Titanium-nitride self-encapsulation of Cu and Ag films on silicon dioxide

Daniel Adamsa,*, T. Laursenb, T.L. Alfordc, J.W. Mayerb

aDepartment of Physics, University of the Western Cape, Private Bag X17, Bellville 7535, South AfricabCenter for Solid State Science, Arizona State University, Tempe, AZ 85287-1704, USA

cDepartment of Chemical, Bio and Materials Engineering, Arizona State University, Tempe, AZ 85287–6006, USA

Abstract

Encapsulation of Ag and Cu films via Ti nitridation of Ag(Ti) and Cu(Ti) alloy systems in an ammonia ambient has been investigated.Silver- and copper-titanium alloys on silicon dioxide of compositions varying from 4 to 27 at.% Ti were annealed at temperatures between350°C and 700°C, for durations of 10–120 min in a flowing NH3 ambient. Annealing of the Ag(Ti) and Cu(Ti) alloys at temperatures≥400°C, resulted in segregation of Ti to the surface to form a TiN(O) encapsulation layer and to the alloy/SiO2 interface. At the interface, Tireacted with the SiO2 to form a TiO/Ti5Si3 bilayer structure. Kinetic studies showed that Ti reactions take place within the first 10 min. Thisself-limiting behavior occurs for all annealing temperatures, but the encapsulation reaction increases with temperature. Four-point probeanalysis of the alloy films suggests that the resistivity is controlled by the residual Ti concentration. Comparable resistivity values wereobtained for Ag and Cu in the composition range considered. Resistivity values of ~2.6mQ-cm were measured in encapsulated Ag filmswith initial low Ti concentrations after nitridation at 600°C. 1997 Elsevier Science S.A.

Keywords:Encapsulation; Ag films; Cu films; Ti nitridation; Silicon dioxide

1. Introduction

The existing metallization schemes for ohmic con-tacts, gate metal and interconnections are found to be inade-quate for the development of ultra large scale integration(ULSI) and giga scale integration (GSI). These inadequa-cies include the reliability of aluminum and its alloys ascurrent carrier, susceptibility to electromigration and therelatively high resistivity of Al (~2.7mQ-cm) [1]. For thedevelopment of faster devices, the resistance-capacitance(RC) delay must be reduced. Advanced metallizationschemes using so-called multilevel metallization (MLM)structures of a low resistivity metal such as Cu have beenproposed to reduce the resistance component of the RCdelays [1].

Currently copper and silver, noted for low resistivitiesand higher resistance to electromigration, are being inves-tigated as future interconnect materials [2,3]. Despite thisthey have not found application in ICs because of their (a)high diffusivity and deep levels in silicon, (b) poor adhesionto SiO2 and polyimide, and (c) reactivity with the environ-

ment [1]. To make copper- and silver metallization manu-facturable, adhesion promoters, protection against corrosiveenvironments and the development of a process to defineinterconnection wiring will be needed [4].

The three most commonly used techniques [1] to enhancethe adhesion between the dielectric and copper are: (i) theuse of adhesion promoters; (ii) increased temperature ofdeposition or providing energy to ionize depositing species;(iii) surface pretreatment, especially plasma oxidation ofpolymers, sputter damage of metal surfaces or use of ionimplantation near the polymer or metal surface. The use ofadhesion promoters seems attractive since they can also actas diffusion barriers at higher process temperatures. Tita-nium, Ti-W and TiN have been most frequently usedbecause of titanium’s excellent chemical reactivity withoxygen, carbon, nitrogen, and fluorine [1].

Passivation of copper includes: (i) exposure of coppersurfaces to silane between 300 and 400°C to form a silicidedsurface [5], with the silicided surfaces providing corrosionprotection in air below 400°C; (ii) implantation of boroninto copper surfaces [6], which leads to a Gaussian depthprofile of boron in copper, leaving the top layers unpro-tected against corrosion; and (iii) alloying the top layersof copper with Ti, Pd and Al [7]. The compounds Cu3Ti,

Thin Solid Films 308–309 (1997) 448–454

0040-6090/97/$17.00 1997 Elsevier Science S.A. All rights reservedPII S0040-6090(97)00502-6

* Corresponding author.

Cu3Pd and CuAl2 are stable in oxidizing ambients; however,selective alloying is difficult and these compounds havevery high resistivities. It has been shown that a TiO2 over-coating of silver successfully protects the silver againstagglomeration in a chlorine ambient, and enhances the elec-tromigration resistance of Ag [8].

To accomplish both surface passivation and diffusionbarrier/adhesion promoter functions in a single processstep, it has been proposed to anneal Cu- and Ag-refractorymetal (Ti, Cr) alloy systems on SiO2 in an ammonia ambientto induce simultaneously a surface nitridation reaction andinterfacial reactions [7].

This investigation aims to contrast Cu-Ti and Ag-Ti alloysystems with respect to: (i) free surface and interfacial reac-tions of diffused Ti with the ammonia ambient and dielec-tric, respectively; (ii) thermal stability of the encapsulatedCu and Ag films; (iii) the effect of residual Ti on the resis-tivity of Cu and Ag; and (iv) the dealloying kinetics of Tiduring annealing.

2. Experimental

Alloy films consisting of ~200 nm Cu (27 at.% Ti), andAg(6–26 at.% Ti) were codeposited by electron-beam eva-poration onto thermally grown SiO2 (100–200 nm) on (100)Si substrates. The stoichiometry and thicknesses of all theas-deposited samples were determined by Rutherford back-scattering spectrometry (RBS). Samples were annealed for10–120 min at temperatures ranging from 300 to 700°C in aLindberg single-zone quartz-tube furnace in a flowing elec-tronic grade (99.99%, with H2O , 33 and O2 + Ar , 10molar p.p.m.) ammonia (NH3) ambient at atmospheric pres-sure, to form the refractory metal nitrides. Before eachanneal the annealing chamber was evacuated to about 10mTorr followed by a 2.5-min purge with NH3. Thissequence was repeated twice with a final 20-min purge.To minimize the chances of oxidation, flow rates of ~2–8l/min were maintained during the annealing.

Free surface as well as interfacial reactions were analyzedby RBS and Auger electron spectroscopy (AES). The RBSanalysis was performed using a 1.7 MV tandem acceleratorwith He+2 beam energies between 2.0 and 4.3 MeV. Thebackscattering angle was 170° and the total accumulatedcharge was 10–20mC. The samples were tilted at 7°. Weutilized the computer-simulation program RUMP for simu-lation and interpretation of RBS spectra [9].

After annealing, the sheet resistance of certain sampleswas measured by the four-point-probe method. The resistiv-ity was determined from the sheet resistance and thicknessof dealloyed Cu or Ag film only. Here, the measured resis-tivity was considered to be that of the metal layers. Thecontribution of the surface Ti-oxynitride layer resistivityto the total resistivity was less than 1%, because the ratioof the nitride thickness to Ag or Cu thickness was approxi-mately 1:10.

3. Results

3.1. Passivation of Ag and Cu films via Ti nitridation inNH3

All anneals in this section were performed for 30 min in aflowing NH3 ambient. Unless otherwise stated the condi-tions mentioned above will be implied. Fig. 1 comparesRBS spectra obtained from the as-deposited alloy withthat from alloys annealed at three different temperatures.Since the atomic masses of Cu and Ti are similar, theirbackscattering signals overlap partially at this beam energy.Simulation of the RBS data of the as-deposited sample indi-cates a composition of Cu(27 at.% Ti) and an alloy thicknessof 163 nm.

The sample annealed at 450°C differs from the as-depos-ited sample only by the presence of a thin surface oxidelayer of ~10 nm. The increased surface Ti peak suggeststhat at this temperature Ti only segregates to the free sur-face, presumably to react with oxygen and nitrogen from theambient. This outdiffusion of Ti makes the subsurface layerof the alloy Cu-rich. The samples annealed at 500°C and650°C clearly show the appearance of two separate Ti sig-nals, resulting from the segregation of Ti to the free surfaceand also to the alloy/SiO2 interface. The titanium at the freesurface is labeled as ‘Surface Ti’ and that at the alloy/SiO2

interface as ‘Interfacial Ti’. Due to the segregation of the Tito the free surface upon annealing, the surface position ofCu shifts to lower energies. At 500°C and higher tempera-tures the integrated surface-Ti signal decreases as comparedto the interfacial Ti signal. This increase in interfacial Ti ispossibly due to reaction between Ti and the underlying SiO2

which competes favorably with the surface reaction in thistemperature regime. At 650°C less than 10% of the Ti seg-regates to the free surface as indicated by the small surface

Fig. 1. 2.0-MeV He+2 spectra obtained from as-deposited Cu(20 at.% Ti)alloy on SiO2 and from the alloy annealed at 450, 500 and 650°C for 30min in NH3. Since the atomic masses of Cu and Ti are close to each other,their backscattering signals overlap.

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Ti peak in Fig. 1. Simulation of the RBS data shows that theCu-rich layer contains less than 4% of residual Ti.

The Ti that diffused to the free surface reacts with theammonia and residual oxygen to form a titanium-nitridelayer, is indicated as TiN(O) in the schematic of Fig. 1.At the alloy/SiO2 interface, the Ti dissociates the SiO2 andsubsequently reacts with the freed Si and O to form aninterfacial Ti-oxide/Ti-silicide bilayer structure. Theselayers are labeled as ‘TiOw’ and ‘Ti5Si3’, respectively.

Fig. 2 compares the RBS spectrum of the as-depositedAg(19 at.% Ti) alloy with that nitrided at 450 and 600°C.After a 450°C anneal, the presence of a ‘Surface Ti’ peakand a distinct ‘Interfacial Ti’ peak indicates that Ti segre-gated to the free surface and also reacted with the SiO2

substrate. The surface and interfacial reactions again resultin the formation of a TiN(O) layer and Ti-oxide/Ti-silicidebilayer structure, respectively. Computer RUMP simulationof the spectrum corresponding to the 450°C anneal, suggeststhat the dealloyed Ag layer contains a residual Ti concen-tration of ~10.8 at.%. The TiN(O) thickness is ~17 nm.Anneals at 600°C result in a more dealloyed Ag layer,with a residual Ti concentration as low as 0.9 at.%.

3.2. Time-dependent dealloying of Ti from Cu- and Ag-based alloys

The dealloying kinetics of Cu(20 at.% Ti) alloys werestudied. The RBS spectrum of the as-deposited alloy iscompared with those annealed 500°C for 30, 60 and 120min (Fig. 3). Ti segregates to the free surface and the alloy/

Fig. 2. RBS spectra showing the depth distributions of Ag and Ti of a 210-nm-thick Ag(19 at.% Ti) alloy, before and after annealing at 450 and600°C for 30 min in NH3. A 2.0-MeV He+2 beam energy was used.

Fig. 3. RBS spectra showing only the depth distributions of Ti for a Cu(20at.% Ti) alloy before and after annealing for 30, 60 and 120 min at 500°Cin NH3 ambient. The spectra were obtained using a 4.3-MeV He+2 beamand a scattering angle of 170°.

Fig. 4. Dealloying kinetics obtained with Cu(20 at.% Ti) alloy films. Theresidual Ti concentration is shown as a function of annealing time for threedifferent temperatures. The annealing took place in a NH3 ambient and thedata were obtained using 4.3-MeV He+2 RBS.

Fig. 5. Dealloying kinetics obtained with Ag(26 at.% Ti) alloy films. Theresidual Ti concentration is shown as a function of annealing time for threedifferent temperatures. The annealing took place in a NH3 ambient and thedata were obtained using 2.0-MeV He+2 RBS.

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SiO2 interface, with a slight preference to the interface. Thespectra of the annealed samples differ only slightly fromeach other, giving a residual Ti concentration of ~5.5at.%. The time-dependence of the dealloying process wasdetermined by measuring the residual Ti concentration inthe alloys after nitridation. The dealloying process is char-acterized by an initial rapid dealloying, followed by a muchslower rate (Fig. 4). For equivalent Cu(Ti) and Ag(Ti) alloycompositions, almost similar residual Ti levels wereobtained. At 500°C for 120 min, the residual Ti concentra-tion for the Ag(19 at.% Ti) alloy is ~4.8 at.%, and ~5.0 at.%for the Cu(20 at.% Ti) alloy.

Fig. 5 shows the residual Ti concentration as a function ofannealing time for an Ag(26 at.% Ti) alloy. It is again clearthat for all temperatures rapid dealloying occurs within thefirst 10 min, thereafter little or no further outdiffusion of Titakes place. On the other hand, a strong temperature depen-dence is evident. Ag(Ti) alloys with initial Ti concentrationsof 6, 10 and 19 at.% showed similar dealloying behavior.

3.3. Effect of Ti alloying on the resistivity of encapsulatedCu and Ag films

The effect of Ti on the resistivity of Cu was studied byannealing Cu(Ti) alloys in NH3 at three different tempera-tures and times. The residual Ti concentration and thicknesswere measured by RBS, and the sheet resistance by the four-point-probe method. The resistivity of a Cu(20 at.% Ti)alloy, nitrided at 400, 450 and 500°C for times varyingbetween 30 and 120 min is depicted in Fig. 6. The datasuggest that the resistivity is strongly temperature depen-dent. At a given temperature, the resistivity decreasesrapidly from the very high as-deposited value (~111.6mQ-cm) to some lower value, and thereafter it decreases at amuch slower rate. For the Cu(Ti) alloy, annealed at 500°C,the resistivity only varies from 8.1 to 5.4mQ-cm in the timeinterval 30–120 min. Fig. 7 shows the resistivity as a func-tion of residual Ti concentration for the Cu(20 at.% Ti)

alloy, nitrided at 350–700°C for 10–120 min. It is clearthat the resistivity of the Cu film increases with the amountof residual Ti. By using a second order polynomial fit, extra-polation of the resistivity to zero residual Ti, gives a value of5.5 mQ-cm. Measured resistivities after annealing correlatewith the residual Ti concentration. The ratio of resistivityincrement to the impurity concentration (Dr/c) for Ti is 8.6mQ-cm/at.% for dilute substitutional Cu alloys [10], which ismuch larger than the slope of the line tangent to the curve ofresistivity versus residual concentrations, 0–15 at.% Ti (Fig.7). Therefore, it is expected that only a small fraction of theTi is substitutional.

Resistivity versus annealing time for an Ag(19 at.% Ti)alloy, annealed at three different temperatures is shown inFig. 8. It is clear that the resistivity drops rapidly within thefirst 10 min from the high value (~109.0mQ-cm) of the as-deposited sample to ~8mQ-cm at 500°C. The initial rapiddrop seems to be strongly temperature dependent and theresistivity change is much slower for longer annealingtimes. This behavior is observed for alloy concentrationsof 6–26 at.%.

4. Discussion

4.1. Passivation of Cu and Ag films via Ti nitridation inNH3

The results obtained from annealing the Cu(27 at.% Ti)/SiO2/Si structure in an NH3 ambient at different tempera-tures for 30 min indicate that Ti segregates out of the Cu toboth the free surface and the alloy/SiO2 interface. A two-layer Ti-nitride/Ti-oxide surface layer is formed at the freesurface due to reaction of Ti with the NH3 ambient andresidual oxygen. The large negative heat of formation ofTiO2 (242.5 kcal/g-mol at 900 K) compared to that of TiN(91.4 kcal/g-mol) is compensated by the high partial pres-

Fig. 6. The variation of resistivity with annealing time is depicted for aCu(20 at.% Ti) alloy annealed at 400, 450 and 500°C in a NH3 ambient.

Fig. 7. The resistivity as a function of residual Ti concentration is shownfor a Cu(20 at.% Ti) alloy, nitrided at different temperatures in NH3. Theresidual Ti concentration was measured by RBS.

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sure of NH3 relative to that of O2 [11], leading to nitriderather than oxide formation.

Previous studies [12] have shown that the Ti:N ratio ofthe TiN layer approaches 1.0 for temperatures greater than550°C, and the nitride contains ~6 at.% O [10]. The amountof nitrogen incorporated increases with temperature with anactivation energy of 0.6± 0.1 eV. The presence of the sub-surface Ti-oxide layer at all temperatures≥450°C, furthersuggests that Ti is the dominant diffusing species duringnitridation. The relatively low temperature required toform TiN from Cu-Ti in an ammonia ambient comparedto N2 [13] could be due to the difference in thermal decom-position energy between the two gas molecules. The decom-position energies of N2 and NH3 are 942 and 432 kJ/mol,respectively [14].

The interfacial reaction starts with the dissociation ofSiO2 in the presence of Ti according to:

11Ti+3Si02 → Ti5Si3 +6TiO (1)

Electron diffraction analysis (not included) revealed thatthe reaction results in the formation of a TiO/Ti5Si3/SiO2

stack in the range of 500–700°C. This is consistent with theresults previously reported for the reaction of pure Ti onSiO2 [15]. The high oxygen solubility of Ti (2.3–34 at.% Oat 800°C) ensures the massive dissociation of SiO2 andsubsequently the diffusion and the interaction of oxygen,metal and silicon atoms. It has been shown that an inter-mixed Cu(Ti, Si, O) layer is present at the TiO/Ti5Si3 inter-face as a result of Cu diffusion across the TiO layer to reactwith Ti5Si3 [16].

Nitridation annealing of Ag-Ti alloys above 300°Cresulted in Ti segregating at the surface as well as at theinterface. The reaction products obtained from the nitrida-tion of the Ag(Ti) alloys are similar to those for Cu(Ti)alloys. At the surface, Ti reacted with NH3 and residualO2 to form a TiN(O) layer, and at the interface with SiO2

to form an oxide-silicide bilayer structure.The data indicates that the Ti-nitride thickness obtained

from nitridation of the Ag(Ti) alloys increases moderatelywith temperature in the range 300–600°C, but reaches afinite thickness (~20 nm) at higher temperatures. Theamount of Ti available for reaction is controlled by thedealloying mechanism, as reflected by the relationshipbetween the residual Ti concentration and annealing tem-perature [3].

In the case of Ag-Ti the Ti/SiO2 reaction occurs at tem-peratures as low as 350°C [3], compared to Cu-Ti alloys forwhich the onset of this reaction was observed to be≥450°C.Earlier studies have indicated that significant reactionbetween pure Ti on SiO2 only occurs at temperatures≥600°C [15]. AES depth profiling analysis supported theRBS analysis that nitridation of the Ag-Ti alloys results inthe formation of a Ti-oxide/Ti-silicide structure at the alloyand SiO2 interface.

RBS analysis indicated no segregation of Ag to the Ti-oxide/Ti-silicide interface for an Ag(19 at.% Ti) alloynitrided in NH3. Calculations based on the heats of reactionusing the heats of formation of binary alloys [16], yieldedpositive values for the change in enthalpy in all cases. Thatis, there is no thermodynamic driving force to initiate theAg-Ti5Si3 reaction (in contrast to the favorable Cu-Ti5Si3reaction observed in Cu(Ti) alloys). This explains the sta-bility of the silver in contact with the interfacial layers.

4.2. Dealloying of Ti from Cu- and Ag-based alloys

Time-dependent dealloying curves for Ag(Ti) and Cu(Ti)alloys (Figs. 4 and 5) are all characterized by an very rapidinitial drop in the Ti concentration within the first 10 minfollowed by a plateau, which is associated with a smalldealloying rate. The residual Ti showed a strong tempera-ture dependence, namely, at higher temperatures lower Ticoncentrations are obtained. Zou et al. have reported a simi-lar behavior in the formation and kinetics of titanium nitridein Ag/Ti bilayers (Y.L. Zou, T.L. Alford, Y. Zeng, F. Deng,S.S. Lau, T. Laursen, A.I. Amali, and B.M. Ulrich, Forma-tion and Kinetics of Titanium Nitride in Ag/Ti Structures,unpublished).

These studies showed that the nitride growth is linear(x ∝ t) in the range 0–15 min and parabolic (x ∝ tl/2) for15–120 min. The latter kinetic model implies that a diffu-sion-controlled layer is likely to be the rate-limiting processgoverning the nitridation reaction. The nitride growth canbe limited by diffusion of the reagents (NH3 or Ti). Forexample, Ti diffuses faster through an Ag(100 nm)/Ti(50nm) bilayer than Ag(200 nm)/Ti(50 nm).

The alloy results further suggest that the rate-limiting stepin the encapsulation is not the mass transport and reactionswithin the encapsulating layers. Otherwise, substantiallylower residual Ti concentrations would be obtained inAg(6 at.% Ti) compared to Ag(26 at.% Ti) alloy. Afterbeing annealed at 500°C for 60 min, these alloys containedresidual Ti concentrations of 3.2 and 6.8 at.%, respectively.More dealloying in Ag(6 at.% Ti) alloy would have resulted

Fig. 8. The resistivity as a function of annealing time is shown for anAg(19 at.% Ti) alloy, nitrided at different temperatures in NH3.

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if the rate-limiting step was in the TiN formation. Thekinetics are instead likely to be controlled by the releaseand transport of the refractory metal in the Ag or Cu film.The factor of 3 increase in grain size, observed after a 500°Canneal is not a significant enough change to consider Timass transport by grain-boundary diffusion as the limitingprocess. However, the effect of this increase in grain sizehas a significant effect on the grain boundary volume avail-able to accommodate the Ti atoms upon annealing. If aspherical shape is assumed for the grains, then the expres-sion for the grain boundary volume fraction (GBVF) isgiven by:

GBVF=3d

r(2)

where the grain boundary width,d ~0.5 nm [17], andr isthe grain radius. For grain sizes of 28 and 77 nm obtainedfrom TEM (not shown) the GBVF is 11 and 4%, respec-tively. Given initial alloy concentrations of 6–26 at.% Ti, itis obvious that upon annealing the Ag cannot accommodatethese large concentrations without precipitation or outdiffu-sion. Given an enhanced Ti diffusivity in supersaturatedgrain boundaries a large amount of Ti is therefore availableat the free surface upon annealing. This scenario is consis-tent with data for the Ag(6 at.% Ti) where only ~50% of theinitial Ti is dealloyed compared to 74% for the Ag(26 at.%Ti) alloy.

4.3. The effect of Ti alloying on the resistivity ofencapsulated Cu and Ag films

After being annealed above 450°C in NH3, the Cu-Ti andAg-Ti alloys consist of a TiN(O) passivation layer, a Cu (orAg) layer and a Ti-silicide/Ti-oxide bilayer at the alloy/SiO2

interface. It is assumed that this structure resembles a par-allel-resistor configuration, that the conductor (Cu or Ag)has the lowest resistance (Rcond) and it is also the thickestlayer. Therefore, 1/Rcond q 1/Rpassivation (Rpassivation, resis-tance of encapsulation) and 1/Rtotal~1/Rcond or Rtotal~Rcond

(after annealing). This assumption simplified the calculationof the resistivity and resulted in an error of less than 1%.

For both alloys, the resistivity of the nitrided samples ishigher than the elemental values. The lowest resistivitiesobtained after nitridation for the Cu(Ti) and Ag(Ti) alloyannealed at 500°C for 30 min are ~5.4 and 7mQ-cm, respec-tively. The higher than elemental resistivity values observedare believed to be attributed to the following two factors.

Firstly, the presence of residual Ti in the encapsulated Cuand Ag films has a great effect on the resistivity. Time- andtemperature-dependent studies of the dealloying processindicated incomplete dealloying. Even alloys with initialcompositions as low as 4–6 at.% contained up to 1% resi-dual Ti after nitridation at ~600°C. The data suggest thatlowering the initial Ti concentration would not necessarilylower the resistivity. Marecal et al. have demonstrated theeffect of grain size on the resistivity of Ag films sputter-

deposited on biased glass and silicon substrates [18]. Agrain size change from 28 to 38 nm resulted in a resistivitychange from 3.7 to 2.2mQ-cm for Ag on glass and from 3.5to 1.6 mQ-cm for Ag/Si. It has also been shown that theresistivity of a TiO2-passivated Ag layer decreased withannealing temperature due to increased grain size, from 50nm for as-deposited sample to 300 nm after annealing at600°C [8]. It is believed that annealing of the Cu(Ti) andAg(Ti) alloys in NH3 initiated microstructural changes andother competing reactions within the first 10 min, inhibitingthe dealloying.

A second factor that may influence the resistivity is theformation of intermetallics as a result of the reactionbetween the Cu (or Ag) and Ti. However, glancing-angleX-ray diffraction analysis of the Ag-Ti and Cu-Ti alloys,could not verify intermetallic formation in either system. Ithas been shown that nitridation of Cu(Ti) alloys results insegregation of Cu to the Ti-oxide/Ti-silicide interface.Therefore, the effective thickness of the Cu layer is reducedand this also led to an increase in resistance of the encapsu-lated Cu films. Comparison of RBS data of the Ag and Cusystems indicated no such segregation of Ag to the Ti-oxide/Ti-silicide interface.

5. Conclusions

Nitridation of Cu-Ti and Ag-Ti alloys on SiO2, in NH3,resulted in a multilayer structure consisting of a TiN(O)surface layer, a dealloyed Cu (or Ag) layer, and an inter-facial Ti oxide-silicide bilayer. The evolution of the finalstructure is therefore governed by a competition between thefree surface nitridation/oxidation and the interfacial reac-tion.

The dealloying behavior for both systems was character-ized by a rapid decrease in residual Ti concentration withinthe first 10 min, followed by a much slower diffusion rate.The residual Ti showed a strong temperature dependence; athigher temperatures lower Ti concentrations are obtained. Itis evident that the rate-limiting step in the encapsulation isnot the mass transport and reactions within the encapsulat-ing layers.

Resistivities.2 mQ-cm have been obtained from theencapsulated Ag and Cu films. The main cause of the higherthan elemental resistivities is due to the incomplete deal-loying that occurs during the nitridation. The relationshipbetween the residual Ti concentration and resistivity indi-cated that the former is controlled by the dealloyingmechanism.

Acknowledgements

The work is partially supported by the National ScienceFoundation (L. Hess, DMR-9624493, DMR-9307662) towhom the authors are greatly indebted.

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