Electron microscopy of the grain structure of metal films and lines

North-HollandUltramicr°sc°py 51 (1993)328-338 ultramicroscopy

Electron microscopy of the grain structure of metal films and lines

David A. Smith 1, Mar t in Small and Carol Stanis

IBM Research Division, Thomas J. Watson Research Center, P.O. Box 218, Yorktown Heights, NY 10598, USA

Received 14 December 1992

The demands of ULSI in the sub-micron regime require the understanding of properties and control of microstructure in a situation where one or more of the characteristic dimensions of the sample is of the same order as the grain size of the material. Grain structure, as determined by deposition conditions, sample geometry and thermal treatment during processing, remains the first-order microstructural factor governing the properties. The grain structure evolves during both deposition and processing. A major practical and scientific issue is the elucidation of the conditions which will result in a stable grain structure. This happens when the forces acting on the grain boundaries are in equilibrium.

1. Introduction

The demands of ULSI in the sub-micron regime require the understanding of properties and control of microstructure in a situation where one or more of the characteristic dimensions of the sample is of the same order as the grain size of the material. The resistance of metal lines to failure by electromigration or stress-void formation depends on microstructure and alloy content [1,2]. In the past great emphasis has been given to alloying as a means to increase lifetime but without a concomitant effort to control or even char- acterize any aspect of the microstructure. How- ever, the continued drive to ULSI has resulted in the need for a deeper understanding of the microstructure of lines now that the grain size is of the same order as the line width. It has been reported that the standard deviation of the mean time to failure from electromigration diverges as the grain diameter becomes of the same order as the line width [3]. This paper reports the applica- tion of a variety of techniques of microscopy to the characterization of metal films and interconnects. The structure of gold films and aluminum-

1 Present address: Stevens Institute of Technology, Hoboken, NJ 07030, USA.

and copper-based films and lines has been inves- tigated by transmission electron microscopy, scanning ion microscopy and environmental scanning electron microscopy. Confinement by "passivating" layers may impede grain growth because of hydrostatic tensile stresses [4]; at the same time it may prevent grain growth stagnation because of the partial or total elimination of surface grooving [5]. Confinement can also result in the formation of voids by a creep process. These voids shrink or grow as the stress state is changed from tensile to compressive as a result of differential thermal expansion during the course of thermal cycling. The ultimate goal of this inten- sive effort in microstructural characterization is to permit an explicit s tructure-property relation- ship to be developed with the intention of pro- ducing interconnects with a higher performance and a greater reliability. Since the grain structure is the first-order determinant of behavior in poly- crystalline materials it is useful first to consider the factors which govern the grain structure of films and interconnects.

2. Microstructure of films

Study of the microstructure of films in the as-deposited condition reveals that all observed

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D.A. Smith et aL / Grain structure of metal films and lines 329

ZONE l i t ZONE T I l ZONE 11" I I ZONE TIT

f : t

0.1 0,2 0.3 0.4 0.5 0.6 0.7 Ts/Trn

Fig. 1. Schematic form of microstructures observed for metals as a function of the substrate temperature expressed as a

fraction of the melting point of the deposit (see text).

structures fall into a small number of classes [6-8]. Fig. 1 illustrates this point by showing in schematic form the microstructures observed for metals as a function of the substrate temperature expressed as a fraction of the melting point of the deposit, T~. The microstructure of an as-deposited film is a constrained stagnant structure for which the surface topography and grain structure have both developed to a stage limited by the deposition time and temperature. All the sketched microstructures are columnar; the dif- ferentiation of the various structures rests on factors such as grain aspect ratio, crystallinity, density and crystallite orientation distribution. These microstructural measures are a result (almost exclusively) of the interplay, through their simultaneous occurrence, of a variety of thermally activated transport processes, principally surface diffusion and grain boundary migration.

The substrate temperature is the dominant factor which determines the as-deposited grain structure of films. Characteristic microstructures are produced in particular temperature ranges; the original classification due to Movchan and Demchisin [6] recognized three regimes: zone I, T s < 0.3T m, equiaxed grains, the film thickness is greater than the grain diameter; zone I1, 0.3T m < T~ < 0.5Tm, columnar grains extending through the film thickness but the grain diameter < film thickness; zone III, 0.5T m < T~, as zone II but the grain diameter exceeds the film thickness. The grain size distribution is monomodal in all three zones. Electron microscopy revealed that at the transition between zones I and II the grain structure has a bimodal size distribution; the earlier investigations utilized X-ray diffraction and scan-

ning electron microscopy which lack the sensitiv- ity combined with spatial resolution necessary to detect this structure.

The transition from zone I structures through zone T to those of zone II occurs because different grain boundaries become mobile at different temperatures. In zone I almost all the boundaries are immobile, whilst in zone II all are mobile. In zone T an increasing proportion of the boundaries present become mobile. Consequently, the probability of any boundary moving across a grain and reacting to form another mobile boundary is increased. Since larger grains have more neighbors, continued growth of particular grains is more probable, once they have established a size advantage. This process is responsible for the bimodal grain size distribution observed in zone T. The onset of surface diffusion in this range of T s also contributes to a decrease in porosity. As the substrate temperature is raised, with a fixed deposition rate, more lateral grain growth occurs relative to the thickening by deposition. A uni- form columnar grain structure results, with the aspect ratio, grain diameter D divided by film thickness h, increasing and ultimately exceeding unity. This is the characteristic zone III structure. The aspect ratio of the grains depends on the ratio of the deposition rate and the rate of grain growth. In this sense zones II and III are not physically distinct. In zones II and III surface energy anisotropy evidently plays a role in the development of preferred orientations which are characterized by fiber textures in which closely packed atom planes are parallel to the film surface.

3. Mobility and forces on grain boundaries

In a metal, grain boundaries are never thermo- dynamically stable defects, i.e. there is no equilibrium concentration of grain boundaries, in contrast to the case of point defects. Grain growth of a material proceeds so as to decrease the interra- cial excess energy associated with the grain boundaries. There are other major influences on grain boundary movement; these are alloy addi- tions in the form of partitioned solute atom con-

330 D.A. Smith et al. / Grain structure of metal films and lines

centrations [9] or precipitates [10], surface grooves [5], grain-to-grain variations in surface energy [11], and constraints arising from passivation or the substrate. The last three factors are responsible for differences between the characteristics of grain growth in thin film and bulk materials. It is useful to make a simple quantitative assessment of the orders of magnitude of the predominant driving forces which may act on a grain boundary in a film as follows [12]. The change in free energy which accompanies grain boundary migration may be thought of as the product of a force acting on the grain boundary and the distance moved by the boundary. Table 1 gives expressions for the major forces acting on a grain boundary in a foil comprised of columnar grains such as are characteristic of the zone II and III microstructures. In table 1, "Yb, Ysv and Ysu are the surface energies of grain boundaries, the so l id /vapor interface and the sol id/substrate interface, re- spectively; h is the film thickness; A is the change in the quantity which follows; D is the grain diameter; d is the precipitate diameter and f the volume fraction of precipitate; C b is the concentration of solute atoms in the grain boundary; AV is-tl~e atomic excess volume at a grain boundary, N a is the atomic density in the grain boundary, tr is the hydrostatic stress and a the atom diameter.

In a line with a bamboo structure some of the forces change to the values indicated in table 2; w is the width of the line and it is assumed that the surface energy for any particular grain is isotropic.

Table 1 Forces on a columnar grain in a film

Origin of force Magnitude of force (per unit area)

Grain boundary energy 2 y b / D a) Surface energy variation Aysv / h + A y s u / h b) Precipitates 4 y b f / d Segregation 4ybC b / a Hydrostatic stress 2o-N.AV/D Grooving y 2 / h Ysv

a) In the three-dimensional situation the grain boundary energy leads to a force 3 y b / D per unit area.

b) For a film not on a substrate the surface energy variation force on a columnar grain is 2Aysv/h; for a grain with only one surface, the force is A y s v / h or Ay~u/h .

Table 2 Forces on a grain boundary in a bamboo structure

Grain boundary energy Y b / D Surface energy variation Aysv / h + 2 Aysv / w + Aysu / h Grooves Y~/Yw + ( 1 / h + 2 / w )

Here it is necessary to include the efforts of grooving and surface energy variation on the sides of the line. The velocity, V, of a boundary is given by the product of the grain boundary mobility and the net force acting on the boundary,

V=MP, (1)

where M is the mobility and P the driving force, which is sometimes written as A/z, the chemical potential change driving the process. On a one- dimensional argument the velocity of the interface may be described more explicitly by the expression [13]:

where r is the distance moved per activated even, v is the Debye frequency, A is a geometric factor, Ag the free energy of activation and kT has its usual meaning; r, A, u and Ag are all, in principle, structure-dependent.

Mullins [5] first showed that in a one-component thin film with columnar grains stagnation of grain growth occurs when the curvature driving force balances the retarding force arising from grooving. This analysis leads to a stagnation condition which can be obtained using the expressions in table 1.

= Y b / h Y s . (3) 2yb/D 2

The limiting grain size D~i m is 6h if 3y b is approximated to Ys; Dlim = 3h if grooving is possible on both film surfaces. However, if there is an additional driving force from surface energy anisotropy, the stagnation condition (eq. (3)) is modified by the addition of a term (Aysv+ Aysu)/h. In a one-component thin film with columnar grains, stagnation of grain growth now occurs when the sum of the curvature and surface energy driving forces balances the retarding force

D.A. Smith et al. / Grain structure of metal films and lines 331

Fig. 2. Grain structure of a gold film annealed to stagnation.

arising from grooving. The condition may be written as follows:

2Yb AYsv 7b 2 - - + - - = . ( 4 )

D h hysv

With the order of magnitude approximations that 37b = 7sv and Ays v = 1 ~Tsv there is no longer a limiting grain size; since the surface energy term is equal in magnitude to the groove pinning term, grain growth can continue without limit, apart from that imposed by the sample size, as long as a distribution of surface orientations is maintained. In this example both grooving and surface energy anisotropy at the substrate/ deposit interface have been neglected, but the conclusion is unchanged if these factors are in- cluded. Further discussion of stagnation conditions is deferred until after the presentation of some observations.

4 . O b s e r v a t i o n s o f g r a i n s t r u c t u r e

A gold film 50 nm in thickness was e-beam evaporated onto a silicon nitride window sub-

Fig. 3. Electron micrograph of a 50 nm {100} thick single-crystal gold film which is recrystallizing from {100} to {111} orientation; the final grain diameters are of the order of tens of microns.


strate at room temperature in a diffusion-pumped vacuum system with a base pressure of 5 × 10 - s Torr. The film was annealed in situ in a Philips EM430 transmission electron microscope at 700°C for 2 h after which the grain structure appeared stable; no neighbor switching was observed in the second hour of isothermal annealing. The as-deposited grain structure was bimodal but the annealed grain structures shown in fig. 2 consist entirely of secondary grains. The secondary grains have a strong {111} preferred orientation but the fine grains in the as-deposited microstructure are oriented randomly. The stagnant structure is strikingly different from a net of regular hexagons which is a stable structure in a two-dimensional system with isotropic grain boundary energy. The numbers inside the grains designate the number of grain neighbors for the labelled grain. When the intercept method is used the grains in the field of view range in diameter from about 100 to 740 nm (2h to 14.8h) where h is the film thickness. Many grain boundaries are curved. The grain-to-grain surface energy anisotropy is very small since the structure is comprised of {111} grains which have grown to impingement.

The effect of a sustained difference in surface energy between the growing and shrinking grains, preventing stagnation, is shown in fig. 3 which is an electron micrograph of a 50 nm thick {100} single-crystal gold film which is recrystallizing from {100} to {111} orientation; the final grain diameters are of the order tens of micrometers ( > 100h).

An aluminum film, 90 nm in thickness, was e-beam evaporated onto a silicon nitride window substrate at room temperature in a diffusion- pumped vacuum system with a base pressure of 5 x 10 - s Torr. The film was annealed in situ at 400°C for 1 h after which the grain structure appeared stable; this mierostructure is shown in fig. 4. Neither the as-deposited nor annealed grain structures revealed any evidence for abnor- mal grain growth nor was any preferred orientation detected by selected-area diffraction. A sur- vey of individual grain orientations was quite consistent with the selected-area diffraction data and offered no evidence for a preference for any particular surface normal. Again the stagnant

Fig. 4. Grain structure of an aluminum film which was an-

nealed in situ at 400°C for 1 h.

structure is not a net of regular hexagons and the numbers indicate the number of neighboring grains. The lines make clearer the location of two somewhat indistinct nodes. With the intercept method the grains in the field of view range in diameter from about 73 to 290 nm (0.8h to 3.2h). Many grain boundaries are curved. Mullins [5] pointed out that a doubly curved surface such as a catenoid has zero overall curvature. Such a surface viewed in projection would appear as curve with a finite width and extended on the convex side. The arrowed boundaries seem to have this characteristic. Addition of 2% copper results in a reduction in average grain size and secondary recrystallization as shown in fig. 5. The film was prepared by sputtering onto a silicon nitride window and is 50 nm thick. Again there are the characteristic doubly curved stagnant structures visible after annealing for 100 h at 200°C (473 K). The diameter of the secondary grain is ~ 3000 nm (6h) and the smaller grains are typically 30 nm (0.6h) in diameter.

Although transmission electron microscopy is a powerful technique for elucidating the grain structure and orientation of films, it suffers from


Fig. 5. Grain structure of annealed AI-2wt% Cu.

the disadvantages that both film and substrate (if present) must be electron transparent. This means that examination of metallization on integrated

circuits usually requires local thinning and that the area available for characterization is limited. Focused ion beam (FIB) imaging in a scanning ion microscope avoids these limitations.

In FIB images, the orientation of individual crystallites relative to the incident beam direction causes large variations in ion-induced secondary electron emission, enabling pronounced crystallographic contrast to be observed [14]. For crystalline materials, both nuclear and electronic stopping depend strongly on the lattice orientation of the crystal relative to the incident beam direction, due to ion-channelling effects [15]. As a result, the secondary electron yield from grains oriented with their low-index crystallographic planes parallel to the ion beam is small and they appear dark in a FIB image.

Ion microscopy was performed on a JEOL JIBL 106D focused ion beam system using a 20 pA, 100 keV Ga ÷ beam with an estimated beam diameter of 70 nm. The pressure in the sample chamber was approximately 5 x 10 -6 Torr. Sam- ples were oriented normal to the ion beam. The semi-angle of convergence of the focused ion beam used in this study is 0.03 ° , much less than the acceptance angle for channeling, typically a

Fig. 6. FIB images of copper deposited by CVD onto a polyimide-coated silicon substrate.


Fig. 7. Some examples of stress voids, which appear dark, in passivated AI -2wt%Cu lines. The images were recorded at elevated temperature using a heating stage in an environmental scanning electron microscope. The light particles are 0-phase.

D.4. Smith et al. / Grain structure of metal films and lines 335

few degrees, and thus the contrast is not limited by the angular spread of the beam. Images were formed from detection of secondary electrons using an annular multichannel plate detector lo- cated directly above the samples. Fig. 6 shows FIB images of copper deposited by CVD onto a polyimide-coated silicon substrate. The film thickness is 500 nm and the copper was deposited after the polyimide was patterned. The three uppermost lines in each panel of the figure reveal the effects of geometry and annealing on microstructure; the lower two "lines" are sections cut from large pads and indicate some differences between the behavior of lines and films. The proportion of bamboo grains increases after annealing for both the 500 nm and 1.5 ~zm lines shown. There is a greater proportion of bamboo grains in the narrower lines. The "lines" cut from pads seem to show a greater propensity for twin- ning and somewhat fewer "bamboo" grains; these grains also appear more equiaxial. In the as-deposited condition the grain diameters are larger in the lines cut from pads than in the deposited lines. The annealing treatment for 2 h at 400°C substantially eliminates this difference. The de- composition of the polyimide precludes heat treatment at a higher temperature, which would be expected to drive the grain growth process nearer to stagnation. On the basis of modelling [16] the stagnant structure in a line with w / h --- 1

is bamboo. Interconnects see service under "passivation"

which is an encapsulating layer, usually silicon dioxide or silicon nitride. The consequence is that post-deposition grain growth process takes place under conditions of stress. The stress state ap- proaches hydrostatic conditions. Coupling to the volume contraction which accompanies grain growth can thus affect both the approach to stagnation and the actual stagnant microstructure [4,17]. An effect of biaxial tension has been found; the stagnant grain size observed after annealing under biaxial tension is less than that for uncon- strained growth [4]. A triaxial compressive stress of the order 1% can, in principle, be relaxed by grain growth; however, a corresponding tensile stress may be relaxed by the formation of voids or the punching of prismatic dislocation loops. Voids

can be produced in aluminum alloys during holding at room temperature [2]. Oxide passivating films are typically deposited at 400°C (673 K); treatment at this temperature permits the mechanical stresses resulting from earlier processing steps to relax by thermally activated processes. Cooling to room temperature results in the gen- eration of hydrostatic tension in the metallization because of differential thermal expansion. The coefficients of linear expansion are 4 × 10 -6 K-1 for the silicon substrate and the passivation, and 27 × 10 -6 K -1 for aluminum-based metallization. Cooling to room temperature results in vol- umetric strains of about 1% which are relaxed by void formation; conversely, heating above the temperature of passivation deposition puts the metallization into compression which can be relaxed by shrinkage of any voids produced during holding at low temperature. Some examples of such voids are shown in fig. 7. These defects were produced without any treatment other than cooling to ambient temperature. The lines in fig. 7 are sputtered AI-2wt%Cu which is deposited onto a silicon wafer coated with 500 nm of silicon oxide, patterned, heated to 400°C (673 K) for 30 min and then coated with 1/zm of silicon dioxide. The images were recorded at elevated temperatures using a heating stage in an environmental scanning electron microscope (ESEM) [18]. The secondary electrons which contribute to the image are produced by back-scattered electrons which penetrate the passivation; these are sometimes called SE2 electrons. The light particles are 0-phase. There are two layers of metallization on this chip and the accelerating voltage needs to be adjusted to optimize the image according to the depth of the feature. Charging of the sample is avoided by the presence of 2.5 Torr of water vapor in the differentially pumped specimen chamber. Heating to a temperature in excess of 400°C (673 K) results in some shrinkage of the voids (and dissolution of the 0-phase precipitates). Stress reversal results in void shrinkage. These processes are illustrated in the environmental SEM images in fig. 7. Separate observations after removal of the passivation show that the majority of the voids and the coarse 0-phase particles are at grain boundaries.

336 D.A. Smith et aL / Grain structure of metal films and lines

5. Discussion

An order of magnitude estimate of the final grain size in a material can be made by writing an expression for the static equilibrium of the forces on a grain based on table 1, and substituting appropriate values for the characteristics of the material. Thus for the case of a columnar grain structure in an unsupported film containing precipitates and solute segregated to the grain boundaries, the equilibrium condition can be written:

2y b 2A7~ 2y~ 4Ybf 4"ybC b + - - - + - - + - - ( 5 )

D h hy~ d a

The factor 2 in the surface energy and groove pinning terms takes account of the fact that the film has two free surfaces. An immediate diffi- culty is that in general C b is not known. Recent measurements using a VG STEM [19] suggest that in the AI-2wt%Cu film shown in fig. 5, C b is of the order 1 at%. Examination of the figure shows that there is a remarkable absence of 0- phase which in thicker films precipitates as coarse particles ~ 100 nm in diameter at grain triple junctions. The discrepancy is a consequence of surface segregation. With the approximation that the grooving and surface energy anisotropy terms balance, the limiting grain size in the example is given by:

D = a / 2 C ~ , (6)

i.e. the limiting grain size is expected to be of the order of 50 atom diameters, i.e. 15 nm. For a thicker sample, where surface segregation does not deplete the grains of copper, d the precipitate diameter is typically 50 nm and the volume fraction of precipitate is about 0.01; f / d is then about two orders of magnitude less than C b / a SO

that Zener pinning [10] is negligible relative to solute pinning. Both expressions are actually overestimates since each is based on the pre- sumption that the grain boundary escapes from all the pinning sites simultaneously. The agree- ment, within a factor 2, between this order-of- magnitude estimate and the observation (for the pinned grains) in fig. 5 is thus very reasonable.

The justification for the neglect of the surface grooving and surface energy anisotropy forces for aluminum needs further elaboration. The experi- ments were done under circumstance in which an oxide film would certainly form. Consequently, the magnitude of Ay, would be substantially re- duced since the sol id /vapor interface has been replaced by a solid-state surface. This considera- tion applies to both surfaces since the specimen was on a silicon nitride substrate. The energetics and kinetics of grooving would both be changed relative to a foil surface, presumably in the direction of decreasing the effect of surface grooving.

The unalloyed aluminum film annealed to stagnation at 400°C (673 K) did not undergo any secondary grain growth, and the range of grain diameters is less than for A1-2wt%Cu but the mean grain size is larger. However, the largest grain in AI-2wt%Cu is larger than that in the pure aluminum.

The influence of surface energy anisotropy is clearest for the case of gold which has no surface oxide. The difference in the behavior of the poly- crystalline film, fig. 2, and the initially single- crystal film, fig. 3, i.e. no stagnation in the latter case, results from the fact that in the polycrystal many favorably oriented grains grow simultaneously while in the single crystal only a few {111} grains nucleate. Once the large surface energy anisotropy is eliminated Mullins' [5] mechanism becomes effective and grain growth stagnates. Consequently, the upper limit on grain size in this case is governed by the location and density of the nuclei for {ll l}-oriented grains.

The largest grains in the gold film heated to stagnation exceed the diameter predicted by the widely used Mullins criterion. It was shown that an anisotropy of surface energy such that Ay~ > y~ is sufficient to prevent stagnation of grain growth. On a broken bond model this corresponds to the difference in surface energy of a {111} and a {100} surface in a face-centered cubic metal. Lesser differences in Ys would lead to stagnation but at grain diameters >> than the film thickness. There is a difference in the size of the largest grains observed in gold and aluminum annealed to stagnation. In gold the largest grains are typically 10h in diameter; in the aluminum described here the


largest grains have a diameter of about 3h. In principle, passivation of a deposit on a substrate may decrease or eliminate grooving and thus eliminate one reason for stagnation. However, in gold the largest grains frequently but not exclusively are formed when an oriented grain con- sumes a random fine-grained structure so that surface energy anisotropy can overcome pinning by grooving. In view of the reactivity of aluminum it is possible that there is some pinning by impuri- ties which would serve to mask any decrease in pinning by grooves.

Additional pinning by surface grooving can occur for lines since some grain boundaries inter- sect the edge of a line. This process has been modelled systematically by Walton et al. [16]; for the present purposes their key result is that the approach to stagnation is exponential and the limiting grain structure of an interconnect, with width w and height h, is bamboo providing w/h < 2.7. The time to achieve stagnation is both material- and geometry-dependent. As a repre- sentative example, an aluminum line with an aspect ratio of 1 and a width of 1 # m is predicted to reach stagnation and to achieve a bamboo structure in 10 4 S at 400°C. This time scale gives a basis with which the data shown in fig. 6 can be interpreted; copper has a higher activation energy for grain growth than aluminum and should evolve to a bamboo structure when w/h < 2.7 but not in a 2 h anneal at 400°C (673 K).

Chaudhari [17] derived a stagnation condition based on the production of a back stress arising from the expected contraction accompanying grain coarsening. Stagnation in an initially stress- free film is predicted to occur only when the starting grain size is less than 1.3 nm. For larger initial grain sizes the stress effect is to retard but not arrest grain growth. The effect on grain growth of the stress arising from differential thermal expansion is minimal in the data presented here for aluminum-based metallization since the passivated samples were relaxed at 400°C (673 K), the annealing temperature. However, for samples for which mechanical constraint is established at ambient temperatures the effect may be significant in perturbing an otherwise equilibrium situation. Using the expression from table 1 with ~r =

109 Pa, AV= 1.6 × 10 -30 m 3, N a = 10 t9 m -2, the fo rce /un i t area on a boundary is about 3 × lO-2/D. This is of order 10% of the strength of the capillary force. However, the more significant microstructural effect of stress is the formation of voids.

6. Conclusion

The final grain structure of films and lines depends on both kinetic and thermodynamic factors. The bamboo structure which is desirable for electromigration resistance is not produced by the heat treatments ordinarily used during the processing of metallization. The final grain size in films and lines is strongly dependent on the alloy content and distribution. Under conditions of sustained anisotropy of surface energy there may be no upper limit on grain size.

Acknowledgement

We are grateful to P. Blauner for assistance with FIB and the staff of Electroscan Corpora- tion for access to and assistance with the ESEM. Fig. 1 was first published in ref. [4] and is repro- duced with permission of Pergamon Press.

References

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[2] See, for example, T. Kwok, K.K. Chan, H. Chan and J. Simko, J. Vac. Sci. Technol. A 9 (199l) 2523.

[3] E. Kinsbron, Appl. Phys. Lett. 36 (1980) 968. [4] D.A. Smith and C.R.M. Grovenor, Trans. Jpn. Inst. Met.,

Suppl. 27 (1986) 969. [5] W.W. Mullins, Acta Met. 6 (1958) 414. [6] B.A. Movchan and A.V. Demchisin, Fiz. Met. Met-

alloyed. 28 (1969) 83. [7] J.A. Thornton, Annual Rev. Mater. Sci. 7 (1977) 239. [8] C.R.M. Grovenor, H.T.G. Hentzell and D.A. Smith, Acta

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[10] J.E. May and D. Turnbull, Trans. AIME 212 (1958) 769. [ l l l C.V. Thompson and H.I. Smith, Appl. Phys. Lett. 44

(1984) 603.

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[12] H.P. Stuwe, in: Recrystallisation of Metallic Materials, Ed. F. Haessner (Reiderer, Stuttgart, 1971).

[13] D. Turnbull, Trans. Met. Soc. AIME 191 (1951) 661. [14] P.H. Larnarche, R. Levi-Setti and K. Lain, IEEE Trans.

Nucl. Sci. NS-30 (1983) 1240. [15] K. Nikawa, K. Nasu, M. Murase, T. Kiato, T. Adachi and

S. Inoue, in: Proc. IRPS'89 (1989) p. 43.

[16] D.T. Walton, H.J. Frost and C.V. Thompson, Appl. Phys. Lett., in press.

[17] P. Chaudhari, J. Vac. Sci. Technol. 9 (1990) 520. [18] G.D. Danilatos, J. Microscopy 160 Pt.1 (1990) 9. [19] M.B. Small, D.A. Smith and A.J. Garratt-Reed, Scr. Met.

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Electron microscopy of the grain structure of metal films and lines