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A TEM study of decomposition behavior of a melt-quenched Al-Zr alloy
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~ E T A L L O G R A P H Y 17:231-252 (1984)
A TEM Study of Decomposition Behavior of a Melt-Quenched AI-Zr Alloy Z A R I F F A. C H A U D H U R Y AND C. S U R Y A N A R A Y A N A
CeHtre of Advanced Study in Metallur~,,y. Department of Metallurgical En~,im,erin~,,, Banaras Hindu University. VARANASI--221 005, India
Melt quenching of an A I - Z r alloy containing 9.4 wt.% Zr resulted in the formation of a supersaturated solid solution. Elevated temperature annealing treatments led to the formation of a metastable AI~Zr phase having an ordered cubic Cu~Au-type structure with a = 0.4073 nm and at a later stage the equilibrium A h Z r phase having a tetragonal structure with a 0.4013 nm and c 1.7321 nm. Electron microscopic investigations on the morphology and structural features of the as-quenched as well as transfc)rmation products formed during decomposition are discussed in detail.
Introduction In recent years melt quenching techniques have gained wide popularity not only because of their ability to produce extremely high cooling rates but because of the unusual properties associated with their products. Melt quenching could lead to the formation of supersaturated solid solutions at composition ranges far beyond those obtained by conventional methods. In a variety of alloy systems, new non-equilibrium crystalline intermediate phases and metallic glasses have also been produced. The innumerable results obtained on several alloys by the application of meltquenching techniques have been well reviewed [I-5]. A large number of investigations have been carried out on the morphology of the phases and the precipitation behavior in AI-Zr alloys. Transmission electron microscopy had been extensively employed lbr this purpose [6-11]. The majority of these investigations were concerned with chill-cast A1-Zr alloys containing up to 2 wt.% Zr. The effect of cooling rate on the grain size of alloys and grain refinement behavior were also studied, Formation of a metastable-ordered cubic AI3Zr phase has been found to be responsible for the improved properties of AI-Zr alloys. Melt quenching of AI-Zr alloys has also yielded some interesting resuits. Varich el al. [12] studied the effect of superheating the melt on the c© Elsevier Science Publishing Co., Inc., 1984 52 Vanderbilt Ave., New York, NY 10017
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structure and properties of rapidly quenched A1-Zr alloys containing 2.2 wt.% Zr. Varich and Sheyko [13] studied the stability and thermal expansion of A1-Zr alloys. Gudzenko and Polesya [14] observed the formation of an amorphous phase in alloys containing 45 to 53 at.% Zr. Recently, it has been shown that the solid solubility of Zr in Al could be enhanced by melt quenching up to about 8.6 wt.% Zr [15]. The maximum solid solubility of zirconium in aluminum under equilibrium conditions is reported to be 0.28 wt.% [16]. The most Al-rich intermediate phases in the AI-Zr system are AI3Zr, AIzZr, A13Zr2 and AIZr [17-19]. As part of an on-going program on the formation and decomposition behavior of metastable phases in aluminum alloys [20-26], we had undertaken detailed investigations on the A1-Zr [27] system. Since not much work was reported on the metastable behavior of rapidly quenched AI-Zr alloys, it was decided to investigate the impact of rapid quenching from the melt on the constitution and decomposition behavior of metastable phases in AI-Zr alloys.
Experimental Procedure An AI-Zr alloy with a nominal composition of 9.4 wt.% Zr was prepared in an arc-melting furnace under a protective argon atmosphere. The alloy was remelted to achieve homogenization. A small quantity of the alloy was induction-melted and rapidly quenched from the melt using the conventional "gun" technique, wherein the cooling rate achieved was estimated to be about 107 K/sec. The as-quenched foils were thin enough in many areas to be examined directly in a transmission electron microscope fitted with a goniometer stage and operating at 100 kV. The as-quenched foils were given proper annealing treatments after encapsulation under vacuum in quartz capsules. In-situ hot stage electron microscopy was also carried out using the hot-stage specimen holder to follow the transformation sequences.
Results AS-QUENCHED STRUCTURE Remarkable extension of solid solubility of Zr in Al has been observed on melt quenching. Figure 1 shows a bright-field electron micrograph featuring elongated grains of the supersaturated solid solution. The corresponding diffraction pattern is shown in an inset. The absence of any second phase in the microstructure and clear single crystal fcc reflections
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FrG. 1. Microstructure of the as-quenched AI-9.4 wt.% Zr alloy showing the formation of supersaturated solid solution. The corresponding diffraction pattern is shown as an inset. in the diffraction pattern suggest that all the 9.4 wt.% Zr had dissolved in AI to form a homogeneous supersaturated solid solution. These observations clearly indicate that the solid solubility of zirconium in aluminium in the melt-quenched condition can be increased to at least 9.4 wt.%.
Defect Structure As reviewed by Jones [28], rapidly-quenched metals and alloys show a variety of defect structures. This has been confirmed recently in our laboratory as well [24, 25, 29]. Though the solid solution grains formed in the present investigation were defect-free many times, interesting dislocation structures were sometimes observed in some of the grains. As a typical example, isolated dislocations could be noticed near the grain boundary in Fig. 2a. It was also possible, occasionally, to observe the dislocations arranged into regular low-angle boundaries as shown in Fig. 2b. The cell structure, surrounded by the low-angle boundaries, was assumed to have formed as a result of the recovery of stresses introduced during the quenching process, a situation similar to what happens in plast-
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(a)
(b) Flo. 2. Typical dislocations observed in the as-quenched microstructure: (a) isolated dislocations, and (b) low-angle boundaries.
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ically deformed and annealed metals. The dislocations might have been formed because of the plastic deformation induced by rapid quenching stresses and/or aggregation of vacancies. ANNEALING BEHAVIOR With a view to evaluating the stability of the supersaturated solid solution formed in this system, the as-quenched foils were heat treated both externally and in-situ. Salient features of these observations are here detailed.
External Heat Treatments The metastable solid solution was quite stable up to at least about 573°K. On annealing the foils at 593°K for 3.6 ks, the discontinuous precipitation of a second phase along the grain boundaries (Fig. 3) could clearly be seen. Simultaneously, it was noticed that the matrix also undergoes some transformation as revealed by the formation of extremely fine rodlike precipitates. Although it was not easy to distinguish between these
Fro. 3. Microstructure showing the discontinuous precipitation along the grain boundary on low-temperature annealing. Also notice the fine structure inside the matrix grains.
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two types of precipitates from electron diffraction patterns, it should be emphasized that diffraction patterns from this area could be indexed on the basis of an ordered cubic structure. Better contrast from the precipitates could be obtained when the foils were annealed at higher temperatures on the order of 773°K for 3.6 ks. Figures 4a to 4c show the bright-field and dark-field micrographs and the corresponding diffraction pattern. This diffraction pattern also could be indexed as arising from an ordered fcc structure. F r o m the available literature, it can be inferred that this is the metastable AI3Zr compound. A point of interest in this micrograph is the shape of the precipitate particles, which is nearly circular in cross section. This observation, coupled with the microstructure obtained on annealing at lower temperatures (Fig. 3), clearly suggests that metastable A13Zr forms as rods. Figure 5a shows that the matrix contains, in some areas, precipitates of metastable AI3Zr which are interconnected along the long axis of the rods. This gives the impression of long arrays of precipitates arranged parallel to each other. It may also be noticed in the micrograph that precipitation of yet another phase is just starting. The size of these particles is much larger than that of metastable A13Zr. When the alloy foils were
(a)
FIc. 4. (a) Precipitate particles uniformly distributed in the aluminum matrix after annealing the foils at 773°K for 3.6 ks; (b) the corresponding dark-field micrograph showing the shape of the precipitates; (c) [ 111] diffraction pattern taken from these particles indicating that this phase is the metastable cubic AIaZr.
Decomposition Behavior of Al-Zr Alloy
(b)
(c) Fl(3. 4.
(Continued)
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238
ipm (a)
d
(b) FIG. 5. (a) Precipitation of the rodlike precipitates after annealing at 773°K for 3.6 ks. Precipitation of another phase seems to just be starting; (b) grain boundary precipitation of the tetragonal equilibrium intermediate phase.
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(a)
(b) FIG. 6. (a) Microstructure showing lhe fan-shape of the precipitates of the melaslable AI~Zr phase; (b) the corresponding dark-field micrograph.
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Fro. 7. is [001].
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Diffraction pattern taken from the metastable rodlike precipitates. Foil normal
FIG. 8. A high magnification micrograph showing the shape and size of the rodlike precipitates after annealing at 773°K for 10.8 ks.
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annealed for 10.8 ks at 793°K, much more precipitation of this new phase along the boundaries of the elongated grains could be observed (Fig. 5b). While the rodlike precipitates have a diameter of about 15 nm and are spaced about 40 nm from each other, the grain boundary precipitates are much coarser, measuring about 80 nm at this stage. Diffraction patterns from this phase indicated a tetragonal crystal structure and a correspondence to the equilibrium AI3Zr phase. Figures 6a and 6b show a pair of bright-field and dark-field electron micrographs bringing out the fan-shape morphology of the metastable AI3Zr phase. A typical diffraction pattern from these precipitates is shown in Fig. 7, which unambiguously confirms that this is the At3Zr-ordered cubic phase with a = 0.4073 nm. A high magnification micrograph showing the morphology of the precipitate phase is presented in Fig. 8. The growth, thickening and branching off of the rods, is worth observing in this micrograph. Figure 9a shows the micrograph recorded from an alloy annealed for 10.8 ks at 773°K. Large-size precipitates of the tetragonal equilibrium
(a) FIG. 9. (a) Bright-field micrograph showing the formation of the equilibrium precipitates. The particles are uniformly distributed; (b) electron micrograph showing the precipitates inside the equi-axed grains; (c) a high magnification bright-field electron micrograph showing the shape of the precipitates within the grains. Note also the absence of precipitation near the grain boundaries.
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(b)
(c) FIG. 9.
(Continued)
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AI3Zr phase can be seen clearly here. Rodlike precipitates can also be seen in the background. However, even for the same heat treatment conditions, a different type of microstructure is observed in some other regions of the foil, as shown in Fig. 9b. Precipitation of a second phase inside the equiaxed grains is worth noticing. Another point of interest is the absence of precipitation near the grain boundaries, although one cannot compare this with the precipitate free zones observed in solid-state quenched and aged precipitation hardening alloys. A high-magnification micrograph of the central grain of Fig. 9b is shown in Fig. 9c, which indicates that the precipitate here also perhaps has a cylindrical shape. It has not been possible to get diffraction patterns from these precipitates, and so nothing can be said about the composition and crystal structure of this phase, except that, from the shape of the precipitate particles, it may be assumed that these belong to the metastable AI3Zr compound. Long annealing treatments at 903°K resulted in the complete disappearance of the metastable AI3Zr phase; the fact that the microstructure at this stage consists of only aluminum solid solution and the tetragonal AI3Zr phase suggests that equilibrium has been established. A typical micrograph corresponding to this situation is shown in Fig. 10. Two rep-
Fl(;. 10. Bright-field electron micrograph showing the distribution of the equilibrium AI3Zr phase along the equi-axed grain boundaries of the aluminum matrix after annealing the foils at 903°K.
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(a)
(b) FIG. ll. Typical single crystal electron diffraction patterns taken from the tetragonal equilibrium Al3Zr phases. Foil normals are (a) [591] and (b) [lll].
Decomposition Behavior of AI-Zr Alloy
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5pro (a)
IIIII
(b) Fie. 12. (a) Bright-field electron micrograph of a big equilibrium AI~Zr particle; (hi the corresponding dark-field micrograph.
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(a)
(b) FIG. 13. (a) Discontinuous precipitation of the metastable cubic AI3Zr phase; (b)-(d) a sequence of micrographs showing the growth of the discontinuous phase on continued annealing.
Decomposition Behavior of Al-Zr Alloy
(c)
(d)
FIG. 13.
(Contintted)
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resentative diffraction patterns recorded from these precipitate particles are shown in Figs. 1 la and 1lb. From these and several other patterns, it had been possible to confirm that this phase is the equilibrium AI3Zr compound having a tetragonal structure with the lattice parameters a = 0.4013 nm and c = 1.7321 nm. Figure 12 shows a pair of bright-field and dark-field micrographs featuring a very large AI3Zr particle indicating that further annealing leads only to coarsening of the precipitate particles.
Hot-Stage Electron Microscopy The as-quenched AI-9.4 wt.% Zr alloy foils have also been heat treated in situ using the hot-stage specimen holder. During continuous heating, it was observed that the supersaturated solid solution was quite stable up to about 593°K, and that precipitation took place beyond this point. On continued heating, discontinuous precipitation of metastable AI3Zr could be noticed at about 668°K (Fig. 13a). Figures 13b to 13d show a series of three micrographs recorded at 750°, 800 ° and 825°K depicting the growth of the discontinuous phase which gradually enveloped the whole of the untransformed matrix. Diffraction patterns from these areas clearly revealed that the phase is metastable AI3Zr, confirming the observation made during external heat treatments. At still higher temperatures the equilibrium tetragonal AI3Zr phase could be detected. Discussion
S U P E R S A T U R A T E D SOLID SOLUTION Earlier investigators [9, 12, 15, 30-32] also reported that the solid solubility of Zr in AI could be enhanced by rapidly quenching the melts of these alloys. Our results indicate that a maximum of 9.4 wt.% Zr could be dissolved in a metastable condition in the solid state. Earlier investigators reported a much lower value. This observation suggests that the extended value is about 35 times the maximum equilibrium value of 0.28 wt.% Zr. Work with more concentrated alloys will indicate whether still higher supersaturation can be achieved. While the previous investigations were carried out using either chill casting or a two-piston device, the present results were obtained by "gun" quenching, wherein the cooling rates obtained are definitely higher than those in any other technique. Thus, it was only natural that the solid solubility extension was more prominent in our studies. This is essentially due to the flow of melt into thin foils. The elongated grains noticed in the
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foils is further evidence that the molten droplets did flow on the substrate. However, in the relatively thicker portions of the foils the grains were equiaxed. Even here precipitation of a second phase could not be detected in the as-quenched condition. DECOMPOSITION BEHAVIOR The decomposition behavior of aluminum-rich AI-Zr solid solutions produced by solid-state quenching [6, 7, 33] and melt quenching [15, 32] has been studied earlier. Sahin and Jones [15] identified three distinct stages during isochronal annealing of an AI-4.9 wt.% Zr alloy: 1. Homogeneous modular separation, 2. Discontinuous precipitation of metstable cubic AI3Zr, and 3. Transformation to coarse plates of equilibrium tetragonal AI~Zr. In their studies on chill-cast alloys, Ohashi and Ichikawa [9] and Hori et al. [32] detected a petal-like structure in low-Zr alloys. This is not, however, confirmed by Sahin and Jones [15]. The supersaturated solid solution was quite stable up to a temperature of 593°K. Beyond this point precipitation started. A point of interest here is the discontinuous precipitation of metastable cubic AI3Zr phase along with a fine rodlike precipitate in the matrix. This fine structure in the matrix cannot be ascribed to the clustering effect observed in solid-state quenched and aged alloys, although Fontaine et al. [34] and Bonefa~i~ et al. [35] observed clustering effects due to the segregation of solute atoms in rapidly quenched A1-Fe and AI-Ni systems, respectively. The diffraction patterns from the A1-Zr specimens subjected to low-temperature annealing, however, do not show any evidence of clustering. This microstructure has been interpreted as consisting of a fine rodlike precipitate. The diffraction patterns clearly indicate that this is the cubic metastable AI3Zr phase. Depending on the orientation in which these rodlike precipitate particles are observed, either rods or round particles are noticeable. Evidence has also been collected to show that the fan-shaped precipitate observed by others is nothing but the rods arranged one next to the other. Figures 6 and 8 bring out these features very clearly. Hotstage experiments indicated that the interface velocity is very high and that the fan-shaped precipitates grow very rapidly into the matrix (Fig. 13). Ryum [6] has reported the formation of fine spherical precipitates. especially between the arms of the fan-shaped precipitates. These may be the rods viewed along the long direction. The petal-shaped precipitates have not been observed in our present investigation.
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Annealing at higher temperatures resulted in the precipitation of the equilibrium tetragonal A13Zr phase. All earlier investigators reported that the equilibrium A13Zr phase has a morphology of coarse plates which could be considered to have transformed from the metastable A13Zr phase. However, in our present investigation, we have noticed that this has a bulky shape and that, given a suitable annealing treatment, these can grow to very large sizes. The diffraction patterns could be unambiguously identified. From the foregoing it appears that, although one can clearly identify two stages in the transformation of the supersaturated solid solution to the equilibrium phases, these two stages seem to overlap quite extensively. Simultaneous precipitation of both the metastable and stable A13Zr phases was detected at as low a temperature as 793°K, while at high temperatures of about 900°K only the equilibrium phases could be observed. Conclusions On the basis of the results presented above the following conclusions can be drawn. 1. The solid solubility of Zr in A1 can be extended by melt quenching at least up to 9.4 wt.%. This is the highest value of supersaturation reported so far in this system. 2. Both in situ and external heat treatments suggest that the supersaturated solid solution is stable up to about 593°K. 3. Annealing at temperatures in excess of 600°K results in the precipitation of metastable ordered AI3Zr phase in the form of rods. Fan shape of these precipitates was also noticed. 4. Higher temperature annealing leads to the formation of the equilibrium tetragonai AI3Zr phase.
The authors are grateful to Professor T. R. Anantharaman, Program Coordinator, Center o f Advanced Study, for encouragement, and to the Head of the Department o f Metallurgical Engineering, Banaras Hindu University, for provision of laboratory facilities. They are also thankful to Dr. Shigenori Hori of Osaka University, Japan for supplying the A l Zr alloy used in this investigation, and to Dr. G. V. S. Sastry for a critical reading of the manuscript. References 1. T. R. Anantharamanand C. Suryanarayana,Review: A decade of quenchingfrom the melt, J. Mater. Sei. 6:1111-1135 (1971).
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2. H. Jones and C. Suryanarayana, Rapid quenching from the melt: An annotated bibliography 1958-72, J. Mater. Sci. 8:705-753 (1973). 3. H. Jones, Splat cooling and metastable phases, Rep. Progr. Phys. 36:1425-1497 (1973). 4. H. Jones, Rapid Solid(l~cation ~/' Metals and Alloys, Monograph No. 8, Institution of Metallurgists, London 11982). 5. C. Suryanarayana, Rapidly quenched metals--a bibliography 1973-1979. IFl/Plenum. New York (1980). 6. N. Ryum, Precipitation and recrystallization in an AI-0.5 wt.C/c Zr alloy, Acta Mct. 17:269-278 (1969). 7. O. lzumi and D. Oelschlagel, Strukturelle Untersuchung der ausscheidung in einer aluminium-legierung mit I. 1 Gew.-'~ Zr, Z. Metallkde. 60:845-851 (1969). 8. V. I. Dobatkin, V. I. Elagin, V. M. Fedorov and R. M. Sizova, Decomposition of saturated solid solutions in granulated aluminium alloys, Russian Met. 2:122-127 (1970). 9. T. Ohashi and R. Ichikawa, Grain refinement in aluminium-zirconium and aluminiumtitanium alloys by metastable phases, Z. Metallkde. 64:517-521 (1973). 10. E. Nes and H. Billdal, Non-equilibrium solidification of hyperperitectic AI-Zr alloys, Acta Met. 25:1031-1037 (1977); The mechanism of discontinuous precipitation of the metastable AI~Zr phase from an AI-Zr solid solution, Acta Met. 25:1039-1(146 (1977). 1I. S. Hori, H. Kitagawa, T. Masutani and A. Takehara, Structure and phase decomposition of rapidly solidified aluminium-zirconium solid solutions. J. Japan lttst. Li~,,ht Metal~ 27:129-137 (1977). 12. N. 1. Varich, R. B. Lyukevich, L, E. Kolomoytseva, A. N. Varich and V. V. Maslov. Effect of superheating the melt on the structure and properties of rapidly cooled AI-Zr alloys, Phys. Met. Metallogr. 27(2): 176-179 (1979). 13. N. I. Varich and T. I. Sheyko, Thermal expansion of Al-Mo, AI-Zr alloys prepared at high cooling rates, Phys. Met. Metallogr. 30(2):231-232 (1970). 14. V. N. Gudzenko and A. F. Polesya, Structure of zirconium-aluminium alloys rapidly cooled from the liquid state, Phys. Met. Metallogr. 39(6): 177-179 (1975). 15. E. Sahin and H. Jones, Extended solid solubility, grain refinement and age-hardening in AI-1 to 13 wt.% Zr rapidly quenched from the melt, in Rapidly Quenched Metah' II1, (B. Cantor, Ed.), The Metals Society. London (1978). pp. 138-146. 16. L, E. Mondolfo, Alaminiam Alloys--Structure and Propertie.~, Butterworths, London (1976). 17. M. Hansen and K. Anderko, Constitution (~['Bipm~3' Alloy.s, McGraw-Hill, New York (1958). 18. R. P. Elliott, Constitution ~f'Binao' Alloys: Fir,st Supplement. McGraw-Hill, New York (1965). 19. F. A. Shunk, Constitution ~fBinao' Alloys: Second Supplement, McGraw-Hill, New York (1969). 20. Z. A. Chaudhury, G. V. S. Saslry and C. Suryanarayana. Phase transformations in rapidly quenched aluminium-ruthenium alloys, Z. Metallkde. 73:201-206 (1982). 21. Z. A. Chaudhury and C. Suryanarayana, Metastable phases in vapour-deposited AIRu alloys, J. Mater. Sci. 17:3158-3164 (1982). 22. Z. A. Chaudhury and C. Suryanarayana, Annealing behaviour of vapour-deposited AIRh thin films, Thin Solid Films 98:233--239 (1982). 23. Z. A. Chaudhury and C. Suryanarayana, Al~3X4-type phases in aluminium-group VIII metal systems, J. Less-Common Metals 91:181-187 (1983). 24. Z. A. Chaudhury and C. Suryanarayana. Electron microscopic investigations on a meltquenched AI-Rh alloy, J. Mater. Sci. 18:3011-31)22 (1983). 25. G. V. S. Sastry and C. Suryanarayana, Metastable effects in melt-quenched AI-Pd alloys, Mater. Sci. & Eng. 47:193-208 {1981).
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26. G. V. S. Sastry, C. Suryanarayana and G. Van Tendeloo, A structural study of vapourdeposited A1-Pd alloys, Phys. Stat. Sol. (a) 73:267-278 (1982). 27. Z. A. Chaudhury, Structure of rapidly quenched aluminium alloys, Ph.D. thesis, Banaras Hindu University, Varanasi, India, 1983. 28. H. Jones, Lattice defects in splat-quenched metals, in Vacancies '76, (R. E. Smallman and J. E. Harris, Eds.), The Metals Society, London (1977), pp. 175-184. 29. L. R. K. Rao, C. Suryanarayana and S. Lele, Lattice defects in melt-quenched aluminium and its alloys, Aluminium 58:214-217 (1982). 30. L. M. Burov and A. A. Yakunin, Effect of the rate of cooling on the composition of solid solutions in binary alloys based on Al, Russian J. Phys. Chem. 42:540-541 (1968). 31. A. K. Kushnereva and I. V. Salli, The metastable crystallization of binary alloys, Metallofizika Desp. Mezhved. 56:115-120 (1968). 32. S. Hori, S. Saji and A. Takehara, Structure of rapidly solidified A1-Zr alloys and its thermal stability, in Proc. Fourth International Conf. on Rapidly Quenched Metals, (T. Masumoto and K. Suzuki, Eds.) The Japan Institute of Metals, Sendai (1982), Vol. II, pp. 1545-1548. 33. E. Nes, Precipitation of the metastable cubic A13Zr phase in subperitectic AI-Zr alloys, Acta Met. 26:499-506 (1972). 34. A. Fontaine, J. Dixmier and A. Guinier, Mechanical properties and structural transitions in splat-cooled AI-Mn and AI-Fe alloys, Fizika, 2, Suppl. 2:23.1-23.6 (1970). 35. A. Bonefa~i6, M. Kerenovi6, A. Kirin and D. Kunstelj, Segregation of solutes in AlNi and AI-Sn and its influence on their mechanical properties, J. Mater. Sci. 10:243251 (1975).
Received October 1983; accepted January 1984.
~ E T A L L O G R A P H Y 17:231-252 (1984)
A TEM Study of Decomposition Behavior of a Melt-Quenched AI-Zr Alloy Z A R I F F A. C H A U D H U R Y AND C. S U R Y A N A R A Y A N A
CeHtre of Advanced Study in Metallur~,,y. Department of Metallurgical En~,im,erin~,,, Banaras Hindu University. VARANASI--221 005, India
Melt quenching of an A I - Z r alloy containing 9.4 wt.% Zr resulted in the formation of a supersaturated solid solution. Elevated temperature annealing treatments led to the formation of a metastable AI~Zr phase having an ordered cubic Cu~Au-type structure with a = 0.4073 nm and at a later stage the equilibrium A h Z r phase having a tetragonal structure with a 0.4013 nm and c 1.7321 nm. Electron microscopic investigations on the morphology and structural features of the as-quenched as well as transfc)rmation products formed during decomposition are discussed in detail.
Introduction In recent years melt quenching techniques have gained wide popularity not only because of their ability to produce extremely high cooling rates but because of the unusual properties associated with their products. Melt quenching could lead to the formation of supersaturated solid solutions at composition ranges far beyond those obtained by conventional methods. In a variety of alloy systems, new non-equilibrium crystalline intermediate phases and metallic glasses have also been produced. The innumerable results obtained on several alloys by the application of meltquenching techniques have been well reviewed [I-5]. A large number of investigations have been carried out on the morphology of the phases and the precipitation behavior in AI-Zr alloys. Transmission electron microscopy had been extensively employed lbr this purpose [6-11]. The majority of these investigations were concerned with chill-cast A1-Zr alloys containing up to 2 wt.% Zr. The effect of cooling rate on the grain size of alloys and grain refinement behavior were also studied, Formation of a metastable-ordered cubic AI3Zr phase has been found to be responsible for the improved properties of AI-Zr alloys. Melt quenching of AI-Zr alloys has also yielded some interesting resuits. Varich el al. [12] studied the effect of superheating the melt on the c© Elsevier Science Publishing Co., Inc., 1984 52 Vanderbilt Ave., New York, NY 10017
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structure and properties of rapidly quenched A1-Zr alloys containing 2.2 wt.% Zr. Varich and Sheyko [13] studied the stability and thermal expansion of A1-Zr alloys. Gudzenko and Polesya [14] observed the formation of an amorphous phase in alloys containing 45 to 53 at.% Zr. Recently, it has been shown that the solid solubility of Zr in Al could be enhanced by melt quenching up to about 8.6 wt.% Zr [15]. The maximum solid solubility of zirconium in aluminum under equilibrium conditions is reported to be 0.28 wt.% [16]. The most Al-rich intermediate phases in the AI-Zr system are AI3Zr, AIzZr, A13Zr2 and AIZr [17-19]. As part of an on-going program on the formation and decomposition behavior of metastable phases in aluminum alloys [20-26], we had undertaken detailed investigations on the A1-Zr [27] system. Since not much work was reported on the metastable behavior of rapidly quenched AI-Zr alloys, it was decided to investigate the impact of rapid quenching from the melt on the constitution and decomposition behavior of metastable phases in AI-Zr alloys.
Experimental Procedure An AI-Zr alloy with a nominal composition of 9.4 wt.% Zr was prepared in an arc-melting furnace under a protective argon atmosphere. The alloy was remelted to achieve homogenization. A small quantity of the alloy was induction-melted and rapidly quenched from the melt using the conventional "gun" technique, wherein the cooling rate achieved was estimated to be about 107 K/sec. The as-quenched foils were thin enough in many areas to be examined directly in a transmission electron microscope fitted with a goniometer stage and operating at 100 kV. The as-quenched foils were given proper annealing treatments after encapsulation under vacuum in quartz capsules. In-situ hot stage electron microscopy was also carried out using the hot-stage specimen holder to follow the transformation sequences.
Results AS-QUENCHED STRUCTURE Remarkable extension of solid solubility of Zr in Al has been observed on melt quenching. Figure 1 shows a bright-field electron micrograph featuring elongated grains of the supersaturated solid solution. The corresponding diffraction pattern is shown in an inset. The absence of any second phase in the microstructure and clear single crystal fcc reflections
Decomposition Behavior of Al-Zr Alloy
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FrG. 1. Microstructure of the as-quenched AI-9.4 wt.% Zr alloy showing the formation of supersaturated solid solution. The corresponding diffraction pattern is shown as an inset. in the diffraction pattern suggest that all the 9.4 wt.% Zr had dissolved in AI to form a homogeneous supersaturated solid solution. These observations clearly indicate that the solid solubility of zirconium in aluminium in the melt-quenched condition can be increased to at least 9.4 wt.%.
Defect Structure As reviewed by Jones [28], rapidly-quenched metals and alloys show a variety of defect structures. This has been confirmed recently in our laboratory as well [24, 25, 29]. Though the solid solution grains formed in the present investigation were defect-free many times, interesting dislocation structures were sometimes observed in some of the grains. As a typical example, isolated dislocations could be noticed near the grain boundary in Fig. 2a. It was also possible, occasionally, to observe the dislocations arranged into regular low-angle boundaries as shown in Fig. 2b. The cell structure, surrounded by the low-angle boundaries, was assumed to have formed as a result of the recovery of stresses introduced during the quenching process, a situation similar to what happens in plast-
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(a)
(b) Flo. 2. Typical dislocations observed in the as-quenched microstructure: (a) isolated dislocations, and (b) low-angle boundaries.
Decomposition Behavior of AI-Zr Alloy
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ically deformed and annealed metals. The dislocations might have been formed because of the plastic deformation induced by rapid quenching stresses and/or aggregation of vacancies. ANNEALING BEHAVIOR With a view to evaluating the stability of the supersaturated solid solution formed in this system, the as-quenched foils were heat treated both externally and in-situ. Salient features of these observations are here detailed.
External Heat Treatments The metastable solid solution was quite stable up to at least about 573°K. On annealing the foils at 593°K for 3.6 ks, the discontinuous precipitation of a second phase along the grain boundaries (Fig. 3) could clearly be seen. Simultaneously, it was noticed that the matrix also undergoes some transformation as revealed by the formation of extremely fine rodlike precipitates. Although it was not easy to distinguish between these
Fro. 3. Microstructure showing the discontinuous precipitation along the grain boundary on low-temperature annealing. Also notice the fine structure inside the matrix grains.
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two types of precipitates from electron diffraction patterns, it should be emphasized that diffraction patterns from this area could be indexed on the basis of an ordered cubic structure. Better contrast from the precipitates could be obtained when the foils were annealed at higher temperatures on the order of 773°K for 3.6 ks. Figures 4a to 4c show the bright-field and dark-field micrographs and the corresponding diffraction pattern. This diffraction pattern also could be indexed as arising from an ordered fcc structure. F r o m the available literature, it can be inferred that this is the metastable AI3Zr compound. A point of interest in this micrograph is the shape of the precipitate particles, which is nearly circular in cross section. This observation, coupled with the microstructure obtained on annealing at lower temperatures (Fig. 3), clearly suggests that metastable A13Zr forms as rods. Figure 5a shows that the matrix contains, in some areas, precipitates of metastable AI3Zr which are interconnected along the long axis of the rods. This gives the impression of long arrays of precipitates arranged parallel to each other. It may also be noticed in the micrograph that precipitation of yet another phase is just starting. The size of these particles is much larger than that of metastable A13Zr. When the alloy foils were
(a)
FIc. 4. (a) Precipitate particles uniformly distributed in the aluminum matrix after annealing the foils at 773°K for 3.6 ks; (b) the corresponding dark-field micrograph showing the shape of the precipitates; (c) [ 111] diffraction pattern taken from these particles indicating that this phase is the metastable cubic AIaZr.
Decomposition Behavior of Al-Zr Alloy
(b)
(c) Fl(3. 4.
(Continued)
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238
ipm (a)
d
(b) FIG. 5. (a) Precipitation of the rodlike precipitates after annealing at 773°K for 3.6 ks. Precipitation of another phase seems to just be starting; (b) grain boundary precipitation of the tetragonal equilibrium intermediate phase.
Decomposition Behavior of AI-Zr Alloy
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(a)
(b) FIG. 6. (a) Microstructure showing lhe fan-shape of the precipitates of the melaslable AI~Zr phase; (b) the corresponding dark-field micrograph.
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Fro. 7. is [001].
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Diffraction pattern taken from the metastable rodlike precipitates. Foil normal
FIG. 8. A high magnification micrograph showing the shape and size of the rodlike precipitates after annealing at 773°K for 10.8 ks.
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annealed for 10.8 ks at 793°K, much more precipitation of this new phase along the boundaries of the elongated grains could be observed (Fig. 5b). While the rodlike precipitates have a diameter of about 15 nm and are spaced about 40 nm from each other, the grain boundary precipitates are much coarser, measuring about 80 nm at this stage. Diffraction patterns from this phase indicated a tetragonal crystal structure and a correspondence to the equilibrium AI3Zr phase. Figures 6a and 6b show a pair of bright-field and dark-field electron micrographs bringing out the fan-shape morphology of the metastable AI3Zr phase. A typical diffraction pattern from these precipitates is shown in Fig. 7, which unambiguously confirms that this is the At3Zr-ordered cubic phase with a = 0.4073 nm. A high magnification micrograph showing the morphology of the precipitate phase is presented in Fig. 8. The growth, thickening and branching off of the rods, is worth observing in this micrograph. Figure 9a shows the micrograph recorded from an alloy annealed for 10.8 ks at 773°K. Large-size precipitates of the tetragonal equilibrium
(a) FIG. 9. (a) Bright-field micrograph showing the formation of the equilibrium precipitates. The particles are uniformly distributed; (b) electron micrograph showing the precipitates inside the equi-axed grains; (c) a high magnification bright-field electron micrograph showing the shape of the precipitates within the grains. Note also the absence of precipitation near the grain boundaries.
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(b)
(c) FIG. 9.
(Continued)
Decomposition Behavior of Al-Zr Alloy
243
AI3Zr phase can be seen clearly here. Rodlike precipitates can also be seen in the background. However, even for the same heat treatment conditions, a different type of microstructure is observed in some other regions of the foil, as shown in Fig. 9b. Precipitation of a second phase inside the equiaxed grains is worth noticing. Another point of interest is the absence of precipitation near the grain boundaries, although one cannot compare this with the precipitate free zones observed in solid-state quenched and aged precipitation hardening alloys. A high-magnification micrograph of the central grain of Fig. 9b is shown in Fig. 9c, which indicates that the precipitate here also perhaps has a cylindrical shape. It has not been possible to get diffraction patterns from these precipitates, and so nothing can be said about the composition and crystal structure of this phase, except that, from the shape of the precipitate particles, it may be assumed that these belong to the metastable AI3Zr compound. Long annealing treatments at 903°K resulted in the complete disappearance of the metastable AI3Zr phase; the fact that the microstructure at this stage consists of only aluminum solid solution and the tetragonal AI3Zr phase suggests that equilibrium has been established. A typical micrograph corresponding to this situation is shown in Fig. 10. Two rep-
Fl(;. 10. Bright-field electron micrograph showing the distribution of the equilibrium AI3Zr phase along the equi-axed grain boundaries of the aluminum matrix after annealing the foils at 903°K.
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(a)
(b) FIG. ll. Typical single crystal electron diffraction patterns taken from the tetragonal equilibrium Al3Zr phases. Foil normals are (a) [591] and (b) [lll].
Decomposition Behavior of AI-Zr Alloy
245
5pro (a)
IIIII
(b) Fie. 12. (a) Bright-field electron micrograph of a big equilibrium AI~Zr particle; (hi the corresponding dark-field micrograph.
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(a)
(b) FIG. 13. (a) Discontinuous precipitation of the metastable cubic AI3Zr phase; (b)-(d) a sequence of micrographs showing the growth of the discontinuous phase on continued annealing.
Decomposition Behavior of Al-Zr Alloy
(c)
(d)
FIG. 13.
(Contintted)
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resentative diffraction patterns recorded from these precipitate particles are shown in Figs. 1 la and 1lb. From these and several other patterns, it had been possible to confirm that this phase is the equilibrium AI3Zr compound having a tetragonal structure with the lattice parameters a = 0.4013 nm and c = 1.7321 nm. Figure 12 shows a pair of bright-field and dark-field micrographs featuring a very large AI3Zr particle indicating that further annealing leads only to coarsening of the precipitate particles.
Hot-Stage Electron Microscopy The as-quenched AI-9.4 wt.% Zr alloy foils have also been heat treated in situ using the hot-stage specimen holder. During continuous heating, it was observed that the supersaturated solid solution was quite stable up to about 593°K, and that precipitation took place beyond this point. On continued heating, discontinuous precipitation of metastable AI3Zr could be noticed at about 668°K (Fig. 13a). Figures 13b to 13d show a series of three micrographs recorded at 750°, 800 ° and 825°K depicting the growth of the discontinuous phase which gradually enveloped the whole of the untransformed matrix. Diffraction patterns from these areas clearly revealed that the phase is metastable AI3Zr, confirming the observation made during external heat treatments. At still higher temperatures the equilibrium tetragonal AI3Zr phase could be detected. Discussion
S U P E R S A T U R A T E D SOLID SOLUTION Earlier investigators [9, 12, 15, 30-32] also reported that the solid solubility of Zr in AI could be enhanced by rapidly quenching the melts of these alloys. Our results indicate that a maximum of 9.4 wt.% Zr could be dissolved in a metastable condition in the solid state. Earlier investigators reported a much lower value. This observation suggests that the extended value is about 35 times the maximum equilibrium value of 0.28 wt.% Zr. Work with more concentrated alloys will indicate whether still higher supersaturation can be achieved. While the previous investigations were carried out using either chill casting or a two-piston device, the present results were obtained by "gun" quenching, wherein the cooling rates obtained are definitely higher than those in any other technique. Thus, it was only natural that the solid solubility extension was more prominent in our studies. This is essentially due to the flow of melt into thin foils. The elongated grains noticed in the
Decomposition Behavior of AI-Zr Alloy
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foils is further evidence that the molten droplets did flow on the substrate. However, in the relatively thicker portions of the foils the grains were equiaxed. Even here precipitation of a second phase could not be detected in the as-quenched condition. DECOMPOSITION BEHAVIOR The decomposition behavior of aluminum-rich AI-Zr solid solutions produced by solid-state quenching [6, 7, 33] and melt quenching [15, 32] has been studied earlier. Sahin and Jones [15] identified three distinct stages during isochronal annealing of an AI-4.9 wt.% Zr alloy: 1. Homogeneous modular separation, 2. Discontinuous precipitation of metstable cubic AI3Zr, and 3. Transformation to coarse plates of equilibrium tetragonal AI~Zr. In their studies on chill-cast alloys, Ohashi and Ichikawa [9] and Hori et al. [32] detected a petal-like structure in low-Zr alloys. This is not, however, confirmed by Sahin and Jones [15]. The supersaturated solid solution was quite stable up to a temperature of 593°K. Beyond this point precipitation started. A point of interest here is the discontinuous precipitation of metastable cubic AI3Zr phase along with a fine rodlike precipitate in the matrix. This fine structure in the matrix cannot be ascribed to the clustering effect observed in solid-state quenched and aged alloys, although Fontaine et al. [34] and Bonefa~i~ et al. [35] observed clustering effects due to the segregation of solute atoms in rapidly quenched A1-Fe and AI-Ni systems, respectively. The diffraction patterns from the A1-Zr specimens subjected to low-temperature annealing, however, do not show any evidence of clustering. This microstructure has been interpreted as consisting of a fine rodlike precipitate. The diffraction patterns clearly indicate that this is the cubic metastable AI3Zr phase. Depending on the orientation in which these rodlike precipitate particles are observed, either rods or round particles are noticeable. Evidence has also been collected to show that the fan-shaped precipitate observed by others is nothing but the rods arranged one next to the other. Figures 6 and 8 bring out these features very clearly. Hotstage experiments indicated that the interface velocity is very high and that the fan-shaped precipitates grow very rapidly into the matrix (Fig. 13). Ryum [6] has reported the formation of fine spherical precipitates. especially between the arms of the fan-shaped precipitates. These may be the rods viewed along the long direction. The petal-shaped precipitates have not been observed in our present investigation.
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Annealing at higher temperatures resulted in the precipitation of the equilibrium tetragonal A13Zr phase. All earlier investigators reported that the equilibrium A13Zr phase has a morphology of coarse plates which could be considered to have transformed from the metastable A13Zr phase. However, in our present investigation, we have noticed that this has a bulky shape and that, given a suitable annealing treatment, these can grow to very large sizes. The diffraction patterns could be unambiguously identified. From the foregoing it appears that, although one can clearly identify two stages in the transformation of the supersaturated solid solution to the equilibrium phases, these two stages seem to overlap quite extensively. Simultaneous precipitation of both the metastable and stable A13Zr phases was detected at as low a temperature as 793°K, while at high temperatures of about 900°K only the equilibrium phases could be observed. Conclusions On the basis of the results presented above the following conclusions can be drawn. 1. The solid solubility of Zr in A1 can be extended by melt quenching at least up to 9.4 wt.%. This is the highest value of supersaturation reported so far in this system. 2. Both in situ and external heat treatments suggest that the supersaturated solid solution is stable up to about 593°K. 3. Annealing at temperatures in excess of 600°K results in the precipitation of metastable ordered AI3Zr phase in the form of rods. Fan shape of these precipitates was also noticed. 4. Higher temperature annealing leads to the formation of the equilibrium tetragonai AI3Zr phase.
The authors are grateful to Professor T. R. Anantharaman, Program Coordinator, Center o f Advanced Study, for encouragement, and to the Head of the Department o f Metallurgical Engineering, Banaras Hindu University, for provision of laboratory facilities. They are also thankful to Dr. Shigenori Hori of Osaka University, Japan for supplying the A l Zr alloy used in this investigation, and to Dr. G. V. S. Sastry for a critical reading of the manuscript. References 1. T. R. Anantharamanand C. Suryanarayana,Review: A decade of quenchingfrom the melt, J. Mater. Sei. 6:1111-1135 (1971).
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2. H. Jones and C. Suryanarayana, Rapid quenching from the melt: An annotated bibliography 1958-72, J. Mater. Sci. 8:705-753 (1973). 3. H. Jones, Splat cooling and metastable phases, Rep. Progr. Phys. 36:1425-1497 (1973). 4. H. Jones, Rapid Solid(l~cation ~/' Metals and Alloys, Monograph No. 8, Institution of Metallurgists, London 11982). 5. C. Suryanarayana, Rapidly quenched metals--a bibliography 1973-1979. IFl/Plenum. New York (1980). 6. N. Ryum, Precipitation and recrystallization in an AI-0.5 wt.C/c Zr alloy, Acta Mct. 17:269-278 (1969). 7. O. lzumi and D. Oelschlagel, Strukturelle Untersuchung der ausscheidung in einer aluminium-legierung mit I. 1 Gew.-'~ Zr, Z. Metallkde. 60:845-851 (1969). 8. V. I. Dobatkin, V. I. Elagin, V. M. Fedorov and R. M. Sizova, Decomposition of saturated solid solutions in granulated aluminium alloys, Russian Met. 2:122-127 (1970). 9. T. Ohashi and R. Ichikawa, Grain refinement in aluminium-zirconium and aluminiumtitanium alloys by metastable phases, Z. Metallkde. 64:517-521 (1973). 10. E. Nes and H. Billdal, Non-equilibrium solidification of hyperperitectic AI-Zr alloys, Acta Met. 25:1031-1037 (1977); The mechanism of discontinuous precipitation of the metastable AI~Zr phase from an AI-Zr solid solution, Acta Met. 25:1039-1(146 (1977). 1I. S. Hori, H. Kitagawa, T. Masutani and A. Takehara, Structure and phase decomposition of rapidly solidified aluminium-zirconium solid solutions. J. Japan lttst. Li~,,ht Metal~ 27:129-137 (1977). 12. N. 1. Varich, R. B. Lyukevich, L, E. Kolomoytseva, A. N. Varich and V. V. Maslov. Effect of superheating the melt on the structure and properties of rapidly cooled AI-Zr alloys, Phys. Met. Metallogr. 27(2): 176-179 (1979). 13. N. I. Varich and T. I. Sheyko, Thermal expansion of Al-Mo, AI-Zr alloys prepared at high cooling rates, Phys. Met. Metallogr. 30(2):231-232 (1970). 14. V. N. Gudzenko and A. F. Polesya, Structure of zirconium-aluminium alloys rapidly cooled from the liquid state, Phys. Met. Metallogr. 39(6): 177-179 (1975). 15. E. Sahin and H. Jones, Extended solid solubility, grain refinement and age-hardening in AI-1 to 13 wt.% Zr rapidly quenched from the melt, in Rapidly Quenched Metah' II1, (B. Cantor, Ed.), The Metals Society. London (1978). pp. 138-146. 16. L, E. Mondolfo, Alaminiam Alloys--Structure and Propertie.~, Butterworths, London (1976). 17. M. Hansen and K. Anderko, Constitution (~['Bipm~3' Alloy.s, McGraw-Hill, New York (1958). 18. R. P. Elliott, Constitution ~f'Binao' Alloys: Fir,st Supplement. McGraw-Hill, New York (1965). 19. F. A. Shunk, Constitution ~fBinao' Alloys: Second Supplement, McGraw-Hill, New York (1969). 20. Z. A. Chaudhury, G. V. S. Saslry and C. Suryanarayana. Phase transformations in rapidly quenched aluminium-ruthenium alloys, Z. Metallkde. 73:201-206 (1982). 21. Z. A. Chaudhury and C. Suryanarayana, Metastable phases in vapour-deposited AIRu alloys, J. Mater. Sci. 17:3158-3164 (1982). 22. Z. A. Chaudhury and C. Suryanarayana, Annealing behaviour of vapour-deposited AIRh thin films, Thin Solid Films 98:233--239 (1982). 23. Z. A. Chaudhury and C. Suryanarayana, Al~3X4-type phases in aluminium-group VIII metal systems, J. Less-Common Metals 91:181-187 (1983). 24. Z. A. Chaudhury and C. Suryanarayana. Electron microscopic investigations on a meltquenched AI-Rh alloy, J. Mater. Sci. 18:3011-31)22 (1983). 25. G. V. S. Sastry and C. Suryanarayana, Metastable effects in melt-quenched AI-Pd alloys, Mater. Sci. & Eng. 47:193-208 {1981).
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26. G. V. S. Sastry, C. Suryanarayana and G. Van Tendeloo, A structural study of vapourdeposited A1-Pd alloys, Phys. Stat. Sol. (a) 73:267-278 (1982). 27. Z. A. Chaudhury, Structure of rapidly quenched aluminium alloys, Ph.D. thesis, Banaras Hindu University, Varanasi, India, 1983. 28. H. Jones, Lattice defects in splat-quenched metals, in Vacancies '76, (R. E. Smallman and J. E. Harris, Eds.), The Metals Society, London (1977), pp. 175-184. 29. L. R. K. Rao, C. Suryanarayana and S. Lele, Lattice defects in melt-quenched aluminium and its alloys, Aluminium 58:214-217 (1982). 30. L. M. Burov and A. A. Yakunin, Effect of the rate of cooling on the composition of solid solutions in binary alloys based on Al, Russian J. Phys. Chem. 42:540-541 (1968). 31. A. K. Kushnereva and I. V. Salli, The metastable crystallization of binary alloys, Metallofizika Desp. Mezhved. 56:115-120 (1968). 32. S. Hori, S. Saji and A. Takehara, Structure of rapidly solidified A1-Zr alloys and its thermal stability, in Proc. Fourth International Conf. on Rapidly Quenched Metals, (T. Masumoto and K. Suzuki, Eds.) The Japan Institute of Metals, Sendai (1982), Vol. II, pp. 1545-1548. 33. E. Nes, Precipitation of the metastable cubic A13Zr phase in subperitectic AI-Zr alloys, Acta Met. 26:499-506 (1972). 34. A. Fontaine, J. Dixmier and A. Guinier, Mechanical properties and structural transitions in splat-cooled AI-Mn and AI-Fe alloys, Fizika, 2, Suppl. 2:23.1-23.6 (1970). 35. A. Bonefa~i6, M. Kerenovi6, A. Kirin and D. Kunstelj, Segregation of solutes in AlNi and AI-Sn and its influence on their mechanical properties, J. Mater. Sci. 10:243251 (1975).
Received October 1983; accepted January 1984.