STP 557, ASTM, Philadelphia, 1974, p 198-219 Phases in Wrought Heat-Resistant Alloys The microconstituents observed in iron-nickel and nickel-base wrought heat-resistant superalloys are identical, with a few exceptions. The cobalt-base alloys are not strengthened by precipitated intermetallics, but share many common features. All the alloys have an austenitic (γ phase) matrix that is strengthened by solid-solution alloying and by carbide precipitation. Most of the phases discussed below have some degree of solubility for other elements; therefore, their true compositions will vary from alloy to alloy and may be altered by heat treatment and thermal exposure. Not all phases permit substitution, however. Eta phase (Ni 3 Ti) has no significant solubility for other elements. Table 8 summarizes data on the commonly encountered constituents in these alloys. Table 8 Constituents observed in wrought heat-resistant alloys (a) Phase Crystal structure Lattice parameter, nm (b) Formula Comments γ' fcc (ordered L1 2 ) 0.3561 for pure Ni 3 Al to 0.3568 for Ni 3 (Al 0.5 Ti 0.5 ) Ni 3 Al Ni 3 (Al, Ti) Principal strengthening phase in many nickel- and nickel-iron- base superalloys; crystal lattice varies slightly in size (0 to 0.5%) from that of austenite matrix; shape varies from spherical to cubic; size varies with exposure time and temperature η hcp (D0 24 ) a 0 = 0.5093 c 0 = 0.8276 Ni 3 Ti (no solubility for other elements) Found in iron-, cobalt-, and nickel-base superalloys with high titanium/aluminum ratios after extended exposure; may form intergranularly in a cellular form or intragranularly as acicular platelets in a Widmanstätten pattern γ'' bct (ordered D0 22 ) a 0 = 0.3624 c 0 = 0.7406 Ni 3 Nb Principal strengthening phase in Inconel 718; γ'' precipitates are coherent disk-shaped particles that form on the {100} planes (avg diam approximately 600 A o , thickness approximately 50 to 90 A o ); metastable phase Ni 3 Nb (δ) orthorhombic (ordered Cu 3 Ti) a 0 = 0.5106-0.511 b 0 = 0.421-0.4251 c 0 = 0.452-0.4556 Ni 3 Nb Observed in overaged Inconel 718; has an acicular shape when formed between 815 and 980 °C (1500 and 1800 °F); forms by cellular reaction at low aging temperatures and by intragranular precipitation at high aging temperatures MC cubic a 0 = 0.430-0.470 TiC NbC HfC Titanium carbide has some solubility for nitrogen, zirconium, and molybdenum; composition is variable; appears as globular, irregularly shaped particles that are gray to lavender; "M" elements can be titanium, tantalum, niobium, hafnium, thorium, or zirconium M 23 C 6 fcc a 0 = 1.050-1.070 (varies with composition) Cr 23 C 6 (Cr, Fe, W, Mo) 23 C 6 Form of precipitation is important; it can precipitate as films, globules, platelets, lamellae, and cells; usually forms at grain boundaries; "M" element is usually chromium, but nickel-cobalt, iron, molybdenum, and tungsten can substitute M 6 C fcc a 0 = 1.085-1.175 Fe 3 Mo 3 C Fe 3 W 3 C- Fe 4 W 2 C Fe 3 NB 3 C Nb 3 Co 3 C Ta 3 Co 3 C Randomly distributed carbide; may appear pinkish; "M" elements are generally molybdenum or tungsten; there is some solubility for chromium, nickel-niobium, tantalum, and cobalt M 7 C 3 hexagonal a 0 = 1.398 c 0 = 0.4523 Cr 7 C 3 Generally observed as a blocky intergranular shape; observed only in alloys such as Nimonic 80A after exposure above 1000 °C (1830 °F), and in some cobalt-base alloys M 3 B 2 tetragonal a 0 = 0.560-0.620 c 0 = 0.300-0.330 Ta 3 B 2 V 3 B 2 Nb 3 B 2 (Mo, Ti, Cr, Ni, Fe) 3 B 2 Mo 2 FeB 2 Observed in iron-nickel- and nickel-base alloys with about 0.03% B or greater; borides appear similar to carbides, but are not attacked by preferential carbide etchants; "M" elements can be molybdenum, tantalum, niobium, nickel, iron, or vanadium MN cubic a 0 = 0.4240 TiN (Ti, Nb, Zr)N (Ti, Nb, Zr) (C, N) ZrN NbN Nitrides are observed in alloys containing titanium, niobium, or zirconium; they are insoluble at temperatures below the melting point; easily recognized as-polished, having square to rectangular shapes and ranging from yellow to orange μ rhombohedral a 0 = 0.475 c 0 = 2.577 Co 7 W 6 (Fe, Co) 7 (MO, W) 6 Generally observed in alloys with high levels of molybdenum or tungsten; appears as coarse, irregular Widmanstätten platelets; forms at high temperatures Laves hexagonal a 0 = 0.475-0.495 c 0 = 0.770-0.815 Fe 2 Nb Fe 2 Ti Fe 2 Mo Co 2 Ta Co 2 Ti Most common in iron-base and cobalt-base superalloys; usually appears as irregularly shaped globules, often elongated, or as platelets after extended high-temperature exposure σ tetragonal a 0 = 0.880-0.910 c 0 = 0.450-0.480 FeCr FeCrMo CrFeMoNi CrCo CrNiMo Most often observed in iron- and cobalt-base superalloys, less commonly in nickel-base alloys; appears as irregularly shaped globules, often elongated; forms after extended exposure between 540 and 980 °C (1005 to 1795 °F) (a) For more information on this subject, see the article "Crystal Structure of Metals" in this Volume. (b) 1 nm = 10 A o Gamma prime, a gcp phase, has an ordered fcc: L1 2 crystal structure and is Ni 3 Al or Ni 3 (Al,Ti), although considerable elemental substitution occurs. For example, cobalt and chromium will replace some of the nickel, and titanium will replace part of the aluminum. Iron can replace nickel or aluminum. The lattice parameters of γ and γ' are similar, resulting in coherency, which accounts for the value of γ' as the principal strengthening agent in iron-nickel- and nickel-base superalloys. Gamma prime is spherical in iron-nickel-base and in some of the older nickel-base alloys, such as Nimonic 80A and Waspaloy. In the more recently developed nickel-base alloys, γ' is generally cuboidal. Experiments have shown that variations in molybdenum content and in the aluminum/titanium ratio can change the morphology of γ. With increasing γ/γ' mismatch, the shape changes in the following order: spherical, globular, blocky, cuboidal (Ref 53). When the γ/γ' lattice mismatch is high, extended exposure above 700 °C (1290 °F) causes undesirable η(Ni 3 Ti) or δ(Ni 3 Nb) phases to form. The volume fraction, size, and distribution of γ' are important parameters for control of properties. The volume fraction of γ' increases with the addition of aluminum and titanium, but the amounts of each must be carefully controlled. Gamma prime contents above approximately 45% render the alloy difficult to deform by hot or cold working. In the iron-nickel- base alloys, the volume fraction of γ' phase is less than 0.2%, and it is usually spherical. Optimum strength results when the γ' is in the size range of 0.01 to 0.05 μm, much too small to be seen using the optical microscope. If the aluminum/titanium ratio is equal to or greater than 1, extended high-temperature exposure results in replacement of γ' by Ni 2 AlTi, NiAl, or Ni(Al,Ti). These phases overage rapidly at moderately high temperatures, forming massive platelike precipitates. Alloys with γ' contents below 20%, such as Nimonic 80A, are heat treated using a simple two-step process of solution annealing and aging. The solution anneal recrystallizes the austenite matrix and dissolves any γ' and M 23 C 6 carbides present. Aging precipitates γ' uniformly throughout the matrix and precipitates M 23 C 6 carbides at grain and twin boundaries. Alloys with γ' contents of approximately 30%, such as Waspaloy or Udimet 500, are solution treated, then given two aging treatments. Alloys with 40 to 45% γ', such as Udimet 700, are solution treated, then given three aging treatments. Positive identification of γ' is usually performed by x-ray diffraction of the residue of bulk extractions or by electron diffraction using extraction replicas. Some of the electrolytes that selectively attack γ' can be quite useful, because the γ' will be recessed below the matrix, and the other second phases will be in relief or plane with the surface depending on the preparation procedure. Gamma double prime has an ordered bct D0 22 crystal structure with an Ni 3 Nb composition and is found in iron- nickel-base alloys containing niobium. It gained prominence as the strengthening phase with the introduction of Inconel 718 (Ref 54). Early studies of the strengthening mechanism produced conflicting results until the precise details of γ''- phase formation, composition, crystallography, and stability were determined (Ref 55, 56, 57, 58, 59, 60). Gamma double prime has a disk-shaped morphology and precipitates with a well-defined relationship to the austenite matrix: [001]γ'' P <001>γ and {100}γ'' P {100}γ. Strengthening is due to the coherency strains produced by the low degree of γ/γ'' lattice mismatch. Although γ'' and γ' are present in Inconel 718 after aging, the amount of γ' is much less, and γ'' is the primary strengthening agent. Other alloys strengthened by γ'' include Inconel 706 and Udimet 630. Because γ'' is not a stable phase, application of alloys such as Inconel 718 is restricted to below 700 °C (1290 °F). Above this temperature, extended exposure produces a loss of strength due to rapid coarsening of γ'', solutioning of γ'' and γ', and formation of the stable orthorhombic form of Ni 3 Nb, which has an acicular, platelike shape. Positive identification of bct γ'' is more difficult than γ', because x-ray diffraction of bulk extraction residues will not detect γ''. The failure to detect bct γ'' is attributed to line broadening due to the very fine particle size that obscures the peaks of interest (Ref 57). Electron diffraction will, however, detect the superlattice lines of bct γ''. Bright-field TEM examination is unsatisfactory for resolving γ'' due to the high density of the precipitates and the strong contrast from the coherency strain field around the precipitates. However, dark-field TEM examination provides excellent imaging of the γ'' by selective imaging of precipitates that produce specific superlattice reflections (Ref 57). In addition, γ'' can be separated from γ' using the dark-field mode, because the γ'' dark-field image is substantially brighter than that of γ' (Ref 57). Eta phase has a hexagonal D0 24 crystal structure with a Ni 3 Ti composition. Eta can form in iron-nickel-, nickel-, and cobalt-base superalloys, especially in grades with high titanium/aluminum ratios that have had extended high-temperature exposure. Eta phase has no solubility for other elements and will grow more rapidly and form larger particles than γ', although it precipitates slowly. Coarse η can be observed using the optical microscope. Two forms of η may be encountered. The first develops at grain boundaries as a cellular constituent similar to pearlite, with alternate lamellae of γ and η the second, intragranularly as platelets with a Widmanstätten pattern (Ref 61, 62, 63). The cellular form is detrimental to notched stress-rupture strength and creep ductility, and the Widmanstätten pattern impairs stress-rupture strength but not ductility. Eta phase is relatively easy to identify due to its characteristic appearance. Most of the general-purpose reagents will reveal η, as will x-ray diffraction of bulk-extracted residues. Carbides, which are important constituents, are present in all the wrought heat-resistant superalloys. Four basic types are encountered: MC, M 23 C 6 , M 6 C, and M 7 C 3 (where M represents one or more metallic elements). Carbides in these alloys serve three principal functions. First, grain-boundary carbides, when properly formed, strengthen the grain boundary, prevent or retard grain-boundary sliding, and permit stress relaxation. Second, if fine carbides are precipitated in the matrix, strengthening results. This is important in cobalt-base alloys that cannot be strengthened by γ'. Third, carbides can tie up certain elements that would otherwise promote phase instability during service. Carbide precipitation in nickel-base alloys has a stronger tendency to form at grain boundaries than in iron-nickel- or cobalt-base alloys. Although grain- boundary carbides, depending on their morphology, can degrade properties, reducing carbon content to low levels substantially reduces creep life and ductility in nickel-base alloys. Aging of iron-nickel- and nickel-base superalloys causes M 23 C 6 to form at the grain boundaries. The optimum situation is a chain of discrete globular M 23 C 6 particles at the grain boundaries. This form benefits creep-rupture life. However, if the carbides precipitate as a continuous grain-boundary film, properties will be seriously degraded. It is not uncommon to observe zones around the grain boundaries that are devoid of γ'. Such precipitate-free zones can significantly influence stress-rupture life, depending on the width of the zones (Ref 64). In these alloys, MC-type carbide is most frequently titanium carbide; other types, such as niobium carbide, tantalum carbide, or hafnium carbide, are less common. Titanium carbide has some solubility for other elements, such as nitrogen, zirconium, and molybdenum. They are large, globular particles observable on the as-polished surface, particularly if some relief is introduced during final polishing. Metal carbides usually are irregular in shape or cubic. They can be preferentially colored by certain etchants. The most important carbide in superalloys is M 23 C 6 , because it forms at the grain boundaries during aging and, when properly formed, increases the strength of the grain boundaries to balance the matrix strength. Although chromium is the primary "M" element, other metallic elements, such as nickel, cobalt, iron, molybdenum, and tungsten, can substitute for it. The discrete globular form is the most desirable morphology; films, platelets, lamellae, and cells have also been observed. The M 6 C carbide is generally rich in molybdenum or tungsten, but other elements, such as chromium, nickel, or cobalt, may substitute for it to some degree. It is the most commonly observed carbide in the cobalt-base superalloys and in nickel-base alloys with high molybdenum and/or tungsten contents. In these alloys, M 6 C is often observed in the as-cast condition randomly distributed throughout the matrix. In wrought alloys, it will usually be dissolved during heating before hot working. It may precipitate at the grain boundaries in a blocky form or intragranularly in a Widmanstätten pattern and can be preferentially stained by certain etchants. Although M 7 C 3 is not widely observed in superalloys, it is present in some cobalt-base alloys and in Nimonic 80A, a nickel-chromium-titanium-aluminum superalloy, when heated above 1000 °C (1830 °F). Additions of such elements as cobalt, molybdenum, tungsten, or niobium to nickel-base alloys prevents formation of M 7 C 3 . Massive Cr 7 C 3 is formed in Nimonic 80A in the grain boundaries after heating to 1080 °C (1975 °F). Subsequent aging at 700 °C (1290 °F) to precipitate γ' impedes precipitation of M 23 C 6 due to the previously formed Cr 7 C 3 , which generally exhibits a blocky shape when present at grain boundaries. Borides. Boron is added in small amounts to many superalloys to improve stress-rupture and creep properties or to retard formation of η phase, which would impair creep strength. Boron retards formation of the cellular grain boundary form of η, but has no influence on the intragranular Widmanstätten η. Consequently, boron influences grain-boundary structures. Boron also reduces the solubility of carbon in austenite, which increases precipitation of finer-sized MC and M 23 C 6 carbides. If the boron addition is sufficiently high, detrimental borides will form. Borides are hard and brittle and precipitate at the grain boundaries. Borides are generally of M 3 B 2 composition with a tetragonal structure (Ref 65). Molybdenum, tantalum, niobium, nickel, iron, or vanadium can be "M" elements. The identification of borides in Udimet 700 has been documented (Ref 65). Laves phase, a tcp phase, has a MgZn 2 hexagonal crystal structure with a composition of the AB 2 type. Typical examples include Fe 2 Ti, Fe 2 Nb, and Fe 2 Mo, but a more general formula is (Fe,Cr,Mn,Si) 2 (Mo,Ti,Nb). They are most commonly observed in the iron-nickel-base alloys as coarse intergranular particles; intragranular precipitation may also occur. Silicon and niobium promote formation of Laves phase in Inconel 718. Excessive amounts will impair room- temperature tensile ductility; creep properties are not significantly affected. Laves phases have been observed in certain cobalt-base alloys and have been identified as Co 2 W, Co 2 Ti, or Co 2 Ta. Sigma phase is a tetragonal intermetallic tcp phase that forms with a wide range of compositions. Various morphologies may be encountered, some of which are quite detrimental to properties. However, the presence of in superalloys is not necessarily damaging to properties. Sigma in the form of platelets or as a grain-boundary film is detrimental, but globular intragranular precipitation can improve creep properties. Considerable effort has been devoted to determining how composition influences σ-phase formation, particularly in nickel-base superalloys. References 66, 67, 68, 69, and 70present examples of the many studies that have been conducted. This work has substantially influenced alloy development. Sigma can be preferentially attacked or stained by a number of reagents. However, because of the wide range of alloy compositions that may contain σ and the variable nature of its composition, positive identification by etching is not always possible. X-ray diffraction of bulk extraction residues is a more reliable technique. Etching procedures are best applied when they can be tested for response on specimens of the alloy known to contain σ phase. Mu phase is a rhombohedral (triagonal) intermetallic tcp phase with a W 6 Fe 7 structure (Ref 71). In general, it has little influence on properties. Mu precipitates as coarse, irregularly shaped platelets in a Widmanstätten pattern. A general formula for μm is (Fe,Co) 7 (Mo,W) 6 . Nickel can substitute for part of the iron or the cobalt. Nitrides are commonly observed in superalloys containing titanium or niobium as titanium nitride (most common) or niobium nitride. Nitrides are not influenced by heat treatment and are insoluble to the melting point. Nitrides are easily identified in the aspolished condition or after etching due to their regular, angular shapes and distinct yellow-to-orange color. Nitrides are quite hard and will appear in relief after polishing. They have some solubility for carbon and may be referred to as Ti(C,N),Nb(C,N), and so on. They should not be confused with carbonitrides, which are much richer in carbon and lower in nitrogen. Nitrides, often duplex, include an embedded phase or a surrounding film; this second phase is generally a darker colored nitride containing considerable carbon. The usual amounts present in superalloys generally have little influence on properties. Other Phases. A few other phases are less frequently observed in wrought heat-resistant alloys. For example, a few cobalt-base alloys have been developed that attain some degree of strengthening by precipitation of intermetallic phases, such as CoAl, Co 3 Mo, or Co 3 Ti. In alloys similar to A-286, G phase (Ni 18 Ti 8 Si 6 ) has been observed (Ref 72, 73). This phase has a globular shape and precipitates in grain boundaries. It is detrimental to stress-rupture life. A chromium-iron niobide, Z phase, has been observed in an Fe-18Cr-12Ni-1Nb alloy after creep testing at 850 °C (1560 °F) (Ref 74). Inclusions, some of which are similar to those found in steels, can also be found in these alloys. However, in the nickel- base alloys, titanium sulfides may be observed. Oxides, such as Al 2 O 3 or magnesia, may also be present. Oxides and sulfides may be observed at the surface of components due to environmental effects. Coatings are also used on same alloys, and their microstructures may be of interest (Ref 75, 76). References cited in this section 53. W.T. Loomis et al., The Influence of Molybdenum on the γ' Phase in Experimental Nickel- Base Superalloys, Met. Trans., Vol 3, April 1972, p 989-1000 54. J.F. Barker, A Superalloy for Medium Temperatures, Met. Prog., Vol 81, May 1962, p 72-76 55. I. Kirman and D.H. Warrington, Identification of the Strengthening Phase in Fe-Ni-Cr-Nb Alloys, J. Iron Steel Inst., Vol 205, Dec 1967, p 1264-1265 56. P.S. Kotval, Identification of the Strengthening Phase in "Inconel" Alloy 718, Trans. AIME, Vol 242, Aug 1968, p 1764-1765 57. D.F. Paulonis et al., Precipitation in Nickel-Base Alloy 718, Trans. ASM, Vol 62, 1969, p 611-622 58. W.J. Boesch and H.B. Canada, Precipitation Reactions and Stability of Ni 3 Cb in Inconel Alloy 718, J. Met., Vol 21, Oct 1969, p 34-38 59. I. Kirman and D.H. Warrington, The Precipitation of Ni 3 Nb Phases in a Ni-Fe-Cr-Nb Alloy, Met. Trans., Vol 1, Oct 1970, p 2667-2675 60. R. Cozar and A. Pineau, Morphology of γ' and γ'' Pre cipitates and Thermal Stability of Inconel 718 Type Alloys, Met. Trans., Vol 4, Jan 1973, p 47-59 61. J.R. Mihalisin and R.F. Decker, Phase Transformations in Nickel-Rich Nickel-Titanium-Aluminum- Alloys, Trans. AIME, Vol 218, June 1960, p 507-515 62. B.R . Clark and F.B. Pickering, Precipitation Effects in Austenitic Stainless Steels Containing Titanium and Aluminum Additions, J. Iron Steel Inst., Vol 205, Jan 1967, p 70-84 63. L.K. Singhal and J.W. Martin, Precipitation Processes in an Austenitic Stainle ss Steel Containing Titanium, J. Iron Steel Inst., Vol 205, Sept 1967, p 947-952 64. E.L. Raymond, Effect of Grain Boundary Denudation of Gamma Prime on Notch- Rupture Ductility of Inconel Nickel-Chromium Alloys X-750 and 718, Trans. AIME, Vol 239, Sept 1967, p 1415-1422 65. H.J. Beattie, The Crystal Structure of a M 3 B 2 -Type Double Boride, Acta Crystallogr., Vol 11, 1958, p 607- 609 66. L.R. Woodyatt et al., Prediction of Sigma- Type Phase Occurrence from Compositions in Austenitic Superalloys, Trans. AIME, Vol 236, 1966, p 519-527 67. E.O. Hall and S.H. Algie, The Sigma Phase, Met. Rev., Vol 11, 1966, p 61-88 68. J.R. Mihalisin et al., Sigma Its Occurrence, Effect, and Control in Nickel-Base Superalloys, Trans. AIME, Vol 242, Dec 1968, p 2399-2414 69. R.G. Barrows and J.B. Newkirk, A Modified System for Predicting σ Formation, Met. Trans., Vol 3, Nov 1972, p 2889-2893 70. E.S. Machlin and J. Shao, SIGMA- SAFE: A Phase Diagram Approach to the Sigma Phase Problem in Ni Base Superalloys, Met. Trans. A, Vol 9, April 1978, p 561-568 71. A. Raman, The μ Phases, Z. Metallkd., Vol 57, April 1966, p 301-305 72. H.J. Beattie and F.L. Ver Snyder, A New Complex Phase in a High-Temperature (Iron - Nickel - Chromium - Molybdenum) Alloy, Nature, Vol 178 (No. 4526), 1956, p 208-209 73. H.J. Beattie and W.C. Hagel, Intermetallic Compounds in Titanium-Hardened Alloys, Trans. AIME, Vol 209, July 1957, p 911-917 74. K.W. Andrews and H. Hughes, discussion of Aging Reactions in Certain Superalloys, Trans. ASM, Vol 49, 1957, p 999 75. G.F. Slattery, Microstructural Aspects of Aluminized Coatings on Nickel-Base Alloys, Met. Tech., Vol 10, Feb 1983, p 41-51 76. G.W. Goward et al., Formation and Degradation Mechanisms of Aluminide Coatings on Nickel- Base Superalloys, Trans. ASM, Vol 60, 1967, p 228-241 Atlas of Microstructures for Wrought Heat-Resistant Alloys Fig. 1 Alloy A-286 (AISI 660, 195 HV), solution annealed 2 h at 900 °C (1650 °F) and oil quenched. Spe cimen has a very fine austenite grain size. Glyceregia. 100× Fig. 2 Same alloy and processing as Fig. 1 , but showing an area near the surface of the specimen with a duplex grain structure. Tint etch: 20 mL HCl, 100 mL H 2 O, 2.4 g NH 4 F · HF, and 0.8 g K 2 S 2 O 5 . 100× Fig. 3 Same alloy and processing as Fig. 1, showing the very fine austenite matrix grains. Tint etch: 20 mL HCl, 100 mL H 2 O, 1 g NH 4 F · HF, 0.5 g K 2 S 2 O 5 . 200× Fig. 4 Alloy A- 286 (AISI 660, 357 HV), solution annealed 2 h at 900 °C (1650 °F), oil quenched, and held 16 h at 720 °C (1325 °F). A very fine-grained structure similar to that shown in Fig. 1. Glyceregia. 100× Fig. 5 Same alloy and processing as Fig. 4, but showing a region near the surface o f the specimen with a fine grain structure. Tint etched same as Fig. 2. 100× Fig. 6 Same alloy and processing as Fig. 4, showing the very fine matrix grain structure. Tint etched same as Fig. 2. 100× Fig. 7 Alloy A- 286 (AISI 660, 150 HV), solution annealed 1 h at 980 °C (1800 °F) and oil quenched, showing a coarser grain structure than in Fig. 1, 2, 3, 4, 5, and 6 due to the higher solutionizing temperature. Glyceregia. 100× Fig. 8 Same alloy and processing as Fig. 7, but tint etched using 20 mL HCl, 100 mL H 2 O, 1 g NH 4 F · HF , and 0.5 g K 2 S 2 O 5 . 100× Fig. 9 Alloy A-286 (AISI 660, 318 HV), solution annealed 1 h at 980 °C (1800 °F), oil quenched, aged 16 h at 720 °C (1325 °F), and air cooled. Glyceregia. 100× Fig. 10 Same alloy and processing as Fig. 9 , but tint etched. Only the matrix phase has been colored. 20 mL HCl, 100 mL H 2 O, 1 g NH 4 F · HF, and 0.5 g K 2 S 2 O 5 . 100× Fig. 11 Fig. 12 Fig. 13 A-286 (AISI 660), solution annealed 1 h at 980 °C (1800 °F) and aged 16 h at 720 °C (1325 °F), then air cooled and creep tested to rupture. Fig. 11: tested 7131 h at 650 °C (1200 °F). Matrix and grain- boundary precipitates have coalesced. 31.5 HRC. Fig. 12: tested 1232 h at 730 °C (1350 °F). Grain- boundary precipitates have coalesced; the matrix is darkened by γ' precipitation. HRC 25.5. Fig. 13: tested 546 h at 815 °C (1500 °F). Grain-boundary precipitates have coalesced, and the overaged matrix contains needlelike η phase (Ni 3 Ti). HRB 88.5. 15 mL HCl, 10 mL HNO 3 , and 10 mL acetic acid. 1000× [...]... 24 h at 70 5 °C (1300 °F) Grain-boundary M23C6 carbide is stabilized, and precipitation of fine γ' particles has increased Glyceregia 15,000× Fig 73 Fig 74 Fig 75 The effects of different etchants on solution-annealed and aged alloy X -7 5 0 Fig 73 : tint etched in 50 mL HCl, 50 mL H2O, and 1 g K2S2O5 Fig 74 : etched using Kalling's reagent 2 Fig 75 : etched using glyceregia See also Fig 76 , 77 , and 78 All... 76 , 77 , and 78 All 100× Fig 76 Fig 77 Fig 78 Different etchants used to delineate the structure of solution-annealed and aged alloy X -7 5 0 Fig 76 : etched using Marble's reagent Fig 77 : etched using aqua regia Fig 78 : etched using HCl + 1% Na2O2 All 100× Fig 79 Extraction replica electron micrograph of Nimonic 80, solution annealed 8 h at 1 075 °C (1965 °F) and aged 16 h at 70 5 °C (1300 °F) The structure... HCl and H2O 500× Fig 29 N-155 (AISI 661), solution annealed same as Fig 28, then aged 5 h at 76 0 °C (1400 °F) and air cooled Precipitated secondary carbide (M6C or M23C6) at grain boundaries and within grains 20% HCl, methanol, and 1% H2O2, 5 s 500× Fig 30 1 6-2 5-6 (AISI 650) alloy, after forging between 650 and 70 5 °C (1200 and 1300 °F) and stress relieving The solid-solution matrix exhibits banding... carbide As-polished 1500× Fig 94 U -7 0 0, solution annealed 4 h at 1 175 °C (2150 °F) and aged 4 h at 1080 °C (1 975 °F) M23C6 has dissolved, and borides have been spheroidized Large crystals are MC carbide; the oriented precipitate is γ' Kalling's reagent 1000× Fig 95 U -7 0 0, solution annealed 4 h at 1 175 °C (2150 °F), aged 4 h at 1080 °C (1 975 °F), aged 24 h at 845 °C (1550 °F), and aged 16 h at 76 0 °C (1400... recrystallized bands (light) and bands containing residual primary γ' Kalling's reagent 100× Fig 101 U -7 1 0 bar, solution annealed 4 h at 1 175 °C (2150 °F) and air cooled Structure is dispersed primary MC carbide and M3B2 boride in a γ matrix γ' is in solution Kalling's reagent 100× Fig 102 U -7 1 0 bar, solution annealed same as Fig 101, aged 24 h at 845 °C (1550 °F), air cooled, aged 16 h at 76 0 °C (1400 °F), and. .. Fig 15, 16, and 17 Fig 21: aged 1 h at 76 0 °C (1400 °F); 248 HV Fig 22: aged 8 h at 76 0 °C (1400 °F); 258 HV Fig 23: aged 128 h at 76 0 °C (1400 °F); 253 HV Glyceregia 500× (R.L Anderson) Fig 24 Incoloy 800 strip, in the mill-annealed condition The micro-structure consists of a solid-solution matrix in which some grains are delineated by precipitated carbide particles at the boundaries and by twinning... °F); 249 HV Fig 17: aged 512 h at 650 °C (1200 °F); 315 HV Glyceregia 500× (R.L Anderson) Fig 18 Fig 19 Fig 20 Same material and solution annealing treatment as Fig 15, 16, and 17 Fig 18: aged 2 h at 70 5 °C (1300 °F); 223 HV Fig 19: aged 64 h at 70 5 °C (1300 °F); 292 HV Fig 20: aged 272 h at 70 5 °C (1300 °F); 295 HV Glyceregia 500× (R.L Anderson) Fig 21 Fig 22 Fig 23 Same material and solution annealing... H3PO4, and H2CrO4 1000× Fig 58 Alloy 71 8, solution annealed 1 h at 955 °C ( 175 0 °F), air cooled, and aged 10 h at 76 0 °C (1400 °F) and at 650 °C (1200 °F) Structure is Laves phase (light gray particles), MC carbide (dark), and needlelike δ The matrix is γ phase See also Fig 60 Electrolytic: H2SO4, H3PO4, and H2CrO4 1000× Fig 59 Replica electron micrograph of same alloy and processing as Fig 57, showing... matrix grains HCl, ethanol, and H2O2 4500× Fig 99 Replica electron micrograph of U -7 0 0, solution annealed same as Fig 97 and aged 24 h at 980 °C (1800 °F) Precipitated carbide at grain boundaries and γ' within grains of the γ solid-solution matrix HCl, ethanol, CUCl2, and H2O2 4500× Fig 100 U -7 1 0 bar, solution annealed 2 h at 1120 °C (2050 °F), aged 1.5 h at 1040 °C (1900 °F), and air cooled Longitudinal... irregular) and γ' precipitates Electrolytic: H2SO4, H3PO4, and HNO3 10,000× Fig 69 Alloy 600 (202 HV), as-forged Specimen is from the center of a 305-mm (12-in.) diam bar Glyceregia 100× Fig 70 Alloy 625 (190 HV), solution annealed 30 min at 980 °C (1800 °F) and air cooled Specimen is a longitudinal section from the mid-radius of a 140-mm (5.5-in.) diam bar 15 mL HCl, 10 mL acetic acid, 5 mL HNO3, and 2 . Alloy 71 8, J. Met., Vol 21, Oct 1969, p 3 4-3 8 59. I. Kirman and D.H. Warrington, The Precipitation of Ni 3 Nb Phases in a Ni-Fe-Cr-Nb Alloy, Met. Trans., Vol 1, Oct 1 970 , p 266 7- 2 675 60. . 1 978 , p 56 1-5 68 71 . A. Raman, The μ Phases, Z. Metallkd., Vol 57, April 1966, p 30 1-3 05 72 . H.J. Beattie and F.L. Ver Snyder, A New Complex Phase in a High-Temperature (Iron - Nickel - . strengthening agent in iron-nickel- and nickel-base superalloys. Gamma prime is spherical in iron-nickel-base and in some of the older nickel-base alloys, such as Nimonic 80A and Waspaloy. In the