Studies of the initiation, growth and dissolution of the discontinuous precipitation product in aluminumzinc alloys

Acta metall. Vol. 32, No. 10, pp. 1709-1717, 1984

0001-6160/84 $3.00+0.00 Copyright © 1984 Pergamon Press Ltd

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STUDIES OF THE INITIATION, GROWTH A N D DISSOLUTION OF THE DISCONTINUOUS PRECIPITATION PRODUCT IN A L U M I N U M - Z I N C ALLOYS I. G. S O L O R Z A N O I t , G. R. P U R D Y l and G. C. WEATHERLY 2 ~Department of Metallurgy and Materials Science, McMaster University, Hamilton, Ontario, Canada L8S 4L7 and 2Department of Metallurgy and Materials Science, University of Toronto, Toronto, Ontario, Canada M5S 1A4 (Received 29 March 1984)

Abstract--The discontinuous mode of precipitation is studied in A1-22 at.% Zn. High resolution STEM X-ray microanalysis has been used to measure the composition profiles in individual depleted lamellae formed under isothermal conditions. These data, combined with local values of cell velocity and interlamellar spacing, have been used to evaluate interfacial diffusion coefficients and residual supersaturations. The results are compared with predictions of existing theories. It is demonstrated that initiation involves diffusion induced grain boundary displacements and that during the growth process, a significant amount of free energy is stored in the product phases. The dissolution of the reheated lamellar structure is volume diffusion-controlled. Rrsum&---Nous 6tudions la prrcipitation discontinue dans A1-22 at.% Zn. Nous avons utilis6 la technique de la microanalyse de rayons X sur un microscope fi balayage en transmission ~t haute rrsolution afin de mesurer les profils de composition dans des lamelles individuelles appauvries formres dans des conditions isothermes. On a utilis6 ces donnres, ainsi que des valeurs locales de la vitesse des cellules et de l'espacement interlamellaire, pour 6valuer les coefficients de diffusion intarfaciale et les sursaturations rrsiduelles. On compare les rrsultats avec les prrvisions des throries existantes. On drmontre que l'initiation implique des drplacements de joints de grains induits par diffusion et qu'au cours de la croissance une quantit6 notable d'rnergie libre est emmagasinre dans les phases produites. La dissolution de la structure lamellaire rrchauffre est contrrlre par la diffusion en volume. Zusammenfassung--Der diskontinuierliche Ausscheidungsprozel3 wurde an der Legierung AI-22 at.% Zn untersucht. Die Profile der Zusammensetzungen wurden an einzelnen, unter isothermen Bedingungen entstandenen und verarmten Lamellen mit hochaufl6sender R6ntgenanalyse im STEM ermittelt. Mit diesen Ergebnissen und mit lokalen Werten fiir Zellgeschwindigkeit und interlamellaren Abstand wurden die Koeffizienten fiir Grenzflfichendiffusion und die Restfibers/ittigungen bestimmt. Vorhandene Theorien werden mit diesen Ergebnissen verglichen. Der Beginn der Ausscheidungsprozesse 1/iuft unter diffusionsinduzierter Korngrenzverschiebung ab, W/ihrend des Wachstums wird ein betr/ichtlicher Anteil der freien Energie in den entstehenden Phasen gespeichert. Das Aufl6sen der nochmals aufgeheizten Lamellenstruktur wird durch die Volumdiffusion kontrolliert.

INTRODUCTION The discontinuous precipitation reactions in A1-Zn have been the subject of numerous studies [1-4]. They remain of current interest, in part because so much is already known about the A1-Zn system; the solution thermodynamics, interfacial energies of the coprecipitating phases [5, 6] and the precipitate morphology and occurrence are all documented to some degree. Thus, the system has become something of a reference or textbook case, within which current theoretical ideas may be tested and against which other systems may be compared. The purpose of the present contribution is to explore the microchemical and microstructural details of the transformation, its

tPresent address: Departamento de Ciencia dos Materiais e Metalurgia, PUC/RJ, C.P. 38008 Gavea, Rio de Janeiro, RJ, 22452, Brazil.

initiation, steady growth and reversion, using the higher microanalytical resolution afforded by the dedicated scanning transmission electron microscopy (STEM) and to make comparisons with theory wherever these seem feasible and worthwhile. EXPERIMENTAL A series of A1-Zn alloys was prepared from high purity starting materials by induction melting under purified argon and chill casting into 10 m m diameter rods. The ingots were sealed in argon and homogenized at 450°C for 500 h. Slabs about 1 m m thick were checked for segregation in the electron microprobe analyser, then cold rolled to 0.1 mm. The alloys had the following compositions: 6.8, 9.5, 12.2, 22.1 and 28.4at.% Zn ( + 0 . 5 at.%Zn). All of the alloys were used for X-ray microanalysis calibration, the last two for studies of discontinuous precipitation. Only the

1709

1710

SOLORZANO et al.: GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION

22 at.~o Zn alloy was employed for analytical electron microscopical studies of discontinuous precipitation products. The cold rolled strips were solution treated at 370°C to obtain a recrystallized microstructure, grain size about 150 #m. The alloys were quenched, then aged at temperatures from 110 to 220°C. For dissolution studies, the aging treatments were followed by isothermal treatments just above the solvus temperature for the alloy. After thinning in 30~ nitric acid-methanol at - 3 0 ° C , the specimens were observed in a Phillips EM300 electron microscope and in a Vacuum Generators HB-5 scanning transmission microscope. Average growth rates were determined from the slopes of average cell width (at least 10 measurements) vs time. Average as well as local values of the spacing were measured in the STEM. In order to measure the composition profiles in the depleted solid solution, energy dispersive X-ray analysis was employed. As noted above, the series of A1-Zn alloys, quenched from homogenizing temperature, served to calibrate the STEM energy dispersive spectrometer, i.e. to determine the Cliff-Lorimer factor kA1-Zn [7], which relates composition to intensity ratios of characteristic spectra, in the approximation that the foil is sufficiently thin that absorption and fluorescence are negligible. The calibration process and the resuiting values of kA1-Zn are given in the Appendix.

by Bouzaher and Simon [8]. Microanalysis of the swept areas then confirmed that the boundary acted as a diffusion path for solute. The swept regions were uniformly depleted of zinc, as shown in Fig. 2. It was noted that the spacing between precipitates was about one order of magnitude greater than the final steady spacings adopted by the discontinuous product in these alloys. In order to study the bowing event more fully, the 12~o Zn alloy was held at 428K for a few seconds. Bowing events similar to those in the 22~o Zn material were observed, although this alloy does not undergo discontinuous precipitation at the temperature chosen. Thus, in this system, grain boundary bowing between initial precipitates appears to be a more general phenomenon than discontinuous precipitation.

(b) Growth of the discontinuous product The morphological aspects of the reaction product will be presented first, followed by results of the microanalyses. From observations of the structure of entire colonies, such as that shown in Fig. 3, we note that the steady parallel growth of/3 lamellae develops rather slowly, after a period of less cooperative growth from the initial array of grain boundary precipitates. The occurrence of both single and double seam morphologies was observed, but no

RESULTS

(a) Initiation of the reaction The speed of the discontinuous reaction in AI-22 at.%o Zn is such that the earliest stages were not readily observable. On occasion, however, the bowing of the grain boundary between pinning precipitates of the hexagonal zinc-rich 13 phase was observed, Fig. 1, in a manner similar to that reported

(a)

50

I

r~ 20

I e--O-o t

0

I

t I-o-e-

J

J

~

I

It

0

200

600

I

--

Z(nrn) (b)

Fig. 1. Bowing of the grain boundary between Zn-rich precipitates; A1-22 at.~ Zn, held for 300 s at 428K.

Fig, 2. (a) A grain boundary in A1-22~ Zn, 300 s at 428K. (b) Compositional STEM scan across the region of (a).

SOL6RZANO et al.:

GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION

Fig. 3. A well-formed colony. Annular dark-field STEM micrograph. Insect, selected area microdiffraction patterns from product and parent grains. A1-22~ Zn held for 300 s at 478K. strong correlation with reaction temperature, of the type suggested by Baumann e t al. [9] was found. Instead, double and single seams were observed about equally at each temperature. Figure 4 shows a well-developed double seam, while Fig. 5 suggests the development of this morphology via the "S" mechanism. The steady growth morphology, characterized by constant spacing of lamellae, was studied in some detail. The morphology of the quenched growth front is shown in Figs 6 and 7. One important and consistent feature of the quenched growth morphologies is the bowing of the grain boundary into the supersaturated matrix. Bending back of the grain boundary was seen only on occasion, Fig. 8, and only in regions where it could be argued that the lamellae were on the point of changing their spacing, perhaps

Fig. 4. TEM image of a typical double seam in A1-22%Zn, 300 s at 428K.

171!

Fig. 5. Development of discontinuous colonies at both sides of a grain boundary, A1-22~/oZn, 200 s at 478K.

Fig. 6. Transmission micrograph, showing the morphology of the quenched discontinuous precipitation front.

Fig. 7. STEM bright field image of a front.

1712

S O L O R Z A N O et al.:

GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION

Fig. 8. An %/ct' interface between widely spaced fl lamellae.

z = -100

nm

z=0

z = 100 nm

z:0

z:8Onm

(a) 11.50 1110

428 K S = 240 nm v ~ 4 0 . 3 n m . s -I

•

1070

Fig. 9. A regularly spaced colony, A I - 2 2 ~ Zn, 300 s at 478K.

0:5.0

9 90

9 so

I/~

91o

by branching or nucleating a new precipitate at the front. Figure 9 shows a colony of parallel lamellae, which have developed a constant spacing for a sufficiently large distance that the morphology can be termed "steady state". In cases like these it must be argued that the transformation takes place under relatively constant driving force, even though there is clear evidence of a well-advanced general precipitation reaction in the matrix ahead of the front, Figs 6 and 7. The coarsened matrix microstructure permitted an estimate of the minimum concentration to be made using

•

830 790

•

7.50

[I -100

I -80

I -60

I -40

I

I 0

-20

1 20

1 f2 ~3 4 5 f6

k7 8 9

I 60

I 80

h 100

Distence (nm)

(b) Fig. 10. Al-22~o Zn, 300 s at 428K. (a) The microstructure, with sets of convergent beam microdiffraction patterns from microanalysis positions in each of the two central ct' lamellae. The large precipitates on the grain boundary formed after prolonged observation in the STEM. (b) Compositional profile across the left central ct' lamella of Fig. lO(a).

Table 1. Details of cells studied at 428K aging temperature Lamella studied

I 40

S,

v

X3

KDb6

(nm)

(nms -1)

a

(at.~ Zn)

(nm3s-I x 10-4)

160 280 310 320 165 200 210 140 175

20.1 14.0 14.0 23.4 43.3 43.3 43.3 16.6 27.0

6.3±0.3 3.7±0.6 3.9±0.8 2.4±0.7 3.7 ± 0.5 3.8±0.2 5.0±0.2 7.2±0.3 7.8±0.6

7.4±0.8 6.9±1.5 6.9±1.5 6.6±1.5 7.0 ± 1.2 7.8±0.6 7.6±0.7 6.2±1.0 6.9±1.4

8.07 29.6 34.5 99.8 31.8 43.3 38.2 4.54 10.6

Lamellae in the same colony are indicated by the brackets.

S O L O R Z A N O et al.: G R O W T H A N D D I S S O L U T I O N O F D I S C O N T I N U O U S P R E C I P I T A T I O N

1713

Table 2. Details of cells studied 478K aging temperature Lamella studied fi

5 6 I 10 11 12 13 14 17

S, (nm)

v (nms -I)

a

X3 (at.~ Zn)

KDb6 (nm3s Lx 10-4)

1065 710 410 430 580 500 245 200 480 510 510 400 520 280

18.6 24.0 24.0 29.3 t6.0 5.8 43.3 43.3 44.0 44.0 44.0 40.6 40.6 41.6

0.65 __+1.45 4.14 __0.3 1.3__+0.3 2.18 _+0.3 2.7 _+0.05 0.75 _+0.3 2.4_+0.11 2.3 _+0.2 4.0 _+0.3 3.8_+0.1 3.3_+0.1 7,6_+0.8 8,6 _+0.8 4.44 _+0.3

6.7 __+0.3 7.0_+0.6 8.2 _+0.6 7.5 __+0.3 8.1 _+0.5 7.6_+0.4 7.4 _+0.7 6.7 _+0.2 7.3__+0.5 7.3__+0.4 7.6_+1.1 7.2 _+ 1.7 6.3 __+0.8

29.2 31.0 24.8 19.9 19.3 10.6 7.5 25.3 29.8 34.2 8.5 12.7 7.35

were made, across the lamellae, approximately 50 nm behind the quenched front. The results are summarized in Tables 1 and 2. In addition, some typical analyses and their corresponding micrographs are shown in Figs 10-12. The analytic scans were then fitted to a solution of the (planar) moving boundary

X-ray microanalysis. The value obtained for 478K was 12 at.~o Zn, which is close to the coherent solubility limit proposed by Ardell et al. [10]. The solute profiles in the depleted alpha lamellae play an important part in the characterization of the reaction. In the present case, microanalytic scans

(a) 12.5

428 K S ='175 n m v = 2 7 n r n s -1

•

,

11.0

~

,

a =7.8

9.5

.-e 8.o

at

6.5 I 5.0

-

0

I

I

I

I

I

-60

-40

-2o

0

;to

Distance

I 40

I

I

60

80

(nm)

(b) Fig. ] ]. A]-22% Zn, 300 s at 428K. (a) CDF image using a product matrix reflection, emphasizing the change in image contrast between neighbouring c~' ]amel]ae. (b) Microana]ytica] scan across c~' ]amel|a of Fig. ll(a).

1714

SOLORZANO et al.: GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION (X3). In all cases, X0 was taken as the original matrix composition, 22 at.%Zn. The zinc-rich phase was etched out in most cases, allowing a closer microanalytical approach to the ct'/fl interface.

(c) Dissolution o f the discontinuous product

Fig. 12. A1-22%Zn, 300 s at 478K. Diverging fl plates, and the corresponding composition profile in ~'.

'~':7~

,(

,

t 300

The main purpose of this study was to see if the dissolution process was continuous or discontinuous. From examination of Figs 13-15, it is clear that the dissolution process involved volume diffusion, and that the grain boundaries remained essentially static during dissolution. A similar result was obtained for the A1-28% Zn alloy. The most striking features of the post-dissolution microstructures are walls of screw dislocations marking the planes of impingement of pairs of alpha/beta dissolution interfaces. Occasionally, edge dislocation walls were formed perpendicular to the original screw walls, Fig. 15.

nm I

Fig. 13. Bright field TEM image of the post-dissolution microstructure; A1-22%Zn, 300 s at 478K and 10 s at 595K.

diffusion equation due to Cahn [11], using a nonlinear regression which gave a best value of the coefficient a in the equation

Xo- X

_

Xo- X3

cosh (,/~ z/s)

Fig. 14. Interfacial dislocation arrangements, each formed where two ot'/fl dissolution interfaces have met.

cosh (x/a/2)

where a--

v S2 k Dh6

and X0 is the original alloy content, X3 is the composition of the ~' phase in contact with the fl lamellae, v is the growth velocity, S the interlamellar spacing, and Z is a coordinate perpendicular to the growth direction and to the plane of the ct'/fl interface, with origin at the centre-line of the ct' lamella. k is a distribution coefficient which relates equilibrium boundary compositions to those in the depleted lamellae, D h and 5 are the grain boundary diffusion coefficient and thickness respectively. This procedure permitted the evaluation of the product kDh6, and allowed the self-consistent extrapolation of the composition profile to the composition of the plane in contact with the fl lamellae,

Fig. 15. Bright field STEM micrograph of the dissolution structure; the recovery process has produced tilt boundaries normal to the twist boundaries of Fig. 14.

SOL6RZANO et al.: GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION 1715 DISCUSSION The initiation of the reaction, in all of the present observations can be related to the presence of grain boundary precipitates, which formed first, and which presumably act to impart to the boundary an initial displacement, which, in turn, permits the further development of the reaction. Thus, the initiation process, like that proposed by Tu and Turnbull [12], is of a type that requires a precursor precipitateboundary interaction. Unlike their result, in which the reduction of interfacial energy caused the displacement, our observations suggest that the interfacial energy is at first increased by the boundary displacement. The boundary precipitates in A1-Zn, rather surprisingly, appear to have no well-defined habit plane to cause a displacement on the Tu-Turnbull model. We therefore propose that the initial displacement is aided by chemically induced migration, a phenomenon which has been shown to occur in more dilute A1-Zn solutions [13]. In this model, the precipitates act as sinks to which solute diffuses, and the diffusion process, perhaps aided or biased by local capillary forces, causes boundary displacement. The displacement is against boundary curvature, and hence in a manner that increases the total interfacial energy early in the process. We appreciate that by invoking chemically induced migration, we are simply throwing the argument back one stage. The origin of the initial displacements in chemically induced migration is obscure, as is the exact formulation for the "chemical force" for steady boundary migration. Nevertheless, both the chemical force and chemically induced grain boundary migration are by now well established experimental phenomena. Before discussing the steady growth process, some comment should be made concerning the competing general decomposition of the matrix. The speed with which this reaction occurred made it impossible to avoid interference from this source. However, the consistent observation of constant interlamellar spacings, developed after relatively short reaction times, suggests strongly that the matrix reaction reached a quiescent state early in the process. The microanalyses of the coarsened matrix structure, while they probably represent a limit of spatial analytical resolution of the STEM, are in accord with this idea, and we will assume that the depleted matrix has come close to the coherent solubility limit. The morphology of the steady transformation front is of particular interest in the light of the proposal, made above, that chemically induced migration plays a role in initiation of the discontinuous reaction. We see from Figs 2, 6 and 7 that the grain boundary is essentially always shaped such that migration takes place against curvature. Again, this is a strong qualitative indication that boundary migration is chemically induced. A number of quantitative or semiquantitative inferences can be drawn from the microanalytical traces

across the a ' lamellae: it is in principle possible to calculate an interface shape, by balancing chemical and capillary forces at every point. We believe that a detailed calculation of this type is not yet justified, as both grain boundary energies and morphologies are subject to uncertainty. However, a semiquantitative comparison seems to be in order. We assume that the parent phase has reached coherent equilibrium. Based on the data of Hultgren et al. [5], this implies a loss, at 420K, of about 270 J/tool of the original 380J/tool available. Then the remaining chemical forces on the ~'/~ ° interface are small, varying from about 20 J/mol at the centre of the lamellae, to about 70 J/tool for ~' in contact with ft. The corresponding radii of curvature, given by 7Vm/AG, would then be 250 and 70 nm respectively, for an assumed interfacial free energy (y), of 0.5 Jim 2, and a molar volume Vm, of 10cm3/mol. These values are in reasonable accord with the more regular quenched growth fronts, e.g. Fig. 6, suggesting that an equilibrium of forces is indeed maintained at the front during growth. From a macroscopic point-of-view, the analytical traces demonstrate that a significant amount of supersaturation remains in the ~' phase, and permit its quantitative evaluation. In a separate contribution, [20], this observation is discussed in terms of the postulate that the system adopts a spacing and growth rate that maximizes the rate of dissipation of free energy [11]. The quantitative treatment of the microanalyses yields two further results of current interest: these are, the values of diffusion coefficients within the moving grain boundaries; and the compositions in the alpha phase at the ~/fl boundary, )(3 in Tables 1 and 2. Concerning the extracted diffusion data, we believe that it is justified to take 2(0 as 22~ Zn for purposes of evaluating diffusion coefficients, effectively ignoring the matrix precipitation reaction. The wavelengths of the short range components resulting from continuous precipitation are sufficiently small that they should not be expected to intrude on the longer range transport of Zn to the fl lamellae. The diffusion data, calculated using this assumption, are summarized in the tables. They indicate general agreement with the tracer coefficients for Zn reported by Hassner [16], with a tendency for our data to lie near the lower range of that rather scattered data. The reason for the scatter in both sets of measurements may lie in the structure-sensitivity of the grain boundary diffusion process itself. If diffusion data from different cells within the same well-formed colony are compared, the differences in S will be small, and the local value of v will be the same. Then the local kDh6 products take approximately the same value. Another somewhat different effect can be seen when local values of v and S differ significantly within a colony. Then the degree of solute partitioning, reflected in the solute profile,

1716 SOLORZANO et al.: GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION enters to moderate the variation of the vS2 product. However, the greatest Variation is found when values of kDh6 from different cell fronts within the same specimen are compared. The same kind of observation was made by Porter and Edington, in their study of the Mg-AI system [18]. There, the variation of diffusion data within a specimen was considerably greater than the error of a given measurement. The concentration of Zn in a at the triple junction, X3, is greater than the equilibrium composition Xe. Even allowing for the uncertainty in the measured values and in the equilibrium phase diagram, we find a significant difference between the two compositions. Porter and Edington, in contrast, reported that the composition of the growing lamellae in contact with the fl phase was equal to the equilibrium value. The discrepancy in the present case can be reduced, but not entirely eliminated, by assuming that the solubility is modified by a uniform hemicylindrical curvature at the tips of the beta precipitates. Subtracting the estimated capillarity contributions from the experimentally determined values of X 3 would yield plane interface equilibrium mole fractions at 428 and 478K of 0.045 and 0.066, to be compared with the published solubility values of 0.040 and 0.060. Considering all of the uncertainties inherent in such a comparison, we conclude that the capillarity term can bring the experimental values of X3 into reasonable accord with equilibrium data. This result is of particular interest, as it allows a distinction between the theoretical treatments of Cahn [11] and Hillert [14]. The observation that X 3 falls close to a capillarity modified equilibrium is the first of its kind, and serves to provide a measure of support for the detailed interface-modelling approach taken by Hillert. Concerning the dissolution microstructures, it is clear that dissolution is not discontinuous, as was thought by Pawlowski and Truszkowski [19]. However, they had no microstructural evidence to support their assumption. The surprising observation of residual dislocation nets is consistent with the evidence of the dark field micrographs of the discontinuous product. These show differences in orientation, often alternating from one a lamella to the next. (Fig. 11). Thus, the misorientation appears to be connected with the growth process, and to be preserved during dissolution. We are uncertain of the origin of this deformation process and at present can only speculate that it is derived from a torque exerted by the more coherent planar a-fl interfaces. The effect appear to be reproducible, and invites further study. CONCLUSIONS We have found evidence for chemically assisted initiation of the discontinuous product in A1-Zn alloys; from this beginning, the reaction gradually seeks a steady lamellar precipitation mode. Once the steady process is established, the product forms with a considerable degree of stored free energy in the ct

phase, as indexed by the composition profiles left behind the sweeping grain boundaries. The steady precipitation process occurs despite a competing fast general reaction, which is thought to reach relative quiescence when the depleted matrix nears the coherent solubility limit. Diffusion data extracted from the microanalytical profiles are consistent with, even a little lower than, tracer data for zinc in aluminum-zinc alloys. Extrapolation of the traces to obtain the composition of ct in contact with the beta lamellae indicates that this concentration has very probably reached a capillarity-modified equilibrium value, determined by the local curvature of the beta lamellae at the triple point. The dissolution process is controlled by volume, rather than grain boundary diffusion. Residual dislocation networks, formed where the two phase boundaries come together, suggest that a small but reproducible twist deformation is accomplished during growth and preserved during dissolution. The origin of this deformation is at present unclear. We conclude that the high-resolution dedicated STEM, in its several modes of operation, is capable of yielding information of value in increasing our understanding of solid-state transformations. In particular, the present study has exploited the microanalysis capability and has thereby allowed the determination of diffusion data for the moving grain boundary and measurement of the residual supersaturation within the depleted lamellae in a discontinuous transformation product. In addition, the microanalyses have yielded information on the theoretically significant quantity, ?(3. Acknowledgements--This research was supported by the

Natural Sciences and Engineering Research Council of Canada. I. G. Sol6rzano is grateful for support from the Canadian International Development Agency and STI (Brazil). REFERENCES

1. V. Melhortha and K. B. Rundman, Metall. Trans. 3A, 1551 (1972). 2. E. P. Butler, V. Ramaswamy and P. Swan, Acta metall. 21, 517 0973). 3. K. N. Melton and J. W. Edington, Acta metall. 22, 1457 (1974). 4. M. Vigayalakshmi, V. Seetharam and V. S. Raghunathan, Acta metall. 30, 1147 (1982). 5. R. Hultgren, R. L. Orr, Anderson and K. K. Kelly, Selected Thermodynamical Properties of Binary Alloys.

6. 7. 8. 9. 10.

Am. Soc. Metals, Metals Park, OH (1973). D. Chetham and F. R. Sale, Acta metall. 22, 333 (1974). G. Cliff and G. W. Lorimer, J. Microsc. 103, 203 (1975). A. Bouzaher and J. P. Simon, Scripta metall. 16, 687 (1982). S. F. Bauman, I. Mitchel and D. B. Williams, Acta metall. 29, 1343 (1981). A. J. Ardell, K. Nuttal and R. B. Nicholson, The

Mechanisms of Phase Transformations in Crystalline Solids, p. 22. The Institute of Metals, London (1969). 11. J. W. Cahn, Acta metall. 7, 8 (1959). 12. K. Tu and D. Tumbull, Aeta metall. 15, 369 (1969).

SOLORZANO et al.: GROWTH AND DISSOLUTION OF DISCONTINUOUS PRECIPITATION 13. K. Tashiro and G. R. Purdy, Scripta metall. 17, 455 (1983). 14. M. Hillert, Ref. [10], p. 231. 15. M. Hillert, Metall. Trans. 3A, 1729 (1972). 16. M. Hillert, Acta metall. 30, 1689 (1982). 17. A. Hassner, Kristall Tech. 9, 1371 (1974). 18. D. Porter and J. W. Edington, Proc. R. Soc. 385A, 335 (1977). 19. A. Pawlowski and W. Traszkowski, Acta metall. 30, 37 (1982). 20. I. G. Solbrzano and G. R. Purdy, Metall. Trans. To be published. APPENDIX

The experimental determination of the Cliff-Lorimer factor for a given system, in a given instrument, requires the

A,M. 32/10~J

1717

utilization of standards. The well-known tendency of A1-Zn alloys to decompose at room temperature, added to the occurrence of surface precipitation in sufficiently-exposed A1-22 at.~ Zn thin foils, demands that certain precautions be taken. The alloy studied here for discontinuous precipitation and other dilute alloys were solution treated at 370°C, quenched in water at 0°C, kept in liquid nitrogen, thinned at -30°C, and immediately taken to the STEM stage. In this manner the alloy solid solution was preserved. Care was taken in order to ensure the applicability of the thin film conditions. Whenever possible the local foil thickness was measured through the CBMD pattern obtained simultaneously with the intensity ratio of characteristic spectra. The kZn-Al factor obtained here has the value of 1.25 + 0.08. Errors are estimated for the present microanalytical results as ___10~o.