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The effect of coherent particles on the annealing behavior of a cold-worked Cu-2% Co alloy
The Effect of Coherent Particles on the Annealing Behavior of a Cold-worked Cu-2°/o Co Alloy* L. E. T A N N E R
AND I. S. S E R V I
LedgemontLaboratory, Kennecott CopperCorporation, Lexington,Mass. (U.S.A.) (Received April 14, 1966)
S UMMAR Y A polycrystalline Cu-2°//o Co alloy, overaged at 600°C to produce a fine dispersion of coherent
particles, was severely dejbrmed by swagin9 and then annealed. The resulting microstructures were examined by electron transmission microscopy; and related mechanical properties were determined by tensile tests. Swaging plastically dejbrmed the matrix and the particles, producing a substructure comprised of "deJbrmation" bands containing a high density of tangled dislocations and diffuse cells separated by narrow "transition" bands of well-defined subgrains. Recovery and recrystallization processes were inhibited by the presence of the second phase and did
not become Jully operative until annealing took place in the vicinity of the original aging temperature, i.e., 500°C and above. Recovery resulted in the polygonization of the deformation bands and the restoration of spherical shape to the coherent particles within these regions. Recrystallization.proceeded by means of boundary migration, initially in the transition bands utilizing the cold-work produced boundaries and then in the deformation bands with the poly9onized boundaries. The migration process involved the coalescence of the coherent phase and appeared to be a variation of the discontinuous precipitation phenomena.
RF.SUMF. Dans un alliage polycristallin de cuivre d 2°//0 de cobalt, une fine dispersion de particules cohOrentes a OtO obtenue par un traitement de survieillissement d 600°C. Le m~tal a ensuite ~tO soumis gt une dOformation intense par dtampage et Ot~ recuit. Les microstructures obtenues gl diff~rents stades ont Otk examinOes par microscopie ~lectronique en transmission et les caractOristiques mdcaniques correspondantes ont ~t~ dOtermin~es au moyen d'essais de traction. Uop~ration dYtampage a pour effet de d~Jbrmer plastiquement dr.la ibis la matrice et les particules et donne naissance dune structure de bandes : bandes de "d~Jbrmation" comportant des zones gt densit~ ~lev~e de dislocations enchev~tr~es et des cellules plus diJJi~ses, bandes de "transition" plus Otroites constituOes par des sous-grains bien dOfinis et sOparant les bandes du premier type. Les phdnombnes de restauration et de recristallisation sont entrav~s ou retardOs par les particules de la seconde phase: ce n'est que Materials Science and Engineering-
lorsque la temperature de recuit est voisine de la tempdrature du vieillissement initial, c'est-~t-dire 5 0 0 ° C ou au-dessus, que ces processus deviennent r~ellement efficaces. La restauration se traduit par une polygonisation des bandes de d~Jbrmation et le retour d la Jorme sph~rique des particules coh~rentes situ~es dans ces r~gions. La recristallisation s' effectue par migration des joints: au ddbut ce sont les joints de la structure d'~crouissage, situds dans les bandes de transition, qui se d~placent, puis la recristallisation se poursuit par mouvement des sous-joints de polygonisation qui se sont form,s dans les bandes de d~Jbrmarion. La migration des joints est accompagn~e d'une coalescence des particules de la phase coh~rente, si bien que ce mdcanisme de recristallisation semble constituer une variante des processus de precipitation discontinue. * Presented in part at the Fall Meeting of the Metallurgical Society of AIME, October 19, 1965, Detroit, Mich., U.S.A.
Elsevier Publishing Company, Amsterdam - Printed in the Netherlands
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Z USAMMENFASSUNG Eine bei 600°C zur Erzeugung einer feinen Verteilung kohiirenter Teilchen angelassene polykristalline Cu2% Co-Legierung wurde durch Hiimmern stark verformt und dann angelassen. Mit Hilfe yon elektronenmikroskopischen Durchstrahlungsaufnahmen wurde die entstandene Mikrostruktur untersucht. Zugeh6rige mechanische Eigenschajten wurden dutch Zugversuche bestimmt. Matrix und Ausscheidungen wurden durch Hiimmern plastisch verjormt. Die erzeugte Substruktur umJaflte "Verformungsb?inder" mit einer hohen Dichte yon Versetzungsnetzwerken und verteilte, dutch enge "t)bergangsbiinder" aus wohldefinierten SubkiJrnern getrennte Zellen. Erholungs- und Rekristallisations-
1. INTRODUCTION The strengthening of metals by the introduction of a fine dispersion of a second phase is dependent upon the ability of the particles to inhibit dislocation movement under applied stress ~. A large variety of dispersions can be produced by the decomposition of supersaturated solid solutions, internal oxidation of alloys, powder metallurgy techniques, etc. These dispersions, in the form of either coherent or incoherent particles, may consist of metallic phases having mechanical properties not too different from those of the matrix or of very hard, insoluble ceramic compounds. A unique characteristic of these alloys after plastic deformation is that their work-hardened state often remains stable to relatively high temperatures 2. That is, depending on the type, size, and volume fraction of the second phase, there may be a definite resistance to recovery and recrystallization processes at temperatures where such behavior would normally occur in single-phase alloys of comparable solid solution composition. In fact, in the case of certain internally-oxidized alloys, structural stability can approach the melting point of the matrix 3. While these effects have been known for some time, it is only recently, with the advances in electron transmission microscopy techniques, that detailed studies of the variables involved in these processes have been undertaken. Doherty and Martin 4-6 investigated the annealing behavior of deformed A1-Cu alloys having dispersions of the incoherent 0 phase (Cu2A1). These investigators
prozesse waren durch die Anwesenheit der zweiten Phase gehemmt und wurden nicht voll wirksam bis zum Anlassen in der Niihe der urspriinglichen Anlafltemperatur, also bei 500°C oder hiiher. Die Erholung ffihrte zur Polygonisation der Verformungsbiinder und zur Wiederherstellung der Kugelgestalt der kohiirenten Teilchen innerhalb dieser Gebiete. Die Rekristallisation erjolgte durch Korngrenzenwanderung, und zwar anfdnglich in den Obergangsb?indern unter Ausnutzung der durch die KaltverJormung erzeugten Korngrenzen und dann in den VerJbrmungsbiindern mit den polygonisierten Grenzen. Der Wanderungsprozefl war mit der Zusammenlagerung der kohiirenten Phase verbunden und schien eine Variation der Erscheinungen bei diskontinuierlicher Ausscheidung zu sein.
found that when particles are spaced less than 1.0# apart there is a measurable reduction in the rate of "nucleation" of strain-free regions through the inhibition of low-angle boundary migration. Similar observations were made by Leslie et al. in Fe-Cu alloys 7 and by Williams for a number of internallyoxidized Cu-base alloys a. In still another investigation, the mechanical properties of a Cu-Co alloy, aged to produce a fine dispersion of coherent particles and then deformed, show its work-hardened state to be retained after each of a series of 1 h annealing treatments at temperatures up to 600°C 9. However, related structural observations were not reported. Thus, the current study utilizing electron transmission microscopy was undertaken with a comparable alloy treated in a similar manner in order to determine how the coherent second phase affects recovery and recrystallization. The results will be discussed in light of the current models proposed for these phenomena in pure metals and singlephase alloys ~°.
2. EXPERIMENTAL PROCEDURE
A Cu-Co alloy of nominal composition 2.0 wt. ~o Co was prepared from 99.999 + ~o pure copper and 99.84~o pure electrolytic cobalt. The Czochralski method for solidification was used and proved to give minimum microsegregation. Chemical and spectrographic analyses indicated that the cobalt content was 2.05 wt. ~ and that impurities consisted of: Mater. Sci. Eng., 1 (1966) 153-161
Cu-2~o Co
ANNEALING BEHAVIOR OF COLD-WORKED
< 100 p.p.m. Ni < 10 p.p.m. Mg, Fe < 1 p.p.m. AI, Ca 10 p.p.m. C The ingot was cold-swaged to wire of 0.040 in. diam. without intermediate annealing. The wire was then solution-treated in vacuo at 1000°C for ¼h followed by iced brine quenching. Aging was carried out at 600°C for 6 h in an inert atmosphere. According to Livingston 11 the particle size and spacing resulting from this treatment gives optimum mechanical properties. A portion of the stock was cold-swaged to a reduction in area of approximately 50~, and individual samples were annealed as follows: Temperature(°C): Time(h):
70
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ALLOY
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~ deformed [ 60 _--aged
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50
d ~ 4o p--.q w x 50 u_
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20
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/aged a deformed --
O RT
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1(30 200
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.
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300
400
500
600
ANNEALING TEMPERATURE (*C)
Fig. 1. Room-temperature tensile properties of Cu-2 ~o Co in the aged, deformed and annealed conditions. The alloy was aged at 600°C for 6 h. Annealing treatments followed cold-swaging to a 50 ~o reduction in area and were for 1 h at the indicated temperatures.
300 400 500 600 700 1 1 1 1, 24 1
Specimens given the 1 h treatments at 300-600°C were tested in tension on an Instron machine to determine yield and ultimate strength, as well as tensile elongation. A portion of each wire was examined in a Hitachi HU-11A electron microscope equipped with a biaxial tilting specimen stage. An accelerating voltage of 100 kV was used. The wires were first mechanically polished parallel to their axis to a thickness of about 0.005 in. Thin foils were prepared by electrothinning utilizing a fine jet technique developed by Hugo and Phillips i2. The electrolyte was a solution of two parts phosphoric acid to one part water.
3.
EXPERIMENTAL RESULTS
The tensile data obtained from the aged C u - 2 ~ Co alloy after the mechanical and thermal treatments are plotted in Fig. 1. These results are in substantial agreement with those reported by Jones 9, showing only small changes from the cold-worked condition. These changes consist of a slight loss of strength, as well as a small, but distinct, increase in ductility with increasing, annealing temperature. Subtle as these may be, they indicate that some structural alterations occurred during the thermal treatments. The microstructure of the aged and undeformed condition is shown in Fig. 2. The cobalt concentration remaining in solid solution is expected to be about 0.35~ 11. This was verified by electrical resistivity measurements (i.e., 4/~f~-cm at room temperature corresponds to 0.35~ Co13). The micro-
Fig. 2. Solution treated at 1000°C for ¼ h, iced-brine quenched and aged at 600°C for 6 h. (z48,000)
graph exhibits a dispersion of coherent particles within a single grain*. The spherical Co-rich particles are revealed by diffraction contrast arising from elastic strains in the matrix due to lattice mismatch between the two phases TM15. The size of the particles was determined from micrographs by taking the length of the line of no contrast as their diameteP 4. An average value of 150 A was observed and may be considered to be accurate to 25~. The foil thickness was assumed to be 1000 A, but in regions examined it may have departed from this estimate by a factor of two. Particle density was * The average grain diameter of the wires was of the order of 0.2 mm.
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Fig. 3. Aged and deformed in tension to 11~o elongation. ( x 48,000)
determined, and then interparticle spacing was calculated in a manner proposed by Ashby and Ebling 16 assuming the particles to be randomly dispersed. A value for the latter of 1100 A was obtained which is 2 5 ~ smaller than that reported by Bonar 17 for an alloy of similar composition and treatment. Interpretation of the deformed microstructure was particularly difficult because of the poor visibility of individual dislocations. The result of 11~ tensile elongation is seen in Fig. 3, and comparison with Fig. 2 indicates that plastic flow produced substantial alteration in particle strain contrast, as well as irregular variations in the orientation of the lines of no contrast. Deformation is believed to change the shape of the particles while coherency is retained is and could account for the former. The latter would be expected from a random distribution of local lattice tilts and thus suggests that the dislocations are in a tangled array rather than in a well-defined cell structure of the type characteristic
Fig. 4. Aged and cold-swaged to a 50~o reduction in area. Banded structure elongated in the direction of the wire axis. Deformation bands "A" separated by transition bands "B". (x 25,000)
Mater. Sci. Eng., 1 (1966) 153-161
ANNEALING BEHAVIOR OF COLD-WORKED
of deformed unalloyed Cu 19. The inhibition of cell formation is consistent with Swann's observation that a fine dispersion of a ductile second phase tends to confine dislocation movement to the primary slip systems and inhibits cross-slip in the initial stages of deformation 2°. Another source for changes in particle strain contrast, which would also account for the difficulties in dislocation resolution, is the superposition of the elastic strain fields of coherent particles and dislocations. This should result in rather complicated contrast that would vary from area to area. More severe deformation, to be discussed below, tends to delineate regions having large variations in dislocation density, but contrast never approached the degree of sharpness observed in most worked single-phase materials. Swaging is a notably complicated and nonuniform method for working metals. It promotes glide on many intersecting slip systems and produces a complex combination of lattice rotations and tilts which develop the elements of a fiber texture 21. The resulting banded structure, seen in Fig. 4, accommodates the. lattice reorientations and is similar to that found in heavily-rolled pure metals and alloys 1°. The regions marked "A" appear to have a high density of dislocations in both relatively uniform tangles and poorly defined cells but also include some sharp boundaries. These regions are separated by narrow bands (designated "B") that are comprised of small, elongated subgrains bounded by discrete subgrain walls. Dislocation density seems to be considerably lower within this substructure. Close examination reveals the coherent particles, but their contrast is markedly altered by the deformation. Selected area diffraction indicated only small orientation variations across the tangled regions (i.e., <<,2°), while there was generally about an 8-10 ° disorientation" across the sharp bands. Furthermore, the orientation changes in these bands were in incremental steps of 1-2 ° from one subgrain to the next. The sharp boundaries are the product of glide polygonization made possible by interactions of dislocations moving on the intersecting slip planes. In a qualitative sense the more tangled regions of this structure are analogous to the "deformation" bands, and the narrow subgrain regions correspond to the "transition" bands described by Walter and Koch 22 for rolled materials. These designations will be adopted here. Only small variations in the dislocation substructure could be observed after heat treatments at 300 and 400°C. In the main, these were confined
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(a)
Fig. 5. Aged and swaged alloy annealed at 500°C for 1 h. (a) Dislocations rearranged to form a distinct cell structure within a deformation band. ( × 28,000) (b) Coalescence of the coherent Co-rich phase and migration of a cold-work produced boundary within a transition band. ( × 40,000)
to local rearrangements which tended to increase and sharpen the cell structure in the deformation bands. Such rearrangements were clearly evident after the 500°C treatment as seen in Fig. 5(a). However, a more dramatic change in microstructure was found in isolated regions within the transition bands. In Fig. 5 (b), sub-boundary A'B'C' is pinned by a few coarse incoherent particles (y)*. It has migrated from position ABC, which consists of another string of incoherent particles (x), and the included area between ABC and A'B'C' is free from both the coherent phase and dislocations. * The particles are no longer in the foil, having been lost during electrothinning, but are identified by the remaining holes which appear white in the micrographs.
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served in this alloy system 23. Subsequent migration of the new sub-boundaries, in the manner described above, eliminated a portion of the coherent phase through particle coalescence (Fig. 7(b)). Annealing at 700°C greatly accelerated polygonization, boundary migration, and particle coalescence to the point where delineation of separate deformation and transition band effects was rather difficult. In addition, the situation was further
Fig. 6. Aged and swaged alloy annealed at 600°C for 1 h. Recrystallized grains at the edge of a transition band growing into a dislocation-dense region of a deformation band. (x 48,000)
In addition, some narrow regions adjacent to the leading side of the advancing interface also appear to be free from coherent particles (e.g., between A' and B'). The above process of particle coalescence and boundary migration is considerably more advanced after the 1 h treatment at 600°C. In Fig. 6, a number of recrystallized grains are shown at the edge of a deformation band. These strain-free, equiaxed grains contain only a small number of incoherent particles and no coherent particles. This region of the deformation band still possesses a high density of dislocations, while others examined showed evidence of polygonization. The 24 h treatment at 600°C carried these processes to completion, and the resulting microstructures in the transition and deformation bands were significantly different. The former were comprised of grains of 2.5/z average diameter containing only incoherent particles ranging in size from 500 to 1000 A that were spaced rather nonuniformly (Fig. 7 (a)). On the other hand, the deformation bands evolved a subgrain structure through polygonization, and the second phase remained in the form of coherent particles, though there was a measurable increase in their diameter (i.e., from 150 to 250 A). This coarsening is consistent with the Ostwald ripening effect* 0b* The growth of larger particles at the expense of smaller ones after the matrix has reached its equilibrium solute concentration.
(a)
(b) Fig. 7. Aged and swaged alloy annealed at 600°C for 24 h. (a) Recrystallized grains within a transition band. (× 10,500) (b) Coalescence of the coherent Co-rich phase and migration of a polygonized boundary within a deformation band. (× 24,000)
Mater. Sci. Eng., 1 (1966) 153-161
ANNEALING BEHAVIOR OF COLD-WORKED
Fig. 8. Aged and swaged alloy annealed at 700°C for 1 h. Migrating boundary pinned by closely spaced incoherent particles (marked "x"). ( x 48,000)
complicated by the change in cobalt solubility from 0.35~ at 600°C to about 0.80~ at 700°C1L This change upset the equilibrium developed during the original aging treatment and resulted in alteration of the size and volume fraction of the coherent phase. In some regions, sites of particle coalescence were still quite closely spaced, and migrating boundaries were often pinned once again. An example is found in Fig. 8, where it is also interesting to note that the advancing boundary segments between the particles appear to be of twin orientation.
4. DISCUSSION The inhibition of recovery and recrystallization in the presence of a fine dispersion of coherent particles in a C u - 2 ~ Co alloy has been shown. The direct observations indicate that both these processes did not become fully operative until the original aging temperature was approached; however, 1 h treatments at 500 and 600°C did not alter the structure enough to markedly reduce the workhardening. This temperature range for "softening" is substantially higher than would be expected based on matrix composition alone. According to Smart and Smith 24, 0.05}/0 Co in solution should increase the softening temperature of high-purity Cu by about 20°C. The unalloyed material, 50~o cold-worked, is half-recrystallized after 1 h at
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150°C. Thus, the maximum effect of 0.35~ Co would be to raise this temperature to just below 300°C. Recovery, defined as the rearrangement and annihilation of dislocations to produce strain-free subgrains zS, was confined to eliminating the dislocation structure in the deformation bands. It is suggested that additional thermal activation was required to assist dislocations in becoming disentangled from individual particles, as well as from themselves, and to migrate through regions containing many particles before coming to rest in lower energy configurations. The exact path followed is unknown, but from the evidence it is felt that one should consider the atomic mobility within the particles, as well as the matrix. Therefore, the activation energy for recovery, which is mainly a dislocation climb process as considered here, i.e., volume diffusion limited, should incorporate an average contribution from the dispersion of the Co-rich particles in addition to that for the copper matrix with 0.35~ Co in solution. Concurrent with the elimination of dislocation tangles was the reappearance of distinct particle strain contrast. The conditions for such contrast would be enhanced by the removal of dislocation strain fields. However, in addition, it is suggested that the particles also regain their original spherical shape. The particles are believed to be sheared into an elliptical shape as the result of plastic deformation 18. Such a variation would produce an increase in surface area without a volume change and thus represents an increase in the surface energy for each particle. With the latter as the driving force, the shape reversal ma~y be accomplished by shortrange intraparticle diffusion and would necessarily occur simultaneously with Ostwald ripening. Bailey and Hirsch z6 proposed that recrystallization is the growth of strain-free regions by means of boundary migration, where the boundaries may be the original grain boundaries, those produced by cold work or those resulting from polygonization. Their relative mobility depends on the angular misfit across them, as well as the type and magnitude of the driving force impelling them to move, Such potentially mobile boundaries in this alloy were pinned by the rather closely spaced particles, i.e., they were unable to pass through or bow between the particles. Boundary migration, first observed at 500°C, was involved with the coalescence of the coherent second phase. This freed the boundary since (a) it reduced the number Mater. Sci. Eng., 1 (1966) 153-161
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and increased the spacing of the pinning points and (b) it seemed to create adjacent particle-free zones. This behavior initiated i,n the transition bands with the cold work produced boundaries and was followed in the deformation bands once sharp boundaries were formed by polygonization. Recrystallization through the rotation and coalescence of subgrains as proposed by Hu 27 must be ruled out here, though it is conceivable that such behavior could have taken place during recovery. The question of the driving force for migration must now be considered. In the absence of a second phase, it is simply derived from differences in strain energy across existing interfaces within the coldworked structure or, if polygonization occurs first, from the surface energy of the sub-boundaries 2s. However, with particle coalescence playing such an important role, it is necessary to consider a chemical driving force as well. In this light, boundary migration in a two-phase alloy may be thought of as a variation of the discontinuous precipitation phenomena, i.e., a type of segregated growth involving a localized reaction at an advancing interface 29. Perhaps the most common mode of discontinuous precipitation has a cellular appearance and occurs in aging systems where grain boundary nucleation of the second phase is initially preferred. These particles tend to coarsen rapidly because of enhanced diffusion along the boundaries. A decrease in free energy results, creating a driving force for boundary migration and continued growth. Generally, the growth rate, or boundary migration rate, decreases as its chemical driving force is diminished by competing continuous precipitation 3°. This often results in the cessation of growth while the generalized reaction goes to completion. Analogous behavior may also occur in the overaged condition, in which case the growth at the boundary competes with Ostwald ripening of the continuous precipitate (e.o., Al-16~o Ag31). Kelly and Nicholson 1 suggest that discontinuous growth under these conditions should eventually consume the original structure, since the rate of free energy change for both coarsening processes is essentially constant but is higher for the one at the boundaries. It is this second mode of discontinuous growth that seems to be operative during recrystallization of the Cu-2~o Co alloy. Limited grain boundary precipitation would be expected to take place in Cu-Co alloys where the atomic volumes of solute and solvent differ by a rather small amount 32. This was confirmed by
Phillips, who observed only isolated evidence of discontinuous precipitation in a Cu-3~o Co alloy 33. None was found after the initial aging treatment in the alloy under study, where the undeformed specimens possessed very little grain boundary area. However, severe cold work introduced a great many new boundaries into the overaged structure (and more were produced by polygonization). Taking the number and spacing of the coherent particles into consideration, these boundaries could not avoid inheriting the particle distribution of the matrix and were thereby pinned. The boundaries remained pinned until there was sufficient thermal activation to promote particle coarsening and discontinuous growth at boundaries. The evidence in Fig. 5(b) gives a reasonable picture of the sequence of events in the boundary migration process. The direct observations made here were qualitative but imply that particle coalescence is the ratelimiting process for boundary migration. However, this is likely to be an incomplete picture since the observations did not provide enough information about the role of strain energy. A quantitative study is currently in progress utilizing a number of complementary experimental techniques that should sort out the relative importance of the above factors.
ACKNOWLEDGEMENTS
The authors wish to thank Professors D. Turnbull and M.F. Ashby of Harvard University for their stimulating discussions and constructive comments. Special thanks are due to Mr. J. A. Hugo and Dr. V. A. Phillips of General Electric Research Laboratory for helping to develop the technique of thinning wire specimens.
REFERENCES 1 A. KELLY AND R. B. NICHOLSON, Progr. Mater. Sci., 10 (1963) 149. 2 A. KELLY, Proc. Roy. Soc. (London), Ser. A, 282 (1964) 63. 3 0 . PRESTON AND N. J. GRANT, Trans. Met. Soc. AIME, 221 (1961) 164. 4 R. D. DOHERXY AND J. W. MARTIN, J. Inst. Metals, 91 (1962-3) 332. 5 R. D. DOHERTY AND J. W. MARTIN, Trans. Am. Soc. Metals, 57 (1964) 874. 6 R. D. DOHERTY AND J. W. MARTIN, Third European Re#, Conf. on Elect. Micr., Prague, 1964, p. 141.
Mater. Sci. Eno., 1 (1966) 153-161
ANNEALING BEHAVIOR OF COLD-WORKED C u - 2 ~ o C o ALLOY 7 W. C. LESLIE, J. T. MICHALAKAND F. W. AUL, Iron and Its Solid Solutions, Interscience Publ., New York, 1963, p. 119. 8 D. M. WILLIAMS,The Structure and Properties of Some Internally Oxidized Copper Alloys, Ph.D. Thesis, Cambridge Univ., 1964. 9 R. L. JONES, unpublished work (see ref. 2, p. 67). l0 K. T. AUST, 7e Colloque de Metallurgie, Saclay, Presses Universitaires de France, Paris, 1963, p. 97 (in French); General Electric Research Rept. 63-RL-3537M, Dec., 1963. 11 J. D. LIVINGSTON, Trans. Met. Soc. AIME, 215 (1959) 566. 12 J. A. HUGO AND V. A. PHILLIPS, Brit. J. Appl. Phys., 40 (1963) 202. 13 A. KNAPPWOST,Z. Physik. Chem., 12 (1957) 30. 14 V. A. PHILLIPS AND J. D. LIVINGSTON, Phil. Mag., 7 (1962) 969. 15 M. F. ASHBYAND L. M. BROWN, Phil. Mag., 8 (1963) 1083. 16 M. F. ASHBY AND R. EBLING, Trans. Met. Soc. AIME, in press. 17 L. G. BONAR, Precipitation Hardening, Ph.D. Thesis, Cambridge Univ., 1963. 18 J. D. LIVINGSTONAND J. J. BECKER,Trans. Met. Soc. AIME, 212 (1958) 316; N. TAMAGAWAAND T. MITUI,J. Phys. Soc. Japan., 20 (1965) 1989. 19 J. E. BAILEY, Phil. May., 8 (1963) 223.
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20 P. R. SWANN,Electron Microscopy and Strength of Crystals, lnterscience Publ., New York, 1963, p. 131. 21 C. S. BARRETT, Structure of Metals, McGraw-Hill Book Co., New York, 1952, pp. 442-446. 22 J. L. WALTERAND E. F. KOCH, Acta Met., 11 (1963) 923. 23 R. A. ORIANI,Acta. Met., 12 (1964) 1399; I. S. SERVI, Acta Met., 14 (1966) 234. 24 J. S. SMARTAND A. A. SMITH, Trans. Met. Soc. AIME, 147 (1942) 48. 25 J. C. M. LI, Recrystallization, Grain Growth and Textures, A S M Seminar, Detroit, 1965, in press. 26 J. E. BAILEYAND P. B. HIRSCH,Proc. Roy. Soc. (London), Ser. A, 267 0962) l l. 27 H. Hu, Recovery and Recrystallization of Metals, Interscience Publ., New York, 1963, p. 311. 28 J. W. CHRISTIAN,The Theory of Transformations in Metals and Alloys, Pergamon Press, Oxford, 1965, pp. 715-6. 29 J. W. CHRISTIAN, The Theory of Transformations in Metals and Alloys, Pergamon Press, Oxford, 1965, pp. 607-609, 643-649. 30 J. W. CAHN, Acta Met., 7 (1959) 18. 31 R. B. NICHOLSONAND J. NUTTING, Acta Met., 9 (1961) 332. 32 E. HORNBOGEN, Z. Metallk., 56 (1965) 133. 33 V. A. PHILLIPS, Trans. Met. Soc. AIME, 230 (1964) 967.
Mater. Sci. Eno., I (1966) 153-161
AND I. S. S E R V I
LedgemontLaboratory, Kennecott CopperCorporation, Lexington,Mass. (U.S.A.) (Received April 14, 1966)
S UMMAR Y A polycrystalline Cu-2°//o Co alloy, overaged at 600°C to produce a fine dispersion of coherent
particles, was severely dejbrmed by swagin9 and then annealed. The resulting microstructures were examined by electron transmission microscopy; and related mechanical properties were determined by tensile tests. Swaging plastically dejbrmed the matrix and the particles, producing a substructure comprised of "deJbrmation" bands containing a high density of tangled dislocations and diffuse cells separated by narrow "transition" bands of well-defined subgrains. Recovery and recrystallization processes were inhibited by the presence of the second phase and did
not become Jully operative until annealing took place in the vicinity of the original aging temperature, i.e., 500°C and above. Recovery resulted in the polygonization of the deformation bands and the restoration of spherical shape to the coherent particles within these regions. Recrystallization.proceeded by means of boundary migration, initially in the transition bands utilizing the cold-work produced boundaries and then in the deformation bands with the poly9onized boundaries. The migration process involved the coalescence of the coherent phase and appeared to be a variation of the discontinuous precipitation phenomena.
RF.SUMF. Dans un alliage polycristallin de cuivre d 2°//0 de cobalt, une fine dispersion de particules cohOrentes a OtO obtenue par un traitement de survieillissement d 600°C. Le m~tal a ensuite ~tO soumis gt une dOformation intense par dtampage et Ot~ recuit. Les microstructures obtenues gl diff~rents stades ont Otk examinOes par microscopie ~lectronique en transmission et les caractOristiques mdcaniques correspondantes ont ~t~ dOtermin~es au moyen d'essais de traction. Uop~ration dYtampage a pour effet de d~Jbrmer plastiquement dr.la ibis la matrice et les particules et donne naissance dune structure de bandes : bandes de "d~Jbrmation" comportant des zones gt densit~ ~lev~e de dislocations enchev~tr~es et des cellules plus diJJi~ses, bandes de "transition" plus Otroites constituOes par des sous-grains bien dOfinis et sOparant les bandes du premier type. Les phdnombnes de restauration et de recristallisation sont entrav~s ou retardOs par les particules de la seconde phase: ce n'est que Materials Science and Engineering-
lorsque la temperature de recuit est voisine de la tempdrature du vieillissement initial, c'est-~t-dire 5 0 0 ° C ou au-dessus, que ces processus deviennent r~ellement efficaces. La restauration se traduit par une polygonisation des bandes de d~Jbrmation et le retour d la Jorme sph~rique des particules coh~rentes situ~es dans ces r~gions. La recristallisation s' effectue par migration des joints: au ddbut ce sont les joints de la structure d'~crouissage, situds dans les bandes de transition, qui se d~placent, puis la recristallisation se poursuit par mouvement des sous-joints de polygonisation qui se sont form,s dans les bandes de d~Jbrmarion. La migration des joints est accompagn~e d'une coalescence des particules de la phase coh~rente, si bien que ce mdcanisme de recristallisation semble constituer une variante des processus de precipitation discontinue. * Presented in part at the Fall Meeting of the Metallurgical Society of AIME, October 19, 1965, Detroit, Mich., U.S.A.
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Z USAMMENFASSUNG Eine bei 600°C zur Erzeugung einer feinen Verteilung kohiirenter Teilchen angelassene polykristalline Cu2% Co-Legierung wurde durch Hiimmern stark verformt und dann angelassen. Mit Hilfe yon elektronenmikroskopischen Durchstrahlungsaufnahmen wurde die entstandene Mikrostruktur untersucht. Zugeh6rige mechanische Eigenschajten wurden dutch Zugversuche bestimmt. Matrix und Ausscheidungen wurden durch Hiimmern plastisch verjormt. Die erzeugte Substruktur umJaflte "Verformungsb?inder" mit einer hohen Dichte yon Versetzungsnetzwerken und verteilte, dutch enge "t)bergangsbiinder" aus wohldefinierten SubkiJrnern getrennte Zellen. Erholungs- und Rekristallisations-
1. INTRODUCTION The strengthening of metals by the introduction of a fine dispersion of a second phase is dependent upon the ability of the particles to inhibit dislocation movement under applied stress ~. A large variety of dispersions can be produced by the decomposition of supersaturated solid solutions, internal oxidation of alloys, powder metallurgy techniques, etc. These dispersions, in the form of either coherent or incoherent particles, may consist of metallic phases having mechanical properties not too different from those of the matrix or of very hard, insoluble ceramic compounds. A unique characteristic of these alloys after plastic deformation is that their work-hardened state often remains stable to relatively high temperatures 2. That is, depending on the type, size, and volume fraction of the second phase, there may be a definite resistance to recovery and recrystallization processes at temperatures where such behavior would normally occur in single-phase alloys of comparable solid solution composition. In fact, in the case of certain internally-oxidized alloys, structural stability can approach the melting point of the matrix 3. While these effects have been known for some time, it is only recently, with the advances in electron transmission microscopy techniques, that detailed studies of the variables involved in these processes have been undertaken. Doherty and Martin 4-6 investigated the annealing behavior of deformed A1-Cu alloys having dispersions of the incoherent 0 phase (Cu2A1). These investigators
prozesse waren durch die Anwesenheit der zweiten Phase gehemmt und wurden nicht voll wirksam bis zum Anlassen in der Niihe der urspriinglichen Anlafltemperatur, also bei 500°C oder hiiher. Die Erholung ffihrte zur Polygonisation der Verformungsbiinder und zur Wiederherstellung der Kugelgestalt der kohiirenten Teilchen innerhalb dieser Gebiete. Die Rekristallisation erjolgte durch Korngrenzenwanderung, und zwar anfdnglich in den Obergangsb?indern unter Ausnutzung der durch die KaltverJormung erzeugten Korngrenzen und dann in den VerJbrmungsbiindern mit den polygonisierten Grenzen. Der Wanderungsprozefl war mit der Zusammenlagerung der kohiirenten Phase verbunden und schien eine Variation der Erscheinungen bei diskontinuierlicher Ausscheidung zu sein.
found that when particles are spaced less than 1.0# apart there is a measurable reduction in the rate of "nucleation" of strain-free regions through the inhibition of low-angle boundary migration. Similar observations were made by Leslie et al. in Fe-Cu alloys 7 and by Williams for a number of internallyoxidized Cu-base alloys a. In still another investigation, the mechanical properties of a Cu-Co alloy, aged to produce a fine dispersion of coherent particles and then deformed, show its work-hardened state to be retained after each of a series of 1 h annealing treatments at temperatures up to 600°C 9. However, related structural observations were not reported. Thus, the current study utilizing electron transmission microscopy was undertaken with a comparable alloy treated in a similar manner in order to determine how the coherent second phase affects recovery and recrystallization. The results will be discussed in light of the current models proposed for these phenomena in pure metals and singlephase alloys ~°.
2. EXPERIMENTAL PROCEDURE
A Cu-Co alloy of nominal composition 2.0 wt. ~o Co was prepared from 99.999 + ~o pure copper and 99.84~o pure electrolytic cobalt. The Czochralski method for solidification was used and proved to give minimum microsegregation. Chemical and spectrographic analyses indicated that the cobalt content was 2.05 wt. ~ and that impurities consisted of: Mater. Sci. Eng., 1 (1966) 153-161
Cu-2~o Co
ANNEALING BEHAVIOR OF COLD-WORKED
< 100 p.p.m. Ni < 10 p.p.m. Mg, Fe < 1 p.p.m. AI, Ca 10 p.p.m. C The ingot was cold-swaged to wire of 0.040 in. diam. without intermediate annealing. The wire was then solution-treated in vacuo at 1000°C for ¼h followed by iced brine quenching. Aging was carried out at 600°C for 6 h in an inert atmosphere. According to Livingston 11 the particle size and spacing resulting from this treatment gives optimum mechanical properties. A portion of the stock was cold-swaged to a reduction in area of approximately 50~, and individual samples were annealed as follows: Temperature(°C): Time(h):
70
i
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155
ALLOY
i
~ deformed [ 60 _--aged
P
e ~ ~ . _ _
=
0.2 % Y.S.
50
d ~ 4o p--.q w x 50 u_
12 *~ _o
as-aged
20
0
IO N
d
/aged a deformed --
O RT
I
t
1(30 200
~
A
.
I
I
I
I
300
400
500
600
ANNEALING TEMPERATURE (*C)
Fig. 1. Room-temperature tensile properties of Cu-2 ~o Co in the aged, deformed and annealed conditions. The alloy was aged at 600°C for 6 h. Annealing treatments followed cold-swaging to a 50 ~o reduction in area and were for 1 h at the indicated temperatures.
300 400 500 600 700 1 1 1 1, 24 1
Specimens given the 1 h treatments at 300-600°C were tested in tension on an Instron machine to determine yield and ultimate strength, as well as tensile elongation. A portion of each wire was examined in a Hitachi HU-11A electron microscope equipped with a biaxial tilting specimen stage. An accelerating voltage of 100 kV was used. The wires were first mechanically polished parallel to their axis to a thickness of about 0.005 in. Thin foils were prepared by electrothinning utilizing a fine jet technique developed by Hugo and Phillips i2. The electrolyte was a solution of two parts phosphoric acid to one part water.
3.
EXPERIMENTAL RESULTS
The tensile data obtained from the aged C u - 2 ~ Co alloy after the mechanical and thermal treatments are plotted in Fig. 1. These results are in substantial agreement with those reported by Jones 9, showing only small changes from the cold-worked condition. These changes consist of a slight loss of strength, as well as a small, but distinct, increase in ductility with increasing, annealing temperature. Subtle as these may be, they indicate that some structural alterations occurred during the thermal treatments. The microstructure of the aged and undeformed condition is shown in Fig. 2. The cobalt concentration remaining in solid solution is expected to be about 0.35~ 11. This was verified by electrical resistivity measurements (i.e., 4/~f~-cm at room temperature corresponds to 0.35~ Co13). The micro-
Fig. 2. Solution treated at 1000°C for ¼ h, iced-brine quenched and aged at 600°C for 6 h. (z48,000)
graph exhibits a dispersion of coherent particles within a single grain*. The spherical Co-rich particles are revealed by diffraction contrast arising from elastic strains in the matrix due to lattice mismatch between the two phases TM15. The size of the particles was determined from micrographs by taking the length of the line of no contrast as their diameteP 4. An average value of 150 A was observed and may be considered to be accurate to 25~. The foil thickness was assumed to be 1000 A, but in regions examined it may have departed from this estimate by a factor of two. Particle density was * The average grain diameter of the wires was of the order of 0.2 mm.
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Fig. 3. Aged and deformed in tension to 11~o elongation. ( x 48,000)
determined, and then interparticle spacing was calculated in a manner proposed by Ashby and Ebling 16 assuming the particles to be randomly dispersed. A value for the latter of 1100 A was obtained which is 2 5 ~ smaller than that reported by Bonar 17 for an alloy of similar composition and treatment. Interpretation of the deformed microstructure was particularly difficult because of the poor visibility of individual dislocations. The result of 11~ tensile elongation is seen in Fig. 3, and comparison with Fig. 2 indicates that plastic flow produced substantial alteration in particle strain contrast, as well as irregular variations in the orientation of the lines of no contrast. Deformation is believed to change the shape of the particles while coherency is retained is and could account for the former. The latter would be expected from a random distribution of local lattice tilts and thus suggests that the dislocations are in a tangled array rather than in a well-defined cell structure of the type characteristic
Fig. 4. Aged and cold-swaged to a 50~o reduction in area. Banded structure elongated in the direction of the wire axis. Deformation bands "A" separated by transition bands "B". (x 25,000)
Mater. Sci. Eng., 1 (1966) 153-161
ANNEALING BEHAVIOR OF COLD-WORKED
of deformed unalloyed Cu 19. The inhibition of cell formation is consistent with Swann's observation that a fine dispersion of a ductile second phase tends to confine dislocation movement to the primary slip systems and inhibits cross-slip in the initial stages of deformation 2°. Another source for changes in particle strain contrast, which would also account for the difficulties in dislocation resolution, is the superposition of the elastic strain fields of coherent particles and dislocations. This should result in rather complicated contrast that would vary from area to area. More severe deformation, to be discussed below, tends to delineate regions having large variations in dislocation density, but contrast never approached the degree of sharpness observed in most worked single-phase materials. Swaging is a notably complicated and nonuniform method for working metals. It promotes glide on many intersecting slip systems and produces a complex combination of lattice rotations and tilts which develop the elements of a fiber texture 21. The resulting banded structure, seen in Fig. 4, accommodates the. lattice reorientations and is similar to that found in heavily-rolled pure metals and alloys 1°. The regions marked "A" appear to have a high density of dislocations in both relatively uniform tangles and poorly defined cells but also include some sharp boundaries. These regions are separated by narrow bands (designated "B") that are comprised of small, elongated subgrains bounded by discrete subgrain walls. Dislocation density seems to be considerably lower within this substructure. Close examination reveals the coherent particles, but their contrast is markedly altered by the deformation. Selected area diffraction indicated only small orientation variations across the tangled regions (i.e., <<,2°), while there was generally about an 8-10 ° disorientation" across the sharp bands. Furthermore, the orientation changes in these bands were in incremental steps of 1-2 ° from one subgrain to the next. The sharp boundaries are the product of glide polygonization made possible by interactions of dislocations moving on the intersecting slip planes. In a qualitative sense the more tangled regions of this structure are analogous to the "deformation" bands, and the narrow subgrain regions correspond to the "transition" bands described by Walter and Koch 22 for rolled materials. These designations will be adopted here. Only small variations in the dislocation substructure could be observed after heat treatments at 300 and 400°C. In the main, these were confined
Cu-2~
C o ALLOY
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(a)
Fig. 5. Aged and swaged alloy annealed at 500°C for 1 h. (a) Dislocations rearranged to form a distinct cell structure within a deformation band. ( × 28,000) (b) Coalescence of the coherent Co-rich phase and migration of a cold-work produced boundary within a transition band. ( × 40,000)
to local rearrangements which tended to increase and sharpen the cell structure in the deformation bands. Such rearrangements were clearly evident after the 500°C treatment as seen in Fig. 5(a). However, a more dramatic change in microstructure was found in isolated regions within the transition bands. In Fig. 5 (b), sub-boundary A'B'C' is pinned by a few coarse incoherent particles (y)*. It has migrated from position ABC, which consists of another string of incoherent particles (x), and the included area between ABC and A'B'C' is free from both the coherent phase and dislocations. * The particles are no longer in the foil, having been lost during electrothinning, but are identified by the remaining holes which appear white in the micrographs.
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served in this alloy system 23. Subsequent migration of the new sub-boundaries, in the manner described above, eliminated a portion of the coherent phase through particle coalescence (Fig. 7(b)). Annealing at 700°C greatly accelerated polygonization, boundary migration, and particle coalescence to the point where delineation of separate deformation and transition band effects was rather difficult. In addition, the situation was further
Fig. 6. Aged and swaged alloy annealed at 600°C for 1 h. Recrystallized grains at the edge of a transition band growing into a dislocation-dense region of a deformation band. (x 48,000)
In addition, some narrow regions adjacent to the leading side of the advancing interface also appear to be free from coherent particles (e.g., between A' and B'). The above process of particle coalescence and boundary migration is considerably more advanced after the 1 h treatment at 600°C. In Fig. 6, a number of recrystallized grains are shown at the edge of a deformation band. These strain-free, equiaxed grains contain only a small number of incoherent particles and no coherent particles. This region of the deformation band still possesses a high density of dislocations, while others examined showed evidence of polygonization. The 24 h treatment at 600°C carried these processes to completion, and the resulting microstructures in the transition and deformation bands were significantly different. The former were comprised of grains of 2.5/z average diameter containing only incoherent particles ranging in size from 500 to 1000 A that were spaced rather nonuniformly (Fig. 7 (a)). On the other hand, the deformation bands evolved a subgrain structure through polygonization, and the second phase remained in the form of coherent particles, though there was a measurable increase in their diameter (i.e., from 150 to 250 A). This coarsening is consistent with the Ostwald ripening effect* 0b* The growth of larger particles at the expense of smaller ones after the matrix has reached its equilibrium solute concentration.
(a)
(b) Fig. 7. Aged and swaged alloy annealed at 600°C for 24 h. (a) Recrystallized grains within a transition band. (× 10,500) (b) Coalescence of the coherent Co-rich phase and migration of a polygonized boundary within a deformation band. (× 24,000)
Mater. Sci. Eng., 1 (1966) 153-161
ANNEALING BEHAVIOR OF COLD-WORKED
Fig. 8. Aged and swaged alloy annealed at 700°C for 1 h. Migrating boundary pinned by closely spaced incoherent particles (marked "x"). ( x 48,000)
complicated by the change in cobalt solubility from 0.35~ at 600°C to about 0.80~ at 700°C1L This change upset the equilibrium developed during the original aging treatment and resulted in alteration of the size and volume fraction of the coherent phase. In some regions, sites of particle coalescence were still quite closely spaced, and migrating boundaries were often pinned once again. An example is found in Fig. 8, where it is also interesting to note that the advancing boundary segments between the particles appear to be of twin orientation.
4. DISCUSSION The inhibition of recovery and recrystallization in the presence of a fine dispersion of coherent particles in a C u - 2 ~ Co alloy has been shown. The direct observations indicate that both these processes did not become fully operative until the original aging temperature was approached; however, 1 h treatments at 500 and 600°C did not alter the structure enough to markedly reduce the workhardening. This temperature range for "softening" is substantially higher than would be expected based on matrix composition alone. According to Smart and Smith 24, 0.05}/0 Co in solution should increase the softening temperature of high-purity Cu by about 20°C. The unalloyed material, 50~o cold-worked, is half-recrystallized after 1 h at
Cu-2~
C o ALLOY
159
150°C. Thus, the maximum effect of 0.35~ Co would be to raise this temperature to just below 300°C. Recovery, defined as the rearrangement and annihilation of dislocations to produce strain-free subgrains zS, was confined to eliminating the dislocation structure in the deformation bands. It is suggested that additional thermal activation was required to assist dislocations in becoming disentangled from individual particles, as well as from themselves, and to migrate through regions containing many particles before coming to rest in lower energy configurations. The exact path followed is unknown, but from the evidence it is felt that one should consider the atomic mobility within the particles, as well as the matrix. Therefore, the activation energy for recovery, which is mainly a dislocation climb process as considered here, i.e., volume diffusion limited, should incorporate an average contribution from the dispersion of the Co-rich particles in addition to that for the copper matrix with 0.35~ Co in solution. Concurrent with the elimination of dislocation tangles was the reappearance of distinct particle strain contrast. The conditions for such contrast would be enhanced by the removal of dislocation strain fields. However, in addition, it is suggested that the particles also regain their original spherical shape. The particles are believed to be sheared into an elliptical shape as the result of plastic deformation 18. Such a variation would produce an increase in surface area without a volume change and thus represents an increase in the surface energy for each particle. With the latter as the driving force, the shape reversal ma~y be accomplished by shortrange intraparticle diffusion and would necessarily occur simultaneously with Ostwald ripening. Bailey and Hirsch z6 proposed that recrystallization is the growth of strain-free regions by means of boundary migration, where the boundaries may be the original grain boundaries, those produced by cold work or those resulting from polygonization. Their relative mobility depends on the angular misfit across them, as well as the type and magnitude of the driving force impelling them to move, Such potentially mobile boundaries in this alloy were pinned by the rather closely spaced particles, i.e., they were unable to pass through or bow between the particles. Boundary migration, first observed at 500°C, was involved with the coalescence of the coherent second phase. This freed the boundary since (a) it reduced the number Mater. Sci. Eng., 1 (1966) 153-161
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and increased the spacing of the pinning points and (b) it seemed to create adjacent particle-free zones. This behavior initiated i,n the transition bands with the cold work produced boundaries and was followed in the deformation bands once sharp boundaries were formed by polygonization. Recrystallization through the rotation and coalescence of subgrains as proposed by Hu 27 must be ruled out here, though it is conceivable that such behavior could have taken place during recovery. The question of the driving force for migration must now be considered. In the absence of a second phase, it is simply derived from differences in strain energy across existing interfaces within the coldworked structure or, if polygonization occurs first, from the surface energy of the sub-boundaries 2s. However, with particle coalescence playing such an important role, it is necessary to consider a chemical driving force as well. In this light, boundary migration in a two-phase alloy may be thought of as a variation of the discontinuous precipitation phenomena, i.e., a type of segregated growth involving a localized reaction at an advancing interface 29. Perhaps the most common mode of discontinuous precipitation has a cellular appearance and occurs in aging systems where grain boundary nucleation of the second phase is initially preferred. These particles tend to coarsen rapidly because of enhanced diffusion along the boundaries. A decrease in free energy results, creating a driving force for boundary migration and continued growth. Generally, the growth rate, or boundary migration rate, decreases as its chemical driving force is diminished by competing continuous precipitation 3°. This often results in the cessation of growth while the generalized reaction goes to completion. Analogous behavior may also occur in the overaged condition, in which case the growth at the boundary competes with Ostwald ripening of the continuous precipitate (e.o., Al-16~o Ag31). Kelly and Nicholson 1 suggest that discontinuous growth under these conditions should eventually consume the original structure, since the rate of free energy change for both coarsening processes is essentially constant but is higher for the one at the boundaries. It is this second mode of discontinuous growth that seems to be operative during recrystallization of the Cu-2~o Co alloy. Limited grain boundary precipitation would be expected to take place in Cu-Co alloys where the atomic volumes of solute and solvent differ by a rather small amount 32. This was confirmed by
Phillips, who observed only isolated evidence of discontinuous precipitation in a Cu-3~o Co alloy 33. None was found after the initial aging treatment in the alloy under study, where the undeformed specimens possessed very little grain boundary area. However, severe cold work introduced a great many new boundaries into the overaged structure (and more were produced by polygonization). Taking the number and spacing of the coherent particles into consideration, these boundaries could not avoid inheriting the particle distribution of the matrix and were thereby pinned. The boundaries remained pinned until there was sufficient thermal activation to promote particle coarsening and discontinuous growth at boundaries. The evidence in Fig. 5(b) gives a reasonable picture of the sequence of events in the boundary migration process. The direct observations made here were qualitative but imply that particle coalescence is the ratelimiting process for boundary migration. However, this is likely to be an incomplete picture since the observations did not provide enough information about the role of strain energy. A quantitative study is currently in progress utilizing a number of complementary experimental techniques that should sort out the relative importance of the above factors.
ACKNOWLEDGEMENTS
The authors wish to thank Professors D. Turnbull and M.F. Ashby of Harvard University for their stimulating discussions and constructive comments. Special thanks are due to Mr. J. A. Hugo and Dr. V. A. Phillips of General Electric Research Laboratory for helping to develop the technique of thinning wire specimens.
REFERENCES 1 A. KELLY AND R. B. NICHOLSON, Progr. Mater. Sci., 10 (1963) 149. 2 A. KELLY, Proc. Roy. Soc. (London), Ser. A, 282 (1964) 63. 3 0 . PRESTON AND N. J. GRANT, Trans. Met. Soc. AIME, 221 (1961) 164. 4 R. D. DOHERXY AND J. W. MARTIN, J. Inst. Metals, 91 (1962-3) 332. 5 R. D. DOHERTY AND J. W. MARTIN, Trans. Am. Soc. Metals, 57 (1964) 874. 6 R. D. DOHERTY AND J. W. MARTIN, Third European Re#, Conf. on Elect. Micr., Prague, 1964, p. 141.
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ANNEALING BEHAVIOR OF COLD-WORKED C u - 2 ~ o C o ALLOY 7 W. C. LESLIE, J. T. MICHALAKAND F. W. AUL, Iron and Its Solid Solutions, Interscience Publ., New York, 1963, p. 119. 8 D. M. WILLIAMS,The Structure and Properties of Some Internally Oxidized Copper Alloys, Ph.D. Thesis, Cambridge Univ., 1964. 9 R. L. JONES, unpublished work (see ref. 2, p. 67). l0 K. T. AUST, 7e Colloque de Metallurgie, Saclay, Presses Universitaires de France, Paris, 1963, p. 97 (in French); General Electric Research Rept. 63-RL-3537M, Dec., 1963. 11 J. D. LIVINGSTON, Trans. Met. Soc. AIME, 215 (1959) 566. 12 J. A. HUGO AND V. A. PHILLIPS, Brit. J. Appl. Phys., 40 (1963) 202. 13 A. KNAPPWOST,Z. Physik. Chem., 12 (1957) 30. 14 V. A. PHILLIPS AND J. D. LIVINGSTON, Phil. Mag., 7 (1962) 969. 15 M. F. ASHBYAND L. M. BROWN, Phil. Mag., 8 (1963) 1083. 16 M. F. ASHBY AND R. EBLING, Trans. Met. Soc. AIME, in press. 17 L. G. BONAR, Precipitation Hardening, Ph.D. Thesis, Cambridge Univ., 1963. 18 J. D. LIVINGSTONAND J. J. BECKER,Trans. Met. Soc. AIME, 212 (1958) 316; N. TAMAGAWAAND T. MITUI,J. Phys. Soc. Japan., 20 (1965) 1989. 19 J. E. BAILEY, Phil. May., 8 (1963) 223.
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20 P. R. SWANN,Electron Microscopy and Strength of Crystals, lnterscience Publ., New York, 1963, p. 131. 21 C. S. BARRETT, Structure of Metals, McGraw-Hill Book Co., New York, 1952, pp. 442-446. 22 J. L. WALTERAND E. F. KOCH, Acta Met., 11 (1963) 923. 23 R. A. ORIANI,Acta. Met., 12 (1964) 1399; I. S. SERVI, Acta Met., 14 (1966) 234. 24 J. S. SMARTAND A. A. SMITH, Trans. Met. Soc. AIME, 147 (1942) 48. 25 J. C. M. LI, Recrystallization, Grain Growth and Textures, A S M Seminar, Detroit, 1965, in press. 26 J. E. BAILEYAND P. B. HIRSCH,Proc. Roy. Soc. (London), Ser. A, 267 0962) l l. 27 H. Hu, Recovery and Recrystallization of Metals, Interscience Publ., New York, 1963, p. 311. 28 J. W. CHRISTIAN,The Theory of Transformations in Metals and Alloys, Pergamon Press, Oxford, 1965, pp. 715-6. 29 J. W. CHRISTIAN, The Theory of Transformations in Metals and Alloys, Pergamon Press, Oxford, 1965, pp. 607-609, 643-649. 30 J. W. CAHN, Acta Met., 7 (1959) 18. 31 R. B. NICHOLSONAND J. NUTTING, Acta Met., 9 (1961) 332. 32 E. HORNBOGEN, Z. Metallk., 56 (1965) 133. 33 V. A. PHILLIPS, Trans. Met. Soc. AIME, 230 (1964) 967.
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