Microstructural evolution during film growthI. Petrova)

Frederick Seitz Materials Research Laboratory and Department of Materials Science, University of Illinois,Urbana, Illinois 61801

P. B. BarnaResearch Institute for Technical Physics and Materials Science, Hungarian Academy of Sciences,Budapest H-1325, Hungary

L. HultmanThin Film Division, Physics Department, Linko¨ping University, S-581 83 Linko¨ping, Sweden

J. E. GreeneFrederick Seitz Materials Research Laboratory and Department of Materials Science, University of Illinois,Urbana, Illinois 61801

~Received 9 June 2003; accepted 17 June 2003; published 2 September 2003!

Atomic-scale control and manipulation of the microstructure of polycrystalline thin films duringkinetically limited low-temperature deposition, crucial for a broad range of industrial applications,has been a leading goal of materials science during the past decades. Here, we review the presentunderstanding of film growth processes—nucleation, coalescence, competitive grain growth, andrecrystallization—and their role in microstructural evolution as a function of deposition variablesincluding temperature, the presence of reactive species, and the use of low-energy ion irradiationduring growth. © 2003 American Vacuum Society.@DOI: 10.1116/1.1601610#

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I. INTRODUCTION

Polycrystalline thin films have found diverse applicatioranging from metallization and dielectric layers to opticmagnetic, and tribological coatings, to diffusion and thermbarriers. Progress in thin film science and technology acerated following the maturation of vacuum technology whgave birth to the American Vacuum Society~AVS! in 1953.The first divisions within the AVS, the Vacuum MetallurgDivision ~today the Advanced Surface Engineering Divisio!and the Thin Film Division both focused on polycrystallincoatings and thin films. In subsequent years, understanof thin film synthesis has benefited tremendously fromvances in the areas of surface science and analysis.

Thin films exhibit a wide variety of microstructures chaacterized in terms of grain size and crystallographic oriention, lattice defects, phase composition, and surface mphology. Industrial demand for ever lower processitemperatures in device and product manufacturing methat films are often deposited at temperaturesTs which areless that 0.2–0.3 of the melting pointTm ~in degrees K!.Thus, film synthesis generally takes place far from thermdynamic equilibrium. As a consequence, microstructure ding deposition typically evolves in a competitive fashion athe kinetic limitations induced by low-temperature growallow for the controlled synthesis of metastable phasesartificial structures such as multilayers and nanophase mrials.

Among the determinant atomic processes controllstructure evolution during film deposition are surface abulk diffusion. These processes are affected by, in additio

a!Author to whom correspondence should be addressed; electronicpetrov@uiuc.edu

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Ts , energetic particle bombardment which can be usedmanipulate adatom mobilities and nucleation rates. The pence of alloying or impurity elements and their segregatto surfaces and grain boundaries can also strongly influethe final result.

Extensive studies of the correlation between film structand deposition parameters have been carried out overpast five decades. From an understanding of film formatfollows the possibility for microstructural and nanostructuengineering in order to design a material for specific techlogical applications. This has led to the development arefinement of structure zone models~SZMs! which system-atically categorize self-organized structural evolution durphysical vapor deposition~PVD! as a function of film growthparameters.1–7 The history of SZMs has been reviewed bThornton8 and Barna and Adamik.9 In 1969, Movchan andDemchishin1 observed that the microstructural evolutionevaporated Ti, Ni, W, ZrO2 , and Al2O3 coatings can be systematically represented by a single SZM diagram plotted vsus film thickness and the homologous temperatureTs /Tm .Significant contributions to the understanding of microstrutural evolution in PVD thin films are documented in Ref10–37. Electro-deposition and chemical vapor depositirepresenting extreme cases of low- and high-temperagrowth regimes, exhibit many features of microstructuevolution which are analogous to PVD due to similaritiesthe atomistics of the growth processes.38

The first SZMs were derived from relatively lowresolution optical and scanning electron microscopy obsetions. Later, cross-sectional transmission electron micrcopy ~XTEM! and scanning probe microscopy~SPM!analyses were employed to provide more detailed structcharacterization.In situ electron microscopy has revealed thil:

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dynamics of film growth. This, together with results frominsitu SPM studies and computational materials science,provided detailed atomistic insights into microstructural evlution during polycrystalline film growth.

This article is a short review of recent progress in undstanding microstructural evolution in thin films synthesizby low-temperature PVD on amorphous and polycrystallsubstrates. The use of amorphous substrates allows theperimentalist to isolate the effects of individual depositivariables on texture development. Polycrystalline substrbias texture through local pseudomorphic epitaxy, howethe overall microstructure will still evolve toward a finastate driven by the extant deposition conditions. On btypes of substrates, film growth proceeds via a thrdimensional~3D! or Volmer–Weber mode. Related areasthin film science and technology not treated here are moing of nucleation kinetics~see companion article by Ratscand Venables39 in this volume!, epitaxial layer formation,and surface morphological evolution during epitaxy~articleby Cahill40 in this volume!.

We begin in Sec. II by introducing the fundamental filgrowth processes, and resulting SZMs, for PVD synthesipure elemental films. Then we describe SZMs for reactdeposition of multicomponent systems and discuss thefects of dopant and alloying elements~whether intentional ornot!, as well as phase formation, on the resultant nano-microstructures~Sec. III!. In Sec. IV, we describe the effectof ion irradiation in modifying film structure during depostion.

II. GROWTH OF PURE ELEMENTALPOLYCRYSTALLINE FILMS

The growth processes controlling microstructural evotion, presented schematically in Fig. 1, include nucleatiisland growth, impingement and coalescence of islands,mation of polycrystalline islands and channels, developmof a continuous structure, and film growth. When surfadiffusion rates are significant, film thickening proceethrough local epitaxy on individual grains. Grain coarsenii.e., recrystallization through grain boundary~GB! migration,can occur both during and after island coalescence.

The nucleation barrier is generally expected to be smleading to randomly oriented islands, for low-temperatudeposition on amorphous substrates.31,36 In situ TEM inves-tigations confirm this for studies of Au/SiO2

41 and In/C.42 Inthe latter case, nucleation of In deposited on amorphousubstrates atTs /Tm,0.4– 0.5 results in solid-phase islanwith random orientation, while at higher temperature theclei are liquid due to the suppressed melting point of nancale particles.43 Nucleation kinetics is affected by the adatobinding energy, crystal structure of the substrate matelattice defects, surface steps, and contamination.

During island coalescence, there is a strong driving fofor coarsening through surface atom diffusion and GB mtion. The island with lower energy per atom consumesother~s!, resulting in a new single-crystal island as the stem attempts to minimize the overall surface and interf

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energy. Thus, coarsening during coalescence is the firstmost active phenomenon leading to selection of preferorientation.44 Islands with the densest planes are typicaselected; that is,~111! for fcc, ~0002! for hcp, and~110! forbcc. Depending on temperature and island size, coarsecan be very fast, often termed liquid-like coalescence, occring by either rapid surface diffusion or by melting~for lowmelting-point materials! upon contact followed by crystalli-zation. The driving force is the release of edge and surfenergy associated with island coalescence. Rapid coacence also results in new open substrate area for seconnucleation. At lower temperatures or larger island sizcoarsening is slower and proceeds through GB migratGrain coarsening during coalescence of the contacting ctals is repeated until the local grain size becomes sufficielarge that grain boundaries are immobile. Figure 2 illustraseveral coarsening events during coalescence of In islandamorphous C.45

The SZM in Fig. 3, characterizing microstructure evoltion in pure elemental films, consists of three regions:9 ZoneI corresponds to very low deposition temperatures at whadatom diffusion is negligible; surface diffusion becomsignificant in the transition Zone T; and Zone II represefilm growth at deposition temperatures for which both sface and bulk diffusion are operative. The boundariestween the zones are diffuse and ‘‘transitions’’ occur graduaover wide ranges inTs /Tm .

During film growth in the low-Ts Zone I regime~Fig. 3!,an underdense structure with a fine fiber texture develoInitial in-plane grain sizes are set by the saturation nucleadensity. Adatom mobilities are low and columns preserverandom orientation of the nuclei as predicted by ballis

FIG. 1. Schematic diagram illustrating fundamental growth processestrolling microstructural evolution: nucleation, island growth, impingemeand coalescence of islands, grain coarsening, formation of polycrystaislands and channels, development of a continuous structure, andgrowth ~see Ref. 9!.

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models.15,46The columns are generally not single grains, bare composed of smaller more equiaxed grains, or cancompletely amorphous. Surface roughness developsfractal geometry47 which, due to wide angular distribution othe deposition flux, atomic shadowing, and limited surfadiffusion, leads to extensive porosity.

At higher film growth temperatures~Zone T!, grain coars-ening occurs during coalescence of small islands with lasurface to volume ratios, while grain boundaries becomemobile in continuous films. Orientation selection durincoarsening is incomplete, thus crystallites are nearly randor only weakly textured and there is a wide distributiongrain sizes. The orientation and size of individual crystalliwill determine their behavior during subsequent growth pcesses characterized by the competition among neighbo

FIG. 2. In situ plan-view TEM micrographs obtained during the growth acoalescence of In islands deposited at 0.5 nm s21 on amorphous C substrateat Ts540 °C. The time lapse between obtaining the left and right image0.3 s. Note the denuded zones surrounding the coalesced islands~see Ref.45!.

FIG. 3. SZM schematically representing microstructural evolution of pelemental films as a function of the reduced temperatureTs /Tm , whereTs

is the deposition temperature andTm is the melting point of the materialboth expressed in degrees K~see Ref. 9!.

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grains. In thisTs /Tm range, adatom surface diffusion is significant resulting in local epitaxial growth taking place oindividual grains. Pronounced columnar structure develin which the columns are actually elongated grains.

The primary features of Zone T competitive grain groware illustrated by the kinetic Monte Carlo~kMC! simulationof Gilmer et al.48 ~Fig. 4! for Al growth. While there areinitially equal distributions of 111 and 001 islands, the latorientation eventually dominates due to anisotropies in sface diffusivities and adatom potential energies. That is,average adatom residence time is significantly higher attice sites on low diffusivity~low potential energy! 001 sur-faces versus high diffusivity~high potential energy! 111 sur-faces. Thus, adatoms which are stochastically depositedgrain boundaries and, through surface diffusion, sample son both sides of the boundary have a higher probabilitybecoming incorporated at the low-diffusivity surface whiprovides the more stable, lower potential energy sites. Cversely, adatoms on high diffusivity planes have longer mefree paths with correspondingly higher probabilities to mooff the plane and become trapped on adjacent grains. T001 grains with low surface diffusivities grow faster. Atomshadowing exacerbates the difference, as protruding surfcapture more off-normal flux. Thus, low-diffusivity 00grains slowly expand, overgrow the high-diffusivity grainand become bounded by 111 facets. This leads to consable surface roughness which scales with the averageplane grain size.

The consequence of competitive growth is a continuochange in morphology, texture, and surface topography~and,hence, film properties!! as a function of film thickness. Neathe substrate, the microstructure consists of randomly

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FIG. 4. Kinetic Monte Carlo simulation of competitive texture evolutioduring low-temperature sputter deposition of an Al film. Islands~and latercolumns! with lighter contrast are 001 oriented, while darker islandcolumns are 111 oriented~see Ref. 48!.

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ented small grains out of which V-shaped columns withfavored orientations slowly emerge and overgrow kineticadisadvantaged columns. This gives rise to increasedferred orientation. The faceted column tops result, as noabove, in surface roughness which increases with thickngiving rise to open column boundaries due to atomic showing.

At still higher Ts /Tm ~Zone II!, bulk diffusion becomessignificant. GB migration takes place not only during coalcence, but throughout the film thickening process. Oriention selection during the coalescence stage is morenounced and is driven by a decrease in the total GB arewell as minimization of interface and surface energy.35 Largegrains with low surface and interface energy grow atexpense of smaller or unfavorably oriented grains. Normgrain growth is impeded if the grains have a strong textui.e., if the orientation selection was completed during coacence, or the mean in-plane grain diameter reaches twthree times the film thickness.35 Secondary recrystallizationalso called abnormal grain growth, may follow in which thgrain size distribution is transformed from monomodthrough bimodal, to a new monomodal distribution wimuch larger in-plane grain size. During secondary recryslization, the degree of texture is further enhanced. The fistructure is homogenous in the growth direction and coposed of columnar crystals with flat surfaces decoratedGB grooves.

III. REACTIVE DEPOSITION PROCESSES:GROWTH OF MULTICOMPONENTANDÕOR MULTIPHASE FILMS

Possibilities for controlling structural evolution duringrowth of pure elemental films are limited and the structuobtained are unstable against temperature increases. Tduring use the layers tend to restructure toward thermonamic equilibrium. The effects of additives~contaminants,dopants, or alloying elements! in controlling the grain sizeare well known in bulk materials. Similarly, polycrystallinthin films synthesized by reactive deposition provide adtional pathways for microstructure control while yielding ehanced thermal and process stability. It is important to nthat even very low concentrations~sometimes below the detection limits of modern analytical techniques! of uninten-tionally incorporated atmospheric contaminants such aster vapor,49 oxygen, and hydrocarbons, may be active‘‘grain refiners.’’50

Consider the case of O-containing Al films depositedroom temperature, aTs value which corresponds to Zonein the pure-Al SZM.51 Observed changes in film structuand orientation as a function of increasing oxygen conctration are summarized in Fig. 5.9 Oxygen has low solubilityin Al and segregates to surfaces and grain boundaries wit forms two-dimensional~2D! oxide layers~oxide tissuephases! which greatly reduce Al surface and GB mobilitieThese layers modify all thin film formation processes, liming grain coarsening during coalescence and film growThey also periodically interrupt the epitaxial growth of ind

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vidual crystallites and cause renucleation.42 By exploitingthese phenomena, new micro- and nanostructures cacontrollably formed.9

At low O/Al arrival rate ratios,JO/JAl;1023, oxygen isincorporated into the grain boundaries and continues tocumulate during GB migration, eventually inhibiting gracoarsening through ‘‘impurity drag.’’ The resulting texturremains Zone II with columns extending through the filmbut with a lesser degree of preferred orientation and a smagrain size as shown in Fig. 5~b!.

With slightly higher oxygen concentration leve(JO/JAl;1022), coarsening during coalescence is seversuppressed, resulting in grains with random orientation. Tcompetitive growth which follows is governed by anisotropcrystallographic effects52 ~O segregates fastest at 111 sufaces!. Oxygen is incorporated into the lattice of 001 and 1crystal faces, while an oxide layer is formed on the 1faces.53 The oxide layer forms as oxygen tends to accumulat step edges on 111 surfaces, blocking step motion

FIG. 5. XTEM images with corresponding schematic diagrams showingmicrostructure of Al films deposited by thermal evaporation on amorphSiO2 at room temperature as a function of the incident O to Al flux raJO /JAl ~see Ref. 9!.

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forming step bunches. These pinning sites serve to nuclthe oxide phase. Neighboring 111 grains form rounded eddue to oxygen segregation while 111-oriented grains in ctact with 001 facets remain sharp, as oxygen is incorporain the latter. The 111 grains eventually develop roundedfaces indicating that local epitaxial Al growth has been intrupted by an oxide layer, above which renucleation of meislands takes place. Crystal growth on 001-oriented grainunimpeded by oxygen; these grains protrude above the aage film surface and eventually win in competitive growThey develop the shape of truncated octahedrons boundea 001 top face and 111 side faces. The degree of 001ferred orientation increases with film thickness and is accopanied by greater surface roughness with increasing oxyconcentrations@Fig. 5~c!#.

At still higher oxygen concentrations (JO/JAl;0.1– 1),the oxide layer completely covers islands of all orientatioat an early stage and coarsening during coalescencblocked. Thus, film growth proceeds by repeated renuation. The film is composed of 3D equiaxed~globular! grainswith random orientation and the Zone III structure@Fig. 5~d!#introduced in Ref. 1. With increasing oxygen concentratithe grain size decreases and can reach the nanometer rAn important byproduct of repeated nucleation for nanogrfilm formation is that surface faceting on individual columnand the related shadowing effects, are eliminated. Thnanophase films are inherently much smoother, and assult, denser. The presence of oxide phases also inhibitsmigration in the bulk of the film, preventing grain coarseniand imparting higher thermal stability.

O/Al has been used as a model system for investigathe growth of nanocrystalline (nc) grains separated by aamorphous tissue phase (a-AlOx in this case!, as indicated inFig. 5~d!, through continuous segregation and renucleatprocesses. This technique has been systematically explin order to synthesize ‘‘superhard’’ nanocomposite filmbased on transition metal~TM! nitrides and carbides, e.gnc-TMN/a-Si3N4 ,54 nc-TMC/a-C,55 TiNxBy andTiCxBy ,56 TMN/Cu,57 and YSZ/Au.58

As the oxygen concentration is further increase(JO/JAl;2 – 5), the role of the oxide and metal phasesreversed: the oxide phase nucleates first while Al segregto the surface and forms 3D islands.51,59 Resulting films arecomposed of metallic grains dispersed in an oxide ma@Fig. 5~e!#.60 Such composite films, consisting of a lowdiffusivity matrix with higher-diffusivity metallic inclusions,are the basis of a class of ceramic-metallic coatings wdiverse applications: resistors,61 sensors,62 solar cellelements,63 low-friction hard coatings~e.g., TM/a-C64!, andtribological coatings that adapt to the environment.65

At very high oxygen fluxes (JO/JAl@1), the films consistentirely of aluminum oxide, which for room temperatugrowth is amorphous.Ts values exceeding 800 °C are rquired for the synthesis of the chemically and mechanicstablek anda crystalline phases of alumina. There has, hoever, been a concerted effort to achieve hard crystalline

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Figures 6 and 7 illustrate two more examples of micstructural evolution in multi-component films. Both of thecases involve only metallic elements. Figure 6 is a plan-viTEM image of a 35-nm-thick Al14 at. % Pt film, depositedat Ts5350 °C. Phase identification by electron diffractioand dark field imaging67 shows that the larger grains witbrighter contrast are fcc Al while the smaller and darkgrains are a mixture of Al5Pt and Al6Pt. In contradistinctionto the 2D ‘‘tissue phases’’ formed in the Al/O system, thadspecies composing all three phases in this systemhigh surface mobilities and the minority phases (Al5Pt andAl6Pt) nucleate 3D islands on the surface of the majophase. Equiaxed grains of the minority phase decorate gboundaries and triple points, thus significantly decreasGB migration and grain coarsening.68 This type of micro-structural evolution has been observed in the Al~Cu! systemused for metallization in microelectronic circuits.37

The addition of Sn to Al has quite the opposite efferather than to decrease the grain size as O and Pt descabove, it acts as a grain size promoter.9 The plan-view TEMimage in Fig. 7, from a co-evaporated Al–Sn film with ain-plane compositional gradient, reveals a continuous

FIG. 6. Bright-field~BF! and dark field~DF! TEM images of a 35-nm-thickAl14 at. % Pt layer co-evaporated on amorphous C atTs5350 °C. The DFimage was obtained using Al5Pt and Al6Pt diffraction rings~see Ref. 68!.

FIG. 7. Bright-field plan-view TEM micrograph of a 100-nm-thick Al–Sfilm deposited on amorphous C at room temperature. The Sn concentrvaries from 0 at the left edge of the image to 10 at. % at the right edge~seeRef. 69!.

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crease in grain size from left to right. The Sn concentratvaries from 0 at the left edge of the micrograph to;10 at. %at the right edge. At low Sn concentrations, the film is cotinuous with an average grain size of 3–5 nm, whilehigher concentrations the film is still in the island growstage with the grain size enhanced by a factor of.5. Snappears to be acting as a surfactant in increasing the Alface mobility.69

IV. EFFECTS OF ION IRRADIATION ONMICROSTRUCTURE AND TEXTURE EVOLUTION

Low-energy ion irradiation during growth is used extesively to overcome the characteristically rough and unddense microstructures~Zones I and T! of refractory materialsdeposited at lowTs ~typically Ts /Tm,0.25).18–20,25,70Underthe correct set of deposition conditions, ion bombardmhas been shown to increase nucleation rates and film dento decrease average grain size, to inhibit the formationcolumnar structures associated with high surface roughnand to controllably affect the defect density and orientatof coatings. Because of their technological importance, Tand related TM nitrides have served as model systemstudy ion-assisted growth and we will use them as examto illustrate the roles of ion flux and ion energy on micrstructure evolution. For ease of discussion, we separateirradiation effects into three regimes characterized by thetio of the ion to metalJi /JMe fluxes incident at the growingfilm surface and the average ion energyEi : ~i! Ji /JMe<1with Ei<20 eV, ~ii ! Ji /JMe<1 with Ei>100 eV, and~iii !Ji /JMe>5 with Ei <20 eV.

A. Low-temperature film growth with low-energy,low-flux ion irradiation

In this section, we consider microstructural evolutionTiN and Ti0.5Al0.5N films deposited on SiO2 or steel sub-strates by reactive sputtering in Ar1N2 mixtures. The N2fraction in the sputtering gas mixture corresponds to sevpercent and is optimized to obtain stoichiometric films whmaximizing the deposition rate by maintaining a nonpsoned target surface. Under such conditions, the domiion incident at the growing film is Ar1 with energy corre-sponding approximately to the applied negative substbias Vs .71 That is, Ei5e(Vs2Vpl)'eVs , whereVpl is theplasma potential which is close to that of the grounded anunder typical magnetron sputtering conditions.

Figure 8 is a XTEM image illustrating the extremely opmicrostructure~Zone I! of TM nitrides deposited at lowTs .The micrograph is from the top portion of a 3-mm-thick TiNfilm grown at Ts5500 °C, with Ji /JTi50.5, Ei5100 eV,and a total pressurePt of 38 mTorr.72 Pt is approximately anorder of magnitude higher than typical values used in mnetron sputtering in order to thermalize the energetic partfluxes~sputtered and reflected! from the target. The thermalization distancedth of sputtered and reflected species is aproximately a factor of 5 smaller that the target substrseparationdts.73 Despite the relatively high ion energy, thmicrostructure is columnar and very porous with both int

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Magnetron sputter deposition is generally carried outpressures for whichdts,dth . The deposited flux is thus hyperthermal~average sputtered atom energies are;10 eVwith a high-energy tail extending to.100 eV!, resulting inZone T microstructures. Figure 9 is a XTEM micrograph oTiN film grown on amorphous SiO2 at Pt55 mTorr withTs5300 °C, Ji /JTi;1, and Ei;20 eV, the latter corre-sponding to the floating potential. The grain size, initiasmall, increases continuously with film thickness while t

FIG. 8. Bright-field XTEM micrograph obtained from the top portion of3-mm-thick TiN film deposited on steel at 500 °C with a total pressurePt

538 mTorr. The ion-to-Ti flux ratioJi /JTi incident at the film surface was0.5 with an ion energyEi5100 eV.

FIG. 9. Bright-field XTEM micrograph from a TiN film deposited on amophous SiO2 at 300 °C with a total pressurePt55 mTorr. The ion-to-Ti fluxratio Ji /JTi incident at the growing film was;1 with an ion energyEi

520 eV.

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column boundaries become increasingly more open.self-organized Zone T columnar microstructure formthrough random nucleation, limited coarsening during clescence, and competitive column growth. The microstrture is consistent with the one predicted by the kMC simlation in Fig. 4. The column tops are faceted due to kineroughening, which in combination with atomic shadowiresults in deep cusps between columns and open colboundaries. The individual columns, however, are densedicating sufficient adatom surface mobility to sustain locrystal growth. Sputtered atom energy contributes to loepitaxy.

A combination of x-ray diffraction, electron diffractionand lattice resolution TEM imaging was employed to invetigate TiN texture evolution under low-ion flux, low-energdeposition conditions (Ji /JTi;1 and Ei;20 eV).74 While111 and 002 grains, approximately 50% each by area, coeduring the early stages of film growth, the 111 grains graally overgrow the 002 oriented grains until at thickness.150 nm, the film has a nearly complete 111 texture. Resconsistent with those described for TiN, in which the highsurface-energy 111 grains emerge under low-temperalow-ion-irradiation deposition conditions, have been reporfor experiments with B1–NaCl-structure Ti0.5Al0.5N,75

ScN,76 TaN77 as well as for kMC simulations of fcc filmgrowth @see Fig. 4#.78 In each of these cases, the results cbe explained as being due to anisotropies in surface diffusties and adatom potential energies as discussed in Sec.

For the growth of TM nitrides by reactive sputter depotion, the interplay and competition among several interlated surface reaction and diffusional processes determfilm growth kinetics, surface morphology, film microstruture, and texture.79 The rate-limiting step for TM nitride filmgrowth is cation incorporation which, in the above expements, proceeds at a rate of several ML s21. The N2 supplyrate is much larger,&103 ML s21. If N2 molecules reach Tadatoms not yet incorporated in the lattice, they undergosociative chemisorption. However, all other incident N2 mol-ecules are only physisorbed and desorb at kinetic ratesestimated lifetimes& 1029 s. This results in~001! surfaceswhich exhibit an essentially bulk-terminated structure wequal cation and anion number densities while 111 surfaare fully N terminated.79 Ti adatoms form one N back bonon ~001! surfaces and three N back bonds on~111! surfaces.Thus, cation diffusivities and potential energies are higher001-oriented grains than on 111 grains and, followingarguments given above, Ti adatoms have a larger chancbecoming trapped at 111 sites. Consequently, 111 grslowly and inexorably expand at the expense of the 0grains under low-Ts , low-ion-flux growth conditions.

B. Effects of increasing Ei with J i ÕJ Me › 1

Increasing the ion energy, while maintainingJi /JMe < 1,affects the film microstructure, primarily through linear cacade effects. Figure 10 is an XTEM image of the middportion of a TiN film, deposited at a temperature and prsure (Ts5300 °C, Pt55.6 mTorr) similar to that of the

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sample shown in Fig. 9. During the growth of the film in Fi10, Ei was varied in steps of 40 eV withJi /JTi < 1.80 Themicrostructure ofEi<80 eV layers consists of dense coumns with open column boundaries. As the ion energyincreased to 120 eV, the voids along column boundariesappear and the film becomes fully dense. This is accomnied, however, by incorporation of intragranular residudamage~manifested in Fig. 10 as a darker average contrasthe Ei5120 eV sublayer! whose concentration increaseshigher voltages~even darker contrast in theEi5160 eV sub-layer!. When the energy is increased above 160–200 eV,defect density becomes so large that local epitaxial groon individual columns is disrupted and renucleation occuFigure 11, from the bottom portion of the film in Fig. 10shows high-energy (Ei5200– 400 eV) ion-irradiation-induced densification together with a highly defective, eqaxed microstructure which forms through renucleation. Tmicrostructure is similar to the one described in Sec. III,which a tissue phase forms due to segregation of a reaccomponent. The common feature in both cases is that locrystal growth on individual columns is disrupted by rpeated nucleation events.

The densification obtained in this regime (Ei.100 eV)comes at a steep price, however; the high ion energy resin correspondingly high defect densities, high compressstresses,81–83and inert gas incorporation.84 Ar concentrationsCAr in TiN layers deposited on amorphous SiO2 at 350 °C asa function ofEi , between 0 and 1800 eV, withJi /JTi < 1,are presented in Ref. 84.CAr is below 0.5 at. % withEi

,100 eV, while at higher ion energiesCAr increases ap-proximately linearly from 1 at. % at 200 eV to 5.5 at. %1800 eV.

Concurrent with the above results, x-ray and electron dfraction patterns reveal a change in TiN preferred orientat

FIG. 10. Bright-field XTEM micrograph obtained from the middle portiona 3.5-mm-thick TiN layer grown by reactive magnetron sputter depositionsteel at 300 °C with a total pressurePt55.6 mTorr. The ion-to-Ti flux ratioJi /JTi incident at the growing film was,1 while the ion energyEi wasvaried in steps of 40 eV from 400 to 0 to 400 eV.

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from 111 to 002 with increasingEi . However, the ion energyrequired to complete the transition is.800 eV, for which thefilms have unacceptably high stress levels. Thus, the ushigh-energy, low-flux ion irradiation is not a practical aproach for controlling film texture. The formation of 00texture under such conditions is directly related to colliscascade effects.18 Grains with open channel directions, suas 001, have higher survival probabilities due to the anisropy of collision cascades; that is, the ion energy is distruted over larger depths in open channels leading to losputtering yields and less lattice distortion. Increasing theenergy at lowJi /JMe values yields only a narrow interva~100–200 eV! over which densification is obtained with aceptable levels of radiation damage and gas incorporatio

Anisotropic collisional cascades effects,37 using ion-beamirradiation at off-normal incidence angle, have been shoto induce in-plane texture orientation in refractory metals85

oxides,86 and nitrides.87

C. Effects of increasing J i ÕJ Me with Ei › 20 eV

In this section, we discuss irradiation effects observwhen the ion flux is increased while maintainingEi

< 20 eV. At such energies, which are below the threshfor bulk lattice atom displacement in TM nitrides, the rsidual stresses remain low, yet the effects on texturemicrostructure are dramatic.

Using independent control88 of the energy and flux of ionsincident at the growing film, it has been demonstrated tthe reaction paths associated with independently varyingEi

andJi /JMe are quite different, even when the average kineenergy deposited per atomEMe&5Ei(Ji /JMe), is main-tained constant.75 In fact, varyingJi /JMe over a wide range,with Ei ' 20 eV, during magnetron sputter deposition hbeen shown to be an effective method for controllingmicrostructural evolution of polycrystalline NaCl-structu

FIG. 11. Bright-field XTEM micrograph obtained from the bottom portionthe film corresponding to Fig. 10.

J. Vac. Sci. Technol. A, Vol. 21, No. 5, Sep ÕOct 2003

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Ti0.5Al0.5N,75 TiN,74 ScN,76 and d-TaN,77 while introducingnegligible concentrations of residual ion-induced defectsresidual stress. In these studies, Ar/N2 sputtering gas mix-tures with N2 exceeding 10% were used, resulting in N2

1

fractions in the ion flux incident at the growing films o.0.05. We use TaN deposition77 as an example to illustratethe primary observations.

The v-2u and glancing-angle x-ray diffraction~XRD!scans from 500-nm-thick TaN layers grown atTs5350 °Cwith Ei520 eV andJi /JTa values ranging from 1.3 to 10.7contain predominantly~111! and/or ~002! d-TaN reflections~no other phases were detected!. Typical v-2u results areshown in Fig. 12~a!. Films grown with high flux ratios(Ji /JTa>7.4) exhibit a complete 002 texture while thosgrown with Ji /JTa, 6.3 have a strong 111 preferred orietation with a small volume fraction of 002, 022, and 11grains. XRD pole figures reveal, in addition, that all diffration peak maxima are aligned along the growth direction athat peak intensities are azimuthally symmetric. The msured residual stresses are low, ranging from tensile,11.4GPa, in 111-textured layers~low Ji /JTa values! to slightlycompressive,20.7 GPa, in 002-oriented films~high Ji /JTa

values!.The above results clearly demonstrate that the incid

FIG. 12. ~a! XRD v-2u scans from 500-nm-thickd-TaN layers grown byreactive magnetron sputter deposition on amorphous SiO2 at 350 °C as afunction of Ji /JTa with Ei520 eV. ~b! Normalized intensities of the~111!,~002!, ~022!, and~113! XRD peaks in~a! as a function ofJi /JTa ~see Ref.90!.

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S125 Petrov et al. : Microstructural evolution during film growth S125

ion-to-metal flux ratio can be used to selectively and contlably vary the preferred orientation of TM films from predominately 111 to 002. This is illustrated in Fig. 12~b! byplotting normalized hkl XRD peak intensities (I hkl /@ I 111

1I 0021I 0221I 113#) as a function ofJi /JTa.While XTEM images from the above-described 11

textured samples reveal an underdense Zone T microsture similar to the one presented above in Fig. 9, results002 d-TaN samples are very different. Figure 13~a! is abright-field XTEM micrograph from a film grown withJi /JTa510. The microstructure is still columnar, but in thcase fully dense in agreement with plan-view TEM resuand film density measurements.77 The surface is also mucsmoother than 111-textured layers. The inset in Fig. 13~a! isa selected area electron diffraction~SAED! pattern exhibit-ing strong 002 and weak 111 reflections along the grodirection. The dark field images in Figs. 13~b! and 13~c!,obtained using the circled portions of the 002 and 111 SAring segments along the film growth direction, show thmicrostructure evolution ind-TaN layers grown with highJi /JTa values also involves competitive growth. Both 00and 111 grains are present at film thicknesses up to.200nm. In this case, however, it is the 002 columns which wby overgrowing the 111 columns.

The above behavior can be understood based upon aments, presented in Sec. II, regarding competitive grobetween grains exposing low- versus high-diffusivity planto the growth front with the high-flux ion irradiation switching the balance between 111 and 002 orientations. Incid20 eV N2

1 ions are collisionally dissociated and providecontinuous source of atomic N which can chemisorb on 0grains~@001# is a nonpolar direction in the NaCl structure!,but not on N-terminated 111 grains. This is in contrastincident thermal N2 species which, as discussed above, hvery short lifetimes on both surfaces in the absence of a lopopulation of free cations. Thus, increasingJi /JTa corre-sponds to raising the steady-state N coverageuN on 001grains whileuN remains at a constant N-terminated value111 grains. This has the important consequence of decreacation mean free paths on~001! surfaces due to capture byadatoms to form TaNi ( i 51 – 4) admolecules which armore strongly bonded to the surface than Ta adatoms,therefore have lower surface mobilities. That is, in the prence of an atomic N source, the effective 001 adatom potial energy decreases below that on the 111. Consequethe net flux of cations from 002 to 111 oriented grainsreversed under high-flux conditions, resulting in the develment of 002 texture.

In addition to controllably switching the texture from 11to 002 with increasing ion flux, there is a correspondingcrease in layer density and a decrease in surface roughfrom underdensed-TaN layers with intercolumnar voids anself-organized growth mounds separated by deep surtrenches89 to fully dense layers with smoother surfaces. Tdensification is attributed to less pronounced kinetic rouing due to ion-irradiation enhanced surface mobilities resing in smoother surfaces with less atomic shadowing.

JVST A - Vacuum, Surfaces, and Films

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Increasing the ion flux from 1 to 10 withEi;20 eV re-sults in a transition from underdense 111 to a densetexture. In both cases, however, competitive texture evotion follows random nucleation on the amorphous substrIt is possible, however, to use texture inheritance37 to selectthe preferred orientation during the nucleation stage. Tconcept was utilized to achieve dense, highly 111-orienTM nitride ~TiN90 and TaN91! films. Ti underlayers 25 nmthick were first deposited on amorphous substrates at 8with Ji /JMe52 andEi511 eV. The layers are fully denswith a Zone II structure consisting of cylindrical grains haing very strong 0002 orientation. The TM nitride overlayewere then deposited with highJi /JMe values~.10! and Ei

;20 eV, which on bare amorphous substrates result in de002 layers. In this case, however, local epitaxy on 0002grains results in fully dense TM nitride layers with a stro111 texture and competitive grain growth is completeeliminated.

V. SUMMARY AND FUTURE DIRECTIONS

The atomic-scale understanding of microstructural evotion, necessary for the controlled manipulation of the proerties of thin film systems, is growing rapidly. However, ware still far from having an accurate quantitative descriptiA concerted program, with the ultimate goal being the abito design, at the atomic scale, new materials and new sttures having a particular desired set of properties, mightclude the following components.

~1! Detailed experimental measurements, for simmodel systems, of orientation-dependent adatom transparameters~activation barriers for surface diffusion, islanedge diffusion, step edge attachment/detachment, stepEhrlich barriers, and adatom formation energies! andorientation-dependent step and kink formation energThese measurements should be carried out on well-defisingle crystal surfaces as has recently been done by Kodbakaet al.92 for TiN.

~2! The above kinetic and thermodynamic parametshould then be incorporated into multiscale~in both time andspace! models of the early stages of microstructure develment including island growth and coalescence. The structevolution models will be based upon combinations of,example, density functional theory calculations, molecudynamics simulations, kinetic Monte Carlo simulations, acontinuum methods with accurate atomistically derivedformation driving each successive scale in the modeling.most difficult part is always the transition to the continuuHowever, the recent introduction of level set methods93 al-lows the input of atomistic information into a continuudescription of in-plane film growth while preserving the dcreteness of each atomic layer and dealing with local sinlarities in a natural manner.48

~3! Model predictions must be verified by complementain situ studies of film-growth dynamics utilizing, for example, scanning tunneling microscopy, low-energy electmicroscopy, and TEM. The experimental results shouldcompared to model predictions and, if necessary, fed b

J. Vac. Sci. Technol. A, Vol. 21, No. 5, Sep ÕOct 2003

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into a further round of model refinement. Recent advancethe development of aberration-corrected electron micscopes94 render these new-generation instruments mumore suitable for such experiments. They provide larworking volumes between the pole pieces for carrying ouinsitu film growth while maintaining atomic resolution imaging and providing ultrafast recording.

~4! The next stage is to extend both experiments amodeling/simulation to include coarsening during coalcence, competitive grain growth, and texture evolution. Twill require the addition of orientation-dependent graboundary energies into the models and accounting forrole of strain in both experiments and models.95

~5! A further extension is to include postdeposition agiand annealing to describe the thermal stability of thedeposited microstructural features to recrystallization andnormal grain growth, metastable multi-component96 thin filmsystems, and self-organized nanostructures~e.g., nanolamel-lae, nanocolumns, etc.!.

While portions of each of these tasks are being carriedat different laboratories around the world on a varietymaterials systems, it will likely require a focused effort bseveral laboratories working collaboratively to achieveabove goal.

ACKNOWLEDGMENTS

I.P. and J.E.G. gratefully acknowledge the financial suport of the Department of Energy, Division of Materials Scence, under Contract No. DEFG02-91ER45439 and theof the facilities of the Center for Microanalysis of MaterialUniversity of Illinois, which is partially supported by thU.S. Department of Energy. P.B.B. was supported byHungarian National Science Foundation OTKA, under Cotract Nos. 1225, 015878, and 030432, and the EU unContract No. ICAI-CT-2000-70029. L.H. acknowledges suport from the Swedish Research Council and the SwedFoundation for Strategic Research.

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