Encapsulated silver for integrated circuit metallization

Encapsulated silver for integrated circuit metallization

Daniel Adamsa,*, T.L. Alfordb

aDepartment of Physics, University of the Western Cape, Private Bag X17, Bellville 7535, South AfricabDepartment of Chemical and Materials Engineering, Arizona State University,

NSF Center for Low Power Electronics, Tempe, AZ 85287-6006, USA

Accepted 21 March 2003

Abstract

Silver has all the properties to complement copper for multilayer metallization in integrated circuit

technology. The obvious drawbacks to Ag in terms of agglomeration and corrosion are overcome by use of

encapsulation layers. The objective of this issue is to provide readers with a survey of some active areas in Ag-based

metallization for interconnects ULSI applications. This work reflects the current research progress in this field

from several different perspectives and provides advanced metallization solutions for future integrated circuit

manufacturing.

# 2003 Elsevier Science B.V. All rights reserved.

Keywords: Silver; Electromigration; Encapsulation; Stress; Dielectrics; Resistivity

1. Introduction

As the complexity of multilayer metallization (MLM) increase, the performance limitingresistive and capacitive signal delays increase accordingly. The development of advanced ultra-large-scale integration (ULSI) and giga-scale integration (GSI) technologies will place stringent demandson future interconnect and metallization schemes [1]. To decrease the resistance–capacitance (RC)signal delays, the circuit can be fabricated with a metal having resistivity lower than currently usedAl(Cu) alloy [2]. As a result, higher current densities can be imposed on the metal lines and fasterswitching speeds can be achieved, due to the lower RC time delay [3].

Silver has the lowest resistivity of all metals, and its high oxidation resistance differentiates itfrom aluminum and copper. On the issue of electromigration, silver promise to provide anelectromigration resistance of at least one order of magnitude better than Al [4]. It has been shownthat Cr and TiO2 overcoatings of silver, greatly enhance the electromigration resistance of Ag [4,5].These properties of Ag make it one of the promising high-conductivity candidates to be consideredas possible interconnect material for ULSI technology. However, before it can be used for thispurpose, there are several issues that need to be addressed in realizing Ag interconnects.

Silver does not reduce SiO2 and is therefore expected not to adhere well to SiO2 surfaces [6]. Itsuffers also from electrolytic migration along surfaces in wet atmospheric environments [4] andsilver has a high diffusivity in SiO2 [7]. A major concern regarding Ag metallization is its

Materials Science and Engineering R 40 (2003) 207–250

* Corresponding author. Tel.: þ27-21-959-3473; fax: þ27-21-959-3474.

E-mail address: [email protected] (D. Adams).

0927-796X/03/$ – see front matter # 2003 Elsevier Science B.V. All rights reserved.

doi:10.1016/S0927-796X(03)00025-1

susceptibility to corrosion in the presence of weak oxidizing agents such as sulfur [8]. The corrosionand agglomeration of Ag in high Cl ambients have also been reported [6,8]. Thereforeimplementation of silver as interconnect material will require adhesion promoters to improve theadhesion to dielectrics; passivation layers to protect it against corrosive environments, and thedevelopment of a process to define the interconnection wiring.

Fig. 1 is a cross-section of a three metal level interconnect-structure using Ag as the conductivematerial. The interlevel dielectrics (ILD) could be conventional SiO2-based materials or moreideally, materials such as Pa-n (or polyimide). If conventional SiO2 is used, then Ag plugs andinterconnects have to be enclosed in diffusion/drift barriers so that Ag will not move into Si or SiO2

under thermal stress or biased temperature stress (BTS). Fig. 2 shows a sectional view of a multilevelAg metallization-insulator structure fabricated for a NLT-HP [4].

As devices continue to shrink, both film thickness and widths of metal interconnects are alsoreduced; hence, the effect of thermal stability becomes more significant because it affects the device

Fig. 1. Cross-sectional diagram of a two metal level interconnect-structure, using Ag as the conductor.

Fig. 2. Sectional drawing of multilevel metallization-insulator for the NLT-HP as fabricated with Ag and Cu compositemetals [4].

208 D. Adams, T.L. Alford / Materials Science and Engineering R 40 (2003) 207–250

reliability. Ag films have been reported to agglomerate at moderate temperatures under oxygencontaining ambient [9–12]. It has also been reported that the degree of agglomeration dependsspecifically on test conditions, for example, ambient and film thicknesses [9–12]. Agglomeration isthe result of atoms and voids diffusion causing a surface restructure during annealing. As the filmmorphology evolves voids and hillocks are formed. Changes in sheet resistance of the films can beused as a measure of formation and growth of voids and hillocks [13].

Fig. 3 shows the in situ sheet resistance of Ag (60 nm)/SiO2, annealed in air, N2 and vacuumambient at 0.15 8C/s ramp [13]. It is evident that the annealing ambient significantly affects filmstability. The sample annealed in air seems to be more susceptible to thermal instability and hence arapid change in sheet resistance occurs. Samples annealed in vacuum appear to be the mostthermally stable and a more gradual change in sheet resistance occurs once it deviates fromlinearity. The scanning electron microscope (SEM) micrograph in Fig. 4 depicts the surfacemorphology of a 75 nm thick Ag layer on SiO2 annealed at a ramp rate of 0.6 8C/s at 450 8C in air[13]. As seen from the micrograph, the film consists of a uniformly distributed connected-networkof islands.

Fig. 3. In situ sheet resistance as function of temperatures for Ag (60 nm)/SiO2 annealed at 0.15 8C/s in air, N2 andvacuum [13].

Fig. 4. SEM micrograph of Ag/SiO2 annealed at a ramp rate of 0.60 8C/s at 450 8C in air [13].

D. Adams, T.L. Alford / Materials Science and Engineering R 40 (2003) 207–250 209

2. Formation of encapsulation

2.1. Formation of encapsulation from different Ag–metal systems

In view of the thermal stability issues discussed in the previous section a passivation layer anddiffusion barrier are required to protect the silver from the fabrication environment, and an adhesionpromoter is needed to enhance adhesion of Ag to the dielectrics.

Titanium–nitride is used in integrated circuit technology as a diffusion barrier and etch stop. Ithas low specific resistance when compared to TiW [6], which makes it particularly attractive inapplications where device speed is critical. Furthermore, TiN has been found to suppress or preventaluminum migration through TiSi2 to the silicon substrate [14]. This makes TiN a viable material forthin-film diffusion or reaction barrier layer. Two common methodologies can be used to produceTiN, i.e. (1) reactive sputtering of titanium in a nitrogen ambient [15] and (2) titanium deposition andsubsequent annealing in either nitrogen or ammonia [14]. The reactive sputtering in the nitrogenambient has suffered from an incomplete control of compositions since the sputter rate changes as afunction of time. Therefore, we have focused on forming a TiN passivation layer in an ammoniaambient to encapsulate Ag films [16,17].

To accomplish both surface passivation and diffusion barrier/adhesion promoter functions ina single process step, it has been proposed to anneal Cu-refractory metal (Ti, Cr) alloy andbilayer structures on SiO2 in an ammonia (NH3) ambient to induce simultaneously a surfacenitridation reaction and interfacial reactions [18]. This process has been applied to Ag-refractorymetal systems and is schematically represented in Fig. 5. The surface reaction results in theformation of a titanium–nitride encapsulation layer (labeled as ‘TiNx’) and a titanium-oxide/titanium-silicide (‘TiO/TiSiy’) bilayer structure which is a direct consequence of the interfacialreaction.

2.1.1. Copper–titanium and silver–titanium alloys

The Ag (or Cu)–titanium depositions were prepared on Si wafers with 210 nm thick thermallygrown oxides. The alloy structures consisting of �200 nm thick Ag (or Cu)–Ti alloys ofcompositions varying from 6 to 26 at.% films were made by dual electron-beam deposition. Allanneals were carried out in a fused-quartz tube furnace using a flowing NH3 ambient of 99.99%purity and flow rate of 2 l/min at atmospheric pressure.

Fig. 5. Schematic of nitride self-encapsulation of silver-bilayers and -alloys. The encapsulation process simultaneouslyprovides a surface protection layer and an interfacial adhesion layer.


The encapsulated Ag films have been characterized by depth profiling: Rutherfordbackscattering spectrometry (RBS), Auger electron spectroscopy (AES) and secondary-ion-massspectroscopy (SIMS). Glancing angle X-ray diffraction (XRD) and four-point-probe resistancemeasurements were used to determine the film’s structural and electrical properties.

The RBS spectrum of an as-deposited Cu (20 at.% Ti) alloy is compared with that annealed at500 8C for 120 min (Fig. 6). Upon annealing the Ti segregated to the free surface and the alloy/SiO2

interface, with a slight preference to the surface. The Ti that diffused to the surface reacted with theammonia and residual oxygen to form a titanium–nitride layer, indicated as TiN(O). The interfaciallayers are labeled as ‘TiOw’ and ‘Ti5Si3’, respectively. The spectrum of the annealed sample gives aresidual Ti concentration of �5.5 at.%.

Fig. 7 compares the RBS spectrum of an as-deposited Ag (19 at.% Ti) alloy with that nitrided at450 and 600 8C [19]. After a 450 8C anneal, the presence of a ‘surface Ti’ peak and a distinct‘interfacial Ti’ peak indicates that Ti segregated to the free surface and also reacted with the SiO2

substrate. The surface and interfacial reactions again result in the formation of a TiN(O) layer andTi-oxide/Ti-silicide bilayer structure, respectively. Computer RUMP [20] simulation of the spectrumcorresponding to the 450 8C anneal, suggests that the dealloyed Ag layer contains a residual Ticoncentration of �10.8 at.%. The TiN(O) thickness is �17 nm. Anneals at 600 8C resulted in a moredealloyed Ag layer, with a residual Ti concentration as low as 0.9 at.%.

2.1.2. Silver–titanium bilayersA bilayer structure corresponding to �15 at.% Ti were prepared by sputter-deposition of a

22 nm thick Ti film followed by 120 nm Ag without breaking vacuum. The encapsulation took placeat 300–700 8C in a fused-quartz tube furnace using a flowing NH3 ambient of 99.99% purity andflow rate of 2 l/min at atmospheric pressure. The characterization and analysis techniques were thesame as that used in Section 2.1.1.

The encapsulation of the Ag-refractory bilayers proceeded in a similar way to that of the alloys.Alford et al. [16] demonstrated the encapsulation of the Ag-refractory bilayers using RBS (as show

Fig. 6. RBS spectra showing only the depth distributions of Ti for a Cu (20 at.% Ti) alloy before and after annealing for120 min at 500 8C in NH3 ambient. The spectra were obtained using a 4.3 MeV He2þ beam and a scattering angle of 1708.


in Fig. 8). Only the Ti and Ag backscattered signals are displayed, and the Ti signal shows clearlythat the 500 8C anneal causes Ti to segregate at the surface (see peak labeled ‘Ti surface’). Based onthe RBS analysis, the TiN(O) surface layer is 10 nm thick. The interfacial bilayer is expected toconsist of TiOw (w � 1:1) and Ti5Si3, the reaction products from the Ti–SiO2 reaction [21].

Zou et al. performed [22] depth profiling to obtain the elemental distribution and investigate theinterfacial reactions. AES depth profiles of Ag (100 nm)/Ti (50 nm) bilayers both before and afterthe nitridation reaction at 600 8C are shown in Fig. 9a and b [22]. The crossovers in Fig. 9a indicate

Fig. 7. RBS spectra showing the depth distributions of Ag and Ti of a 210 nm thick Ag (19 at.% Ti) alloy, before and afterannealing at 450 and 600 8C for 30 min in NH3. A 2.0 MeV He2þ beam energy was used [19].

Fig. 8. Ag encapsulation on SiO2 prepared from Ag (120 nm)/Ti (22 nm) bilayers. RBS spectra show the depthdistributions of Ag and Ti before and after annealing at 500 8C in an ammonia ambient for 30 min. The spectra wereobtained using a 2.0 MeV He2þ beam and a scattering angle of 1708 [16].


interfaces in the Ag (100 nm)/Ti (50 nm) bilayer structure on an oxidized silicon substrate. The Siand O concentrations in the as-deposited film are nearly zero and are within the noise levels. Depthprofiles after annealing at 600 8C for 30 min are shown in Fig. 9b. Significant peak-to-peak heightsof Ti, N and O during the first minutes of sputtering indicates that Ti nitride formed at the surfacewith oxygen present, is labeled as TiN(O) in Fig. 9b. The oxygen in TiN(O) can be attributed toresidual oxygen in the reaction chamber. Although the N (KLL) and Ti (LMM) Auger signalsoverlap at 380 eV, the formation of Ti nitride enhances the Ti peak intensity. For this reason, thesurface Ti nitride is labeled ‘Ti þ N’ in Fig. 9b. At the interface between Ti and SiO2, pure Ti reactswith SiO2 to form a Ti-oxide (TiOw) and a Ti-silicide (Ti5Si3) bilayer [23]. The reaction promotes theadhesion of Ag to the underlying SiO2.

Zeng et al. [24] used cross-sectional transmission electron microscopy (XTEM) to confirm thefinal structure of the nitrided samples based on RBS and AES results. Various layers are identifiedand labeled on the electron micrographs (Fig. 10) for samples before and after anneals at 600 8C. ATi-nitride film of about 25 nm thick is formed on top of the Ag layer after anneals. By comparing thecross-sectional view before and after anneals, a reduction in the Ti layer thickness as well as a

Fig. 9. Auger depth profiles of the Ag (100 nm)/Ti (50 nm) bilayers (a) as-deposited and (b) nitrided at a temperatureof 600 8C in NH3. The surface Ti nitride is labeled as Ti þ N due to the overlap of the N (KLL) and Ti (LMM) signals[22].


contrast difference at the Ti/SiO2 interface is observed. This supports the observations made fromRBS and AES.

The silver grains tend to be equiaxed in the as-deposited samples, and appear more oriented andextend through the thickness of the Ag film after anneals. In the as-deposited sample viewed byXTEM, the grain sizes are ranging from 0.03 to 0.2 mm, which are increased from 0.1 to 0.4 mm after600 8C annealing. Plane-view TEM (not shown) was also used to observe lateral grain growth. Theaverage grain sizes are 0.08 and 0.17 mm for the as-deposited samples and those annealed at 600 8C,respectively. It is noted that the Ti–SiO2 interface appears relatively uniform and shows no voids ordelamination at the interface [22]. This is a positive indicator of Ti acting as an adhesion promoter.

Zeng [25] also used XRD to confirm the texture of Ag films in Ag/Ti bilayers. Silver filmsdeposited on Ti layers exhibit a strong h1 1 1i texture (Fig. 11) [25]. After the nitridation process, the

Fig. 10. Cross-sectional TEM micrographs obtained from the Ag (100 nm)/Ti (50 nm) bilayers (a) as-deposited and (b)nitrided at 600 8C for 30 min in NH3 [24].

Fig. 11. X-ray diffraction pattern of the Ag (200 nm)/Ti (50 nm) bilayer structure before and after nitridation at 600 8C for30 min in NH3 [25].


texture of Ag films is further enhanced. The improved texture is accomplished by a grain growthduring the thermal annealing. The near bamboo structure with preferred orientation and large grainsis desirable to improve the electromigration resistance of silver metallization.

The results obtained from nitridation of Cu(Ti) and Ag(Ti) alloys and Ag–Ti bilayers indicatedthat in the presence of ammonia (NH3) and residual oxygen (O2), Ti diffuses to the free surface toform a titanium–nitride encapsulation layer, indicated as TiN(O). It has been shown that the Ti:Nratio of the TiN layer approaches 1.0 for temperatures greater than 550 8C, and the nitride contains�6 at.% O [26]. The lower decomposition energy of NH3 (432 kJ/mol) compared to N2 (942 kJ/mol)makes it possible to form TiN at relatively lower temperatures [27]. At the alloy- or bilayer-dielectric(SiO2) interface, the Ti dissociates the SiO2 and subsequently reacts with the freed Si and O to forman interfacial Ti-oxide/Ti-silicide bilayer structure or adhesion/diffusion barrier layer. Consistentwith the results from the reaction of Ti with SiO2, the interfacial reaction products were identified asthe TiOw/Ti5Si3 bilayer stack formed in the temperature range 500–700 8C [28]. In contrast to hightemperature (�600 8C) reaction for the Ti/SiO2 system, the Ti–SiO2 reaction in the Ag (or Cu)/Tibilayers and alloys occurs at temperatures in the range 350–450 8C.

2.2. Nitride growth kinetics

2.2.1. Ag–Ti bilayersZou et al. [22] investigated the influence of the initial Ti or Ag thickness on the nitride growth

using two different experimental configurations: Ag (100 nm)/Ti (50 nm) and Ag (200 nm)/Ti(50 nm) bilayers. The nitridation reaction was carried out at three temperatures (400, 500 and600 8C) and four times (15, 30, 60 and 120 min). An ammonia flow rate of 2 l/min was maintainedso that the partial pressure of NH3 in the reaction chamber is sufficiently high to supply the nitridingspecies.

The TiN thickness (x) obtained from RBS measurements was plotted as a function of annealingtime for times ranging from 15 to 120 min. The best fit to the experimental data is shown in Fig. 12by the parabolic function, x / t1=2, since the coefficient of determination is nearly one [22]. Thestraight lines in Fig. 12 are least-square fits of the experimental data to the parabolic law. Typicaluncertainties in the RBS results are �3% for areal densities and a few tenths of atomic percent forthe stoichiometric ratios [29]. As the thickness is linearly proportional to the areal density, theuncertainty in the film thickness (Dx/x) is also �3% based on large signals counts (200,000) obtainedin the study.

The nitride growth is described by x2 ¼ Bt after t > 15 min, where B is a parabolic rate constantfor the growing nitride phase [30]. The rate constants were determined by taking squares of theslopes of the plots in Fig. 12a and b [22]. It is evident from the results that the nitride growth rates forthe Ag (100 nm)/Ti (50 nm) are much higher than those in Ag (200 nm)/Ti (50 nm) bilayers. Giventhat the only difference in these two bilayer structures is the thickness of the Ag films, it implies thatthe diffusion of Ti through silver plays an important role in the nitride formation.

The parabolic rate constants are plotted against the reciprocal of temperature in Fig. 13 for Ag(100 nm)/Ti (50 nm) and Ag (200 nm)/Ti (50 nm) bilayer systems [22]. Based on the least-square fitof the slopes in Fig. 13, the activation energy (Ea) was determined from the Arrhenius plot. Theactivation energies are Ea ¼ 0:31 � 0:06 eV for Ag (200 nm)/Ti (50 nm) and Ea ¼ 0:44 � 0:10 eVfor Ag (100 nm)/Ti (50 nm) structures. As the parabolic rate constant is defined by B ¼ x2=t andDx=x ¼ �3%, the experimental uncertainty of the rate constants ðDB=BÞ ¼ 2Dx=x ¼ �6%. Zou et al.[22] suggest that the same mechanisms operate in Ag (100 nm)/Ti (50 nm) and Ag (200 nm)/Ti(50 nm) bilayer structures. The kinetics data suggest that the nitridation reaction takes place in two


steps, namely, a fast growth during the first 15 min followed by a much slower reaction at timesgreater than 15 min. The large offsets shown in Fig. 12, confirm the initial fast growth rate whichfollows a x / t relationship, compared to the parabolic dependence x / t1=2 for times >15 min.Based on the parabolic relationship, the nitridation reaction is limited by the rate at which either NH3

or Ti diffuses to arrive at the TiN–Ag interface. The experimental data after 15 min anneals suggestthat the TiN is indeed limited by the diffusion of the titanium since the supply of nitriding species isin abundance. These arguments support a linear-parabolic relationship (x2 þ Ax þ Bt) to describe thekinetics of the nitridation reactions in the Ag/Ti bilayer systems, where x is the nitride thickness at athe time t, and A and B are constants [31]. It was shown that for a parabolic growth, the parabolic rateconstant (B) is linearly proportional to the diffusivity (D) of the moving species [32].

The values of the activation energy obtained by Zou et al. [22] are 0.31–0.44 eV, which agreewell with the reported values (0.4 eV) for surface diffusion on silver [5]. Based on the Boltzmanndiffusion relationships, and literature values [5,33], the diffusivities at 600 8C are DLS � 1�10 12 cm2/s, DGB � 2 � 10 6 cm2/s and DSD � 1 � 10 6 cm2/s, for lattice diffusion (LS), grain

Fig. 12. Ti-nitride thickness as a square root function of time for two bilayer systems: (a) Ag (200 nm)/Ti (50 nm) and Ag(b) (100 nm)/Ti (50 nm) [22].


boundary (GB) diffusion and surface diffusion (SD), respectively, in or on Ag. As the surfacediffusivity is smaller than that in the GB diffusion, the limiting step for the nitride growth is believedto be the surface/interface diffusion on top of the silver films.

2.2.2. Ag–Al bilayers

2.2.2.1. Passivation layer formation. Wang and Alford [34] demonstrated the encapsulation of silverfrom Ag–Al systems. Bilayer structures of 100 and 200 nm Ag on 8 nm Al were depositedsequentially by electron-beam evaporation onto thermally oxidized Si(1 0 0) substrates withoutbreaking vacuum. The preparation and annealing procedures were the same as for the Ag–Tibilayers. Self-encapsulation anneals were performed for 30 min at temperatures ranging from 400 to725 8C.

Upon annealing aluminum diffuses through the silver layer to the free surface where it reactswith the NH3 and the residual O2 to formed a thin layer of Al oxynitride on the surface. Fig. 14compares the RBS spectrum of the sample nitrided at 725 8C with the as-deposited one [34].Although the surface segregation of Al as denoted by the ‘‘Al peak’’ in Fig. 14, was apparent for allthe annealed samples, only those annealed for temperatures greater than 675 8C showed significantnitrogen and oxygen peaks as well.

From RUMP [20] simulations it followed that the Al-oxynitride layers contained �35 at.%oxygen and the thicknesses were about 10 nm. The thermodynamic data supported these results sincethe Gibbs free energy for consumption of one mole of Al to form Al2O3 at 725 8C is much lower, 606.079 kJ/mol ( 6.292 eV/Al atom), than that to form AlN, 217.923 kJ/mol ( 2.262 eV/Alatom) [35]. It is hence expected that oxygen-rich, Al-oxynitride films form at this temperature rangeinstead of pure Al-nitride layers.

Wang and Alford showed that during X-ray diffraction analysis, the shift of Ag diffraction peakcorrelates with the residual concentration of Al atoms in the silver layer after the encapsulationprocess [34]. The authors reported that a lower residual Al concentration was obtained for hightemperature anneals, and that it is indicated by the peak shifts dependence on annealing temperature.

Fig. 13. Temperature dependence of parabolic rate constants for the Ag (200 nm)/Ti (50 nm) and Ag (100 nm)/Ti (50 nm)bilayer structures [22].


Fig. 15 shows the Ag(1 1 1) peaks versus annealing time at 725 and 500 8C, respectively, for thesample described above [36]. The vertical bar indicates the position where the original Ag(1 1 1)peak should be. It is observed that at 725 8C, the Ag peaks gradually shifted back towards theoriginal direction, while at 500 8C, no apparent movement was observed. This comparison confirmsthe previous expectation that the higher annealing temperature accelerates Al diffusion through Agfilms. It is also noticed that the h1 1 1i peaks are stronger at 725 8C than at 500 8C for all the

Fig. 14. RBS spectra (3.7 MeV He2þ, 708 tilt) obtained from a Ag (200 nm)/Al (8 nm) bilayer on oxidized Si substratebefore (dashed line) and after (solid line) for a 30 min anneal at 752 8C in flowing NH3 [34].

Fig. 15. The shift of Ag(1 1 1) X-ray diffraction peaks with anneal time at different temperatures: (a) 725 8C and (b)500 8C [36].


annealing times. This is in agreement with the previous observation that the higher annealingtemperature (725 8C) enhances the h1 1 1i fiber texture of the Ag films [34].

2.2.2.2. Modeling of the Al migration in Ag–Al systems. Wang et al. [37] developed a model toexplain the Al migration during the encapsulation process. A schematic diagram depicting thediffusion of Al through Ag and the resulting encapsulation process are shown in Fig. 16 [37]. Theproposed diffusion model is based on the following:

1. The encapsulation is a direct consequence of plane sheet diffusion within a limited volume.2. Accumulation of Al occurs at the Ag–Al oxynitride interface (shown in Fig. 16 as the Al atoms

trapped in the bottom of the energy valley), requires that the customary diffusion equation bemodified as follows [38]:

@C

@t¼ D

@2C

@x2 @S

@t(1)

where S is the trapped Al concentration, C the concentration of the free-diffusing Al atoms.If a constant self-diffusivity of Al (D) instead of grain-boundary diffusivity is employed, the

diffusion equation becomes

@C

@t¼ D

R þ 1

@2C

@x2(2)

where the constant, R ¼ S=C [38]; which relates to the chemical effects governing the diffusionof Al in Ag. R equal to 500 gives a good fit with the experimental data in Fig. 17.

3. Assume a uniform initial concentration (C2) in the Ag film.4. The accumulation of the Al atoms at the Ag–Al oxynitride interface and subsequent diffusion to

the surface retards the depletion of Al from the Ag matrix and imposes an additional boundarycondition. This retardation mechanism is controlled by the interfacial energy between the newlyformed Al-oxynitride barrier and Ag layer.

Fig. 16. A Schematic depiction of the Ag/Al bilayer encapsulation process. The retardation mechanisms are representedby the trapped Al atoms (solid circle at the bottom of the valley) and the interfacial energy barrier DE at the reactionfront [37].


Based on these assumptions and by considering the total amount of Al transported across the Aglayer instead of calculating the Al concentration profile C(x), the ratio of transported Al in time t(M(t)) to that after infinite time (M(1)), is given by

MðtÞMð1Þ ¼ 1

X1

n¼1

2L2 exp½ b2nDt=l2ðR þ 1Þ�

b2nðb2

n þ L2 þ LÞ(3)

The results calculated from this expression is depicted in Fig. 17 for anneals at 500 and 725 8C,respectively [37]. The very good correlation between the experimental and calculated results (Fig. 17)supports the derived segregation model of Al in Ag. The model confirms the effect of temperature onthe encapsulation process and is in agreement with previous observations that the higher theannealing temperature the higher is the depletion and the lower the probability of inter-metallicAg–Al compound formation.

In Fig. 17, a plateau appears at times >10 min for the 725 8C anneal, which reflects the self-passivating characteristics of the Al-oxynitride films. These characteristics of the Ag/Al bilayerencapsulation process ensures good control on the thickness of diffusion barriers, and will benefitfuture ULSI technology when multiple Ag metallization layers are employed in integrated circuitdevices.

The results from the nitridation of the Ag/Al bilayers in flowing N2 þ NH3 were used to derivea model to explain the kinetics of the self-encapsulation process as a combination of a competitivebehavior between the diffusion in Ag and the trapping of Al atoms at the Ag–Al oxynitride interface.Excellent agreement between the theoretical and experimental results was obtained.

2.3. Electrical properties of encapsulated Ag films

2.3.1. Nitridation in ammonia ambientThe effect of Ti on the resistivity of Ag was studied by annealing Ag(Ti) alloys and Ag–Ti

bilayers in NH3 at different temperatures and times. The thickness of the encapsulated Ag films wasmeasured by RBS and, the sheet resistance by the four-point-probe method.

Fig. 17. Transportation ratio of Al atoms through Ag layer annealed at 725 8C (solid circle) and 500 8C (open circle),respectively, in the flowing gas mixture of N2 and 1% NH3. The overlapped curves are the calculated values from Eq. (3)for 725 8C (dashed line) and 500 8C (solid line), respectively [37].


The resistivity of a Ag (19 at.% Ti) alloy, nitrided at 400, 450 and 500 8C for times varyingbetween 30 and 120 min is depicted in Fig. 18 [37]. The results show that the resistivity dropsrapidly within the first 10 min from the high value (�109.0 mO cm) of the as-deposited sample to�8 mO cm at 500 8C. The initial rapid drop seems to be strongly temperature dependent and theresistivity change is much slower for longer annealing times. This behavior is observed for alloyconcentrations in the range 6–26 at.% Ti in Ag.

Fig. 19 shows the resistivity versus annealing temperature for a Ag (120 nm)/Ti (22 nm) bilayer.Upon annealing the sample at 450 8C, the resistivity increases from a value of 3.1 mO cm(as-deposited) to 3.5 mO cm. For the temperature range 450–600 8C, the resistivity decreases to3.2 mO cm (at 600 8C). It has been shown that substantial dealloying of Ti from Ag(Ti) alloys tolevels below 1 at.% is necessary to obtain resistivities below 3 mO cm. From RBS analysis of the Ag(120 nm)/Ti (22 nm) bilayer it was determined that the Ti accumulations after annealing did notexceed the detection limit of �1 at.%.

The 3 mO cm resistivity level in the case of bilayers is therefore, not controlled by the residualTi concentration, but by the limitation of the deposition process. Ag films with near-bulk resistivityhave been prepared by sputter-deposition [39].

The resistivity measurements showed that the resistivities of the alloys after annealing arecontrolled by the residual Ti concentration. Strong temperature dependence is apparent but isprimarily caused by the residual Ti concentration. Resistivity values of <5 mO cm suggest that thebulk elemental phase and not an intermetallic phase (TiAg or Ti2Ag) [40] is the dominant filmcomponent. Phase identification was done by glancing angle X-ray diffraction (results not shown),only the Ag(1 1 1), Ag(2 0 0) and Ag(2 2 2) signals were observed in both the unannealed andannealed samples.

Wang and Alford [34] reported the variation of resistivity versus annealing time of Ag/Albilayers annealed at 500 and 725 8C, respectively (Fig. 20). It is evident that for both temperatures,the resistivity increases rapidly within the first few minutes. For the 500 8C anneal the resistivityremains almost constant after the initial increase. After the initial increase in the resistivity of thesamples annealed at 725 8C, it decreases dramatically to some constant value. The resistivity reaches

Fig. 18. The resistivity as a function of annealing time is shown for an Ag (19 at.% Ti) alloy, nitrided at differenttemperatures in NH3 [37].


a constant value when the Al atoms stopped segregating to the surface or the residual concentrationis a minimum. At 500 8C the higher resistivity is due to the higher residual Al concentration.

2.3.2. Effect of different annealing ambients on electrical properties of Ag/Al bilayers

Both RBS and AES analyses of the Ag/Al bilayer indicated that Al segregates to the free surfacewhen annealed at different temperatures in Ar, He–H and NH3, respectively. For the three ambients,

Fig. 19. Room-temperature resistivity of encapsulated Ag films prepared from Ag (120 nm)/Ti (22 nm) bilayers. Theresistivity was obtained from four-point-probe and RBS thickness measurements. The Ti concentration was in all casesbelow the RBS detection limit, i.e. (<1 at.% Ti).

Fig. 20. Variations of Ag resistivity values vs. annealing times at different temperatures and with different Ag thicknesses:(a) Ag (100 nm) at 500 8C (open circle); (b) Ag (100 nm) at 725 8C (solid circle); and (c) Ag (200 nm) at 725 8C(open triangle) [34].


Al accumulates in the Ag during annealing (Fig. 21) [43]. The accumulated Al versus annealingtemperature profiles is very similar for the three ambients.

Fig. 22 shows the resistivity as a function of annealing temperature for the Ag (200 nm)/Al(8 nm) bilayer annealed in three different ambients (Ar, He–H and NH3) [43]. Annealing the samplesin Ar at temperatures ranging from 300 to 700 8C shows that the resistivity is higher than that

Fig. 21. Plot of Al concentration vs. annealing temperature for three different ambients [43].

Fig. 22. The Ag resistivity of Ag (200 nm)/Al (8 nm) bilayer on SiO2 vs. annealing temperature for three differentambients [43].


annealed in He–H and ammonia, respectively. The higher resistivity in this case is due to the higheramount of residual Al in the Ag layer.

For the samples annealed in He–H and NH3 the resistivity remains almost constant fortemperatures up to 300 8C; whereafter, it increases linearly to a maximum value of about 6 mO cm at500 8C. The resistivity then decreases to the value of the as-deposited samples at 700 8C. It is evidentthat annealing the Ag/Al bilayers in these ambients gives rise to silver films with almost the sameresistivity. The resistivity of the samples annealed in Ar does not show the plateau for temperatures<300 8C, but rather a linear increase from room temperature.

The behavior of the resistivity with annealing temperature resembled that observed for the Alconcentration versus temperature (Fig. 21) [43]. Therefore, it seems that the accumulated Alconcentration dictate the resistivity values. It is clear that the resistivity of the Ag films increaseswith the amount of Al that remains in the Ag film after annealing.

ULSI applications will require that the encapsulated Ag films have low resistivities, comparableto the elemental resistivity of Al (2.7 mO cm) and approaching those of Cu (1.7 mO cm) and Ag(1.6 mO cm) [41]. The resistivity’s dependence on the Ti concentration is weaker than expected fromdilute substitutional-alloy theory [42]. The resistivity contribution from the Ti impurities is primarilyonly a problem for the alloys.

Results therefore suggests that highly conductive Ag films can be prepared from alloys andbilayers by optimizing the process variables associated with the film deposition and theencapsulation process. However the high processing temperatures for the self-encapsulation processwill have to be avoided if Ag is to become more attractive for future interconnect schemes.

2.4. Thermal stability of encapsulated silver films

2.4.1. Effectiveness of TiN(O) diffusion barrier

Based on RBS analysis it followed that annealing a TiN(O) encapsulated Ag film in H2S atvarious temperatures for 30 min, no corrosion occurs up to 300 8C (Fig. 23) [44]. The sample

Fig. 23. Ag (19 at.% Ti) alloy was nitrided at 600 8C, 30 min in NH3. The RBS spectra show depth distributions of thenitrided alloy, after annealing at 300, 400 and 500 8C for 30 min in H2S [44].


surfaces retained their smoothness, and golden color, and no difference between the sample at300 8C (in H2S) and the nitrided (no sulfur anneal) one could be detected. After a 400 8C anneal, theRBS spectrum displays the following features: (1) a reduction in the height and thickness of the Agsignal; (2) a weak S peak near channel 240; (3) shift of the interfacial Ti signal to higher energies;and (4) a broad tail. The first three features point to the formation of a silver-sulfide compound as aresult of Ag migration through the Ti-nitride to react with the H2S at the free surface. However, thebroad tail on the spectrum indicates the presence of a discontinuous Ag surface layer. Annealing thenitrided sample at 500 8C for 30 min in H2S results in an almost complete diffusion of Ag to thesurface. The two Ti signals (interfacial Ti and surface Ti) merge as a result of this outdiffusion. Amuch larger S peak is also present.

The evolution of the corrosion of encapsulated Ag films is depicted schematically in Fig. 24[44]. After a 400 8C anneal, Ag diffuses through the Ti-nitride layer to form Ag2S crystallites, on thesurface. The well-defined sulfur peak on the spectrum of the sample annealed at 500 8C, suggest thatthe crystallites are Ag covered with a 133 nm thick layer of Ag2S. If the crystallites consisted only ofAg2S nonuniformly distributed over the surface, then the sulfur peak would have been smeared out.At 500 8C, nearly all the Ag had migrated to the surface to react with H2S. It seems that theTi-nitride barrier fails by allowing diffusion of Ag through it, without itself being dissolved. It isbelieved that the diffusion occurs along grain boundaries and possibly microvoids, pinholes, cracks andweak spots in the barrier layer. The fact that the TiN passivation layer essentially remains intact uponfailure is also an indication that neither Ti (nor N) react with Ag or the H2S ambient at this temperature.Therefore, the migration of Ag to the surface is driven by the thermodynamically favorable reactionbetween Ag and H2S.

2.4.2. Effectiveness of AlxOyNz diffusion barrierWang and Alford [34] evaluated the effectiveness of the AlxOyNz encapsulation layer as

diffusion barrier. This was done by depositing a 50 nm thick Cu film on the AlxOyNz/Ag structurefollowed by a high temperature anneal in both vacuum and forming gas. Fig. 25 shows the RBSspectra of the Cu/AlxOyNz/Ag structure before and after annealing at 620 8C for 30 min [34].Analysis of the data revealed that no interdiffusion occurred between the Cu and Ag. These resultswere the same for anneals in both vacuum and forming gas. Compared to the results obtained fromthe Ag/TiN system [45], the AlxOyNz diffusion barrier is stable up to 620 8C, which is about 200 8Chigher than the TiN processed in a similar manner.

Fig. 24. Schematic showing the nitridation and evolution of the corrosion process of a Ag (19 at.% Ti) alloy. At 600 8C, inNH3, the Ti forms TiN at surface and interfacial TiO/Ti5Si3 bilayer structure to encapsulate the Ag. At 400 8C, in H2S,some Ag diffuses through TiN to form faceted crystallites; at 500 8C most of the Ag has moved to the surface to react withthe H2S ambient [44].


3. Microstructural analysis

3.1. The effect of metal-underlayer on texturing in Ag films

X-ray diffraction patterns obtained with the Bragg–Brentano geometry (y–2y scan) for Ag/Al,and Ag/Cr samples, before and after annealing are shown in Figs. 26 and 27, respectively [34,36].Fig. 26 shows that no Ag peaks other than strong {1 1 1} and {2 2 2} reflections appear in thepatterns of the as-deposited Ag (100 nm)/Al (30 nm) bilayer. This is not the case for the Ag/Cr

Fig. 25. RBS spectra (3.7 MeV He2þ, 78 tilt) of the diffusion barriers before (broken line) and after (solid line) testing at620 8C for 30 min in flowing atmospheric N2–5% H2 ambient [34].

Fig. 26. y–2y X-ray diffraction spectra from several Ag (100 nm)/Al (30 nm) bilayers annealed at: 725, 675, 625 8C andas-deposited. All the Ag peaks are indexed, except the one denoted as (0 0 4) which is due to the silicon substrate [34].


bilayers. All Ag diffraction peaks listed in the Joint Committee for Powder Diffraction Standard(JCPDS) card appear when Cr is the underlayer. Ag on Ta (not shown) also displays a similar patternto that of Ag on Cr. These results seem to suggest strong h1 1 1i texture in the as-deposited Ag(100 nm)/Al (30 nm) bilayer, but an almost random orientation for the Ag films in Ag/Cr and Ag/Tastructures. Enhancement of the {1 1 1} peak in the Ag/Al bilayer at 725 8C, further suggests that thetexture in Ag is promoted upon annealing. In contrast, as shown in Fig. 27 [36], most of the Ag peakintensities of the Ag/Cr samples do not change significantly with an exception in {1 1 1} orientation.A 25% increase of peak intensity is observed after the annealing process. This indicated that randomoriented grains dominate with slight texture growth in h1 1 1i direction during anneals. These resultsindicate that Ag film texture has metal underlayer dependence.

3.2. Quantifying the Texture—Texturing Model

Zeng et al. [46–48] extensively investigated the evolution of crystallographic texture and stressin Ag thin films during the encapsulation process. The crystallographic texture was characterized byX-ray diffraction using a Philips materials research diffractometer. The pole figure technique wasused. It has proven to be one of the most informative X-ray techniques for texture characterization inthin films [49]. In this technique, the incident and diffracted beams make the same angle with thefilm surface, which equals to the Bragg angle corresponding to the specific diffraction plane ofinterest {h k l} (Fig. 28) [46]. The diffraction intensity data of this plane were collected as the sampleis rotated in its own plane along the normal corresponding to each fixed tilt angle c (Fig. 28). Therotation angle f ranged from 0 to 3608 at an interval of 38, while the tilt angle c varied from 0 to 808with an interval of 18. The recorded intensity data after background subtraction were corrected forgeometric defocusing and sample absorption and then normalized to the average intensity. The{h k l} pole figure was mapped by plotting the normalized intensity data on a stereographicprojection, with the projection plane parallel to the film plane and the intensity values expressed inmultiples of the average intensity. Plotting the pole intensity as a function of tilt angle, which isreferred to as a {h k l} fiber plot, also simplified the description of the texture [46,48]. The fibertexture can be quantified in terms of the volume fraction and distribution of each texture component.According to Jetter et al. [50] both the influences of the angle c and the multiplicity of diffraction

Fig. 27. X-ray diffraction of Ag (200 nm)/Cr (50 nm) bilayers before and after nitridation at 600 8C for 30 min in NH3

[36].


planes on the diffraction intensity need to be taken into account when calculating the volumefraction. For a homogeneous two-component fiber texture, the volume fraction associated with eachcomponent can be calculated by the equation [47]:

Fc ¼PR c2

c1InormðcÞ sinc dc

2Pc

R p=2

0InormðcÞ sinc dc

(4)

where c1–c2 is the angular range attributed to the texture component in Inorm(c) sin c versus c plot,Pc/P is the fraction of the multiplicity of the selected diffraction plane associated with this range. Inthis study the Jetter et al. method was used to quantify the texture [50].

Zeng et al. [47] illustrate schematically a hypothetical two-component fiber texture hu1v1w1iand hu2v2w2i in Inorm(c) sin c versus c plot. It is assumed that the two peaks correspond to thehu2v2w2i component and the third one corresponds to the hu1v1w1i component. According to Eq. (4),the volume fraction of each component is proportional to the area under the peak(s) due to thiscomponent, with the corresponding multiplicity considered. In order to consider the volume fractionof random grains in this case, the minimum intensity as indicated by the arrows in Fig. 29 wassubtracted before Eq. (4) was used [47]. The volume fraction of random grains was obtained bydividing the minimum integrated intensity by the overall integrated intensity under the curve inFig. 29. Fig. 30 shows the Ag{1 1 1} fiber texture plots of three encapsulated Ag (200 nm)/Ti 50 nm)samples [47]. The h1 1 1i peaks around c ¼ 0� were truncated due to their extremely highintensities. The high h1 1 1i intensity along the fiber axis (c ¼ 0�) in all samples indicates a strongh1 1 1i texture in Ag films. Moreover, three other peaks were observed at 38.9, 56.3 and 70.58 tiltswith respect to the fiber axis (Fig. 30). In terms of the angle relations between {1 1 1} and othercrystallographic planes in a cubic system, these peaks were identified as h5 1 1i texture, with thethird peak at 70.58 coinciding exactly with the h1 1 1i texture component.

Zeng et al. [46] show that the two h5 1 1i peaks at 38.9 and 56.38 in {1 1 1} fiber texture plots(Fig. 30) were combined to calculate the volume fraction of the h5 1 1i texture. The correspondingPc/P ratio is equal to 3/8. The various peak widths o was used to represent the distribution of aspecific texture component along the fiber axis. Although the change in peak width may also becontributed to the internal stress and grain size, these two factors can be ignored in the present case.Upon annealing, the residual stress in Ag films is much higher than that in the as-deposited films dueto the thermal mismatch between Ag films and the substrate [46]. The peak width is expected tobroaden if this stress is nonuniform. In the present study the observed narrowing of peaks afterannealing is not attributed to the internal stress. There is also a grain size range from 10 (or 1) to0.1 mm, where X-ray diffraction is quite insensitive to variation in grain size [51]. The measured

Fig. 28. Schematic diagram showing the X-ray diffraction geometry used for the texture measurement of Ag films [46].


grain sizes in the as-deposited and annealed Ag films (approximately 0.15–0.35 mm) fall into thisrange, therefore the effect of the grain size on the peak width is negligible.

3.3. Texture development

The {1 1 1} pole figures of Ag films annealed at 600 8C for 120 min are schematically depictedin Fig. 31 [48]. Since fiber textures exhibit rotational symmetry about the fiber axis, only one quarterof the pole figure is indicated. For the as-deposited Ag films (not shown), high {1 1 1} intensityaround the fiber axis (c ¼ 0�) suggests a strong h1 1 1i texture. It is evident that the h1 1 1i texture issubstantially enhanced after 600 8C annealing and displays a sharper distribution about the fiber axiswhen compared to that in the as-deposited Ag films. The pole intensities around c ¼ 39 and 568 are

Fig. 29. Schematic representation of I(c) sin c vs. c plot for a hypothetical two-component fiber texture (hu1v1w1i andhu2v2w2i). I(c) is the normalized diffraction intensity of specific {h k l} plane and c is tilt angle. This plot is used tocalculate the volume fraction of each texture component [47].

Fig. 30. {1 1 1} fiber texture plots of Ag films showing the normalized {1 1 1} intensity as a function of tilt angle c for Ag(200 nm)/Ti (50 nm) samples annealed at 400, 500 and 600 8C for 30 min [47].


believed to be the contribution of h5 1 1i texture, while the pole intensity centered at c ¼ 70:5� isdue to both h1 1 1i and h5 1 1i textures. The presence of h5 1 1i texture in the Ag films results fromthe twining of h1 1 1i-oriented grains [47]. The pole intensity about c ¼ 56� is also partiallyattributed to the newly formed h2 0 0i texture for the sample annealed at 600 8C. It is evident that atelevated temperatures the h5 1 1i texture becomes weaker whereas the h1 1 1i texture is enhanced.

Zeng et al. [48] also used (in plane-view) TEM in conjunction with the pole figures to analyzeAg/Ti bilayers microstructure after vacuum anneals (Fig. 32a and b). The average grain sizesmeasured were 80, 84 and 173 nm, respectively, with considerable fast grain growth observed at600 8C. In addition, the silver grains in the as-deposited samples tend to be equiaxed, while thegrains of the samples annealed at 600 8C appear more oriented and extend through the entire Agthickness. These observations were also confirmed by selected area diffraction.

It is, therefore, possible to grow near-bamboo structures with large grains, which are desirablestructures for better electromigration resistance of the silver metallization.

The plan-view TEM micrographs of as-deposited Ag/Ta bilayers are compared with thatannealed at 600 8C for 15 min (Fig. 33) [48]. Compared to Ag/Ti, which showed a mono-modaldistribution, a bimodal distribution is observed for Ag/Ta. The bimodal grain size distribution for Agon Ta is associated with abnormal grain growth, which is shown in Fig. 33b. The selected areadiffraction results show that the abnormal grown grains are primarily h1 1 0i and h1 1 2i oriented,with a mean grain size four to five times larger than that of normal grains. For Ag films in the Ag/Tasystem, vacuum annealing only induces slight Ag h1 1 1i texture from the original randomorientation in the as-deposited films.

Fig. 31. Schematic {1 1 1}-pole figure for the 600 8C, 120 min annealed Ag films. The intensity values are expressed inmultiples of the average intensity [48].

Fig. 32. Plan view micrographs of Ag/Ti bilayer samples: (a) as-deposited and (b) annealed at 600 8C for 15 min [48].


Zeng et al. proposed a simple model of the grain growth and texture evolution in Fig. 34a–d[48]. Based on the texture measurement results, the as-deposited Ag film consists mainly of twotypes of grains, i.e. h1 1 1i and randomly oriented grains. A cross-sectional view for the as-depositedAg film is schematically illustrated in Fig. 34a, showing the configuration of two types of grainslabeled by the letters A–D. The estimated mean grain size (�85 nm) is close to the film thickness

Fig. 33. TEM plane-view micrographs of Ag films in Ag/Ta structures: (a) as-deposited and (b) annealed at 600 8C for15 min [48].

Fig. 34. A proposed model to illustrate the grain growth and texture evolution of Ag films with a strong pre-existingh1 1 1i texture upon annealing: (a) schematic cross-sectional view of an as-deposited film on the substrate, showing theconfiguration of two (1 1 1) oriented grains (B and C) and two randomly oriented grains (A and D); (b) at the early stage ofannealing, two new grains (B0 and C0) form as a result of the growth of the initial (1 1 1) oriented grains at the expense ofthe random grains. Two possible consequences for further annealing: (c) grain B0 grows, while grain C0 shrinks; (d) grain C0

grows, while grain B0 shrinks [48].


(�100 nm), hence a columnar grain structure is assumed. Due to the fact that not all h1 1 1i-texturedgrains are strictly h1 1 1i-oriented, the (1 1 1) plane of grain C in Fig. 34a is not parallel to the filmsurface but shows a slight deviation. At the early stage of annealing, microstructure evolution occurswith some grains growing at the expense of the adjacent ones. Grains B and C have similarcrystallographic orientation, therefore, the misorientation between them is small. This results in alow mobility of the grain boundary between them [52]. In contrast, the grain boundary betweengrains A and B, or grains C and D has higher mobility due to larger misorientation than the grainboundary between grains B and C. The Ag(1 1 1) plane has the lowest surface energy; hence thegrowth of h1 1 1i grains is favored to minimize the surface energy. Consequently, the random grainsA and D are partially or completely consumed by the two adjacent h1 1 1i-oriented grains B and C,respectively (Fig. 34b). It has been shown earlier that the strain energy does not play a dominant rolein the determination of which grain would grow, so the possible transformation of h1 1 1i-orientedgrains into h2 0 0i-oriented grains has been ignored. The consequence of the initial grain growth isthe enhancement of the h1 1 1i texture.

The case for the grain growth during further annealing is somewhat different. Apparently,there is no significant surface energy advantage for either of the grown grains B0 and C0, since theyare both h1 1 1i-oriented. Whether grain B0 or grain C0 expands at the expense of the other, dependsprimarily on their grain sizes or curvatures. Thus there are two possibilities: grain B0 grows whilegrain C0 shrinks (Fig. 34c); or vice versa (Fig. 34d). During the growth there is no change ofcrystallographic orientation for a grain, therefore, different h1 1 1i texture intensities will resultcorresponding to these two consequences. For the grain in Fig. 34c, the measured h1 1 1i intensityis obviously higher than that for the case in Fig. 34b, while the h1 1 1i intensity of the grain inFig. 34d is lower than that for the case in Fig. 34b. Since these two outcomes have identicalpossibility to occur, the overall h1 1 1i intensity measured by X-ray diffraction over a large samplearea should remain unchanged. The presented model correlates well with the electron diffractionresults.

3.4. Residual stress in encapsulated silver films

To examine the texture and stress evolution in Ag patterns after the encapsulation process, Zenget al. [48] used XRD rocking curves and sin2 c analyses of both unencapsulated and encapsulatedAg patterns as shown in Fig. 35. Similar to the rocking curves of blanket Ag films before and afterencapsulation process, strong h1 1 1i texture evolution is also observed in annealed Ag patterns. Thestress state in self-encapsulated Ag (200 nm)/Ti (50 nm) structures was investigated by using the‘‘sin2 c’’ technique [53]. In this technique, an elastic lattice strain ehkl corresponding to the specific{h k l} plane in the thin-film material can be experimentally determined by measuring the change ininterplanar spacing of this plane:

ehkl ¼d d0

d0

(5)

where d0 is the interplanar spacing in the unstressed material and d the interplanar spacing in thematerial under stress. Based on certain assumptions and by using the ‘‘c-splitting’’ effect, therelation between the strain ehkl and strain tensor coefficients with respect to the film coordinatesystem can be reduced to a simple linear relationship [25,54–56]:

ehkl ¼ ðexx ezzÞ sin2 cþ ezz (6)


or

d ¼ d0ðexx ezzÞ sin2 cþ d0ðezz þ 1Þ (7)

If the d values of various crystallographic planes corresponding to different tilting angles c aremeasured by using Bragg’s law, a d versus sin2 c plot can be obtained. The values of exx and ezz canbe calculated from the slope and intercept of this plot if the value of d0 is known. Since a dominanth1 1 1i fiber texture was observed in all Ag films of Ag/Ti structures, and it is also well known thatthe elastic constants E and n (Young’s modulus and Poison’s ratio, respectively) are invariants for alldirections within {1 1 1} planes of cubic crystals, the Ag film is isotropic in the plane parallel to thesurface [25]. As a result, the stress tensor coefficients can be calculate according to the elasticstiffness coefficients in the film coordinate system by using the following simplified equations [54]:

sxx ¼ syy ¼ ðc110 þ c12

0Þexx þ c130ezz (8)

szz ¼ 2c130exx þ c33

0ezz (9)

The stiffness coefficients c110, c12

0, c130 and c33

0 are obtained by a transformation from the laboratorycoordinate system to the film coordinate system [56].

For pure Ag, the corresponding stiffness tensor coefficients are given in Table 1 [25,57]. In thisinvestigation, y–2y geometry scans were made using three high-angle reflections from Ag{3 3 1},{4 2 0} and {4 2 2} planes of Ag (200 nm)/Ti (50 nm) samples. Due to strong h1 1 1i texture,{1 1 1} planes of most grains in Ag films are approximately parallel to film surface. The angles cassociated with these high-angles were therefore, practically equal to the inter-angles between{1 1 1} and these planes. These values are listed in Table 2.

Fig. 35. XRD rocking curve of Ag{1 1 1} for unencapsulated and encapsulated Ag patterns, showing the improvement oforiginal strong h1 1 1i texture [35].

Table 1

Calculated stiffness tensor coefficients for laboratory and film coordinate systems

Laboratory coordinate system c11 ¼ 124 GPa c12 ¼ 93.4 GPa c44 ¼ 46.1 GPaFilm coordinate system c11

0 ¼ 154:8 GPa c120 ¼ 83:1 GPa c13

0 ¼ 72:9 GPa c330 ¼ 165:1 GPa


The plots of d200 versus sin2 c for the as-deposited and annealed Ag/Ti samples are shown inFig. 36 [48]. It can be seen that, for each sample, the measured d200 values follow a linearrelationship with sin2 c, which is consistent with Eq. (7). Two straight lines corresponding to 15 and30 min annealed samples coincide completely. They possess a larger slope than the as-depositedsample and a slightly smaller slope than the 60 min annealed sample.

Zeng et al. [48] used Eq. (7) to determine the strain tensor components (exx, eyy and ezz) of theunstressed value of the {2 0 0} interplanar spacing d0. Since there is generally no out-of-plane stressfor a uniform thin film with or without encapsulated layer, one can combine Eq. (6) (assumingehkl ¼ 0) and Eq. (9) (assuming szz ¼ 0) to obtain the condition for measuring d0:

sin2 c ¼ 2c130

2c130 þ c33

0 ¼ 0:469 (10)

The d0 values can be measured from the d200 versus sin2 c linear plots (Fig. 36) by satisfying theabove sin2 c value. For the as-deposited and annealed samples, the values of d0 are 0.20413 and0.20416 nm, respectively. The results are given in Table 3 [25].

The small difference between them might be due to the residual Ti in Ag films for the annealedsamples or just due to the experimental error. Based on Fig. 36 and the measured d0, the strain andstress were calculated using Eqs. (7) and (9), respectively. The study showed that the stress and straincomponents in the encapsulated Ag films are much higher than those of the as-deposited film.

Table 2

Experimental parameters used for X-ray diffraction measurements of the strain and stress in Ag films [25]

Reflection planes 2y0 (8) Tilt angles c (8)

{3 3 1} 110.5 22.0, 48.5{4 2 0} 114.9 39.2, 75.0{4 2 2} 134.9 19.5, 61.9

The measured inter-planar spacing for these planes was normalized relative to the {2 0 0} inter-planar spacing, in order to

obtain a plot of d vs. sin2 c.

Fig. 36. The interplanar spacing of Ag{2 0 0} planes, d200, is plotted vs. sin2 c for the as-deposited and three annealed Ag/Ti samples at 600 8C [48].


According to the sign of the strains, the Ag lattice is in tension within the film plane and incompression along the normal direction.

Zeng et al. [48] showed that a low tensile stress of approximately 61.0 MPa primarily arisingfrom the non-equilibrium growth during the film deposition was present in the as-deposited Ag films[58]. Evaporated films are initially characterized by high disorder and entropy, and then they evolvetowards a thermodynamic equilibrium at later deposition stage. Frequently, this can give rise to adecrease in volume or an increase in atomic density, thereby generating a tensile stress [58]. Thisstress resulted in a lattice tension state in the film plane and a lattice compression state along the filmnormal [25].

The encapsulation process generated thermal mismatch stress at 600 8C. Due to the differencebetween the coefficients of thermal expansion of Ag film and silicon substrate, a thermal mismatchstress is produced when the sample is cooled from elevated temperature. The amount of thermalmismatch stress is generally given by

sth ¼ MðaAg asubÞDT (11)

where M is the biaxial elastic modulus of the thin film, aAg and asub are the linear coefficients ofthermal expansion for silver and silicon substrate (aAg ¼ 19:0 � 10 6/8C and asub ¼ 3:0 � 10 6/8C),and DT is the temperature difference (DT ¼ 580 8C) [25]. It should be noted that, in the estimation ofthe thermal stress, the effect of the underlayer of Ti was ignored, because it is very thin.

Due to strong h1 1 1i texture, the Ag film is isotropic in the film plane. The value of M can thusbe obtained by [59]

M ¼ 6c44ðc11 þ 2c12Þc11 þ 2c12 þ 4c44

¼ 173:6 GPa (12)

With the use of Eqs. (11) and (12), a thermal stress of approximately 1.61 GPa is estimated. Thestrong h1 1 1i texture increases the thermal stress by changing the elastic modulus in the film plane.Most of the stress relaxed during the cooling stage and a residual tensile stress of �320 MPa in thefilm plane was determined.

4. Dry etch techniques

Zeng et al. [24] investigated oxygen (O2)-based dry etching of Ag. Nguyen et al. investigatedtwo etch chemistries, namely, oxygen-based (O2 and Cl2/O2) [60] and fluorine-based (CF4 andCF4/O2) [61] glow discharges to pattern silver metallization.

In Fig. 37a–d [24], the reactive ion etch mechanism of Ag in an oxygen plasma is schematicallydepicted. Under the bombardment of reactive species in the oxygen plasma, the Ag uncovered by the

Table 3

Strain and stress measurement results for Ag films in Ag/Ti bilayers

Samples Strain (%) Stress (MPa)

exx eyy ezz sxx syy szz

As-deposited þ0.035 þ0.035 0.031 þ61.0 þ61.0 0.0600 8C, 15 min þ0.176 þ0.176 0.155 þ306.0 þ306.0 0.0600 8C, 30 min þ0.176 þ0.176 0.155 þ306.0 þ306.0 0.0600 8C, 60 min þ0.192 þ0.192 0.169 þ334 þ334.0 0.0


resist oxidizes and this causes thickening of the Ag film. The Ag oxide layer then flakes off layer-by-layer or part-by-part (Fig. 37a) due to the strain-induced cracking and ion bombardment. The strainis caused by the significant volume expansion of the oxide layer relative to the original Ag film. Thecorresponding plane-view micrograph of the Ag samples undergoing dry etch in a 50 W-powered O2

plasma for 4 min is shown in Fig. 37b. From energy dispersive spectroscopy (EDS) mapping (notshown), the white pattern area is silver, while the surrounding dark area is Ti/SiO2. The arrows showthat the white flake-like substance on the SEM micrograph refers to the Ag oxides formed over theuncovered areas. All Ag uncovered by the resist eventually oxidizes and flakes off completely(Fig. 37c). The Ag dry etch is thus completed after 5 min. This is confirmed by the SEM micrographshown in Fig. 37d, which indicates that all flake-like material is gone and only the Ag patternprotected by the photoresist remains.

The Ag etch process using O2 or Cl2/O2 glow discharges produce a significant volume changeand generate stresses in the Ag layers [24,62–65]. The exfoliation of these films is a directconsequence of the excessive stress levels. Ag etching by O2 glow discharge also produced roughfilm topography, low removal rates, and high resist erosion rates [60]. O2 glow discharge is,therefore, not suitable for the patterning of narrow silver lines and small size features. In Cl2/O2 glowdischarges, the etch rate is five to nine times higher than in O2 only, however, the corrosion andagglomeration of Ag in chlorine ambients make this process also undesirable [60,66,67].

In view of the aforementioned, it was decided to investigate the feasibility of CF4 and CF4/O2

glow discharges for the patterning of Ag metallization. The rest of this section will therefore bedevoted to the presentation of Nguyen et al. [61] etch results and mechanism of CF4 and CF4/O2

chemistries [61].

Fig. 37. Schematic diagram showing the progression of Ag dry etch and the corresponding plane-view SEM micrographs(with 10 mm width for the horizontal test line) of Ag samples subjected to a 50 W-powered O2 plasma: (a) and (b)show exfoliation of Ag oxide layer by layer or part by part (100 mTorr, 4 min); and (c) and (d) show complete exfoliationof Ag oxide and completion of Ag dry etch (100 mTorr, 5 min). The photoresist had been removed prior to the SEMobservations [24].


For the CF4 and CF4/O2 glow discharge etching experiments, silver films of about 200 nm weredeposited on SiO2 via electron-beam evaporation. Optical lithography by using a positive photoresist(Olin OCG 825), and a Karl Suss aligner was used to transfer the pattern from the chrome mask tothe wafer. The samples were etched in a plasma technology, single wafer; parallel plate reactor.Design of experiments was employed to determine the dependence of the removal rates androughness on the process operating conditions. The reaction-product-film removal occurred in awarm (80 8C) Microstrip-3001 solvent manufactured by Olin Chemicals. A Dektak Microscan (step-height scanning profiler) was used to measure the etch depth. SEM, RBS and X-ray photoelectronspectroscopy (XPS) were employed to analyze post-etch surfaces, films and surface chemistries.Further experimental details can be found in [60].

The SEM micrographs, Fig. 38a–d, show the removal-rate dependence on pressure and powerfor high CF4/O2 ratios [61]. The removal rate dependence on the power is compared in Fig. 38a–c.For the same pressure of 50 mTorr, the sample exposed to a power of 150 W (Fig. 38c), shows afaster removal rate than that of the sample exposed to 50 W (Fig. 38a). A similar behavior of removalrate with power is shown in Fig. 38b and d for a pressure of 150 mTorr. By comparing Fig. 38a and b,or Fig. 38c and d the dependence of removal rate on pressure can be observed. High pressures lead tofaster etch rates and more uniform surfaces. The strong effect of power and pressure on the removalrate and surface roughness indicates that the reactive neutrals and ions are important in the gas-substrate reactions. The removal rate is proportional to the applied power and the increase in powerresulted in higher plasma heating. Heating is most likely the cause for faster removal rate, and henceinfluences surface reaction. The etch rate also depends on the pressure; at higher pressures, higherspecies fluxes arrive at the surface; therefore, the surface reaction rate is increased which in turnleads to higher removal rates. This suggests that the removal rate depends on the species fluxesarriving at the surface and hence also influences the fluorination rate.

Fig. 39 shows a SEM micrograph of the sample etched in a CF4 discharge [61]. The resultsshowed that the samples etched in pure CF4 displays the highest etch rate and the cleanest post-cleansurface. The sample etched at high power (150 W) and pressure (150 mTorr) in CF4/O2 (Fig. 38d)

Fig. 38. Qualitative comparison of removal rates and uniformity dependence on pressure and power for high CF4/O2 (25/5 sccm) ratios. These samples are exposed to a plasma for 80 s. The SEM micrographs show samples exposed to (a)150 mTorr, 150 W; (b) 150 mTorr, 50 W; (c) 50 mTorr, 150 W; and (d) 50 mTorr, 50 W [61].


also shows a very high etch rate and clean surface. However, the CF4/O2 (at 150 W, 150 mTorr) isnot favorable, because the photoresist etch rate is also high.

Blanket samples were etched in pure CF4 or CF4/O2 mixture (with a ratio of 1:1) at 100 W and100 mTorr to study the effect of different surface chemistries. XPS was used to analyze the as-etchedsamples, and the post-clean samples (i.e. after a solvent clean in warm Microstrip) [13]. For samplesetched in CF4 discharge, the XPS results of the atomic concentrations of the species present areshown in Fig. 40 [61]. The data indicates that the constituents of the as-etched surface are carbon andoxygen. Silicon, silver and fluorine represented trace amounts. The results further suggest that aportion of the Ag has been removed because a silicon signal is detected, which is believed to comefrom the underlying SiO2 layer. It is evident that a post-clean in warm Microstrip solvent (80 8C)

Fig. 39. SEM micrograph depicts the clean and uniform pattern of a post-clean sample previously exposed to a CF4 glowdischarge. The conditions of the plasma-exposed process are 30 sccm CF4, 100 mTorr and 100 W, for 80 s [61].

Fig. 40. Histogram shows the surface concentrations of as-etch and post-clean samples exposed to CF4 discharge (30 sccm,100 mTorr, 100 W, 45 s). The post-clean surface concentrations of Ag and F drops to an insignificant amount [61].


further enhances the removal of the reaction products. This is noticeable by the drop in the atomicconcentration of Ag on the surface to an insignificant amount and the complete removal of fluorine.The carbon concentration also decreases during this process, while both the oxygen and siliconconcentration increases.

To further demonstrate the effect of power on the removal rate, RBS was employed to analyzepost-cleaned blanket samples, that is, after exposure to a CF4 plasma of power of 50 and 100 W,respectively. Fig. 41 shows the RBS spectra of two etched samples and the as-deposited sample [61].For illustration purposes, the RBS signals of the samples exposed to plasmas of 50 and 100 W aremultiplied by a factor of 5 and 50, respectively. The data of the spectra in Fig. 41, reveals that the Agis completely removed in the sample etched at 100 W; whereas, portions of the Ag are still presenton the surface of the sample etched at 50 W. The lower concentration indicates that the surface Ag isnot pure and the sloping back edge is due to surface roughness [68].

A simple schematic of the removal mechanism of Ag during patterning with CF4 or CF4/O2

glow discharges is shown in Fig. 42 [61]. The key to this process is the high solubility of silverfluorides and oxyfluorides in warm water or solutions containing ammonia [69]. The Microstripsolvent contains hydroxylammonium sulfate, methyl ammonium hydroxide and water. Silverfluorides and its derivatives are soluble in warm water and ammonium-based solutions [69].Therefore, the resist strip in the Microstrip solvent enhances the removal of the reacted materials.

Based on the results of Nguyen et al. [61], the effect of the CF4 discharge, power and pressureon the etch rate is summarized as follows.

4.1. The CF4 glow discharge

The analyses confirmed that Ag etch in CF4 is superior over CF4/O2 discharges despite theexpected increase in fluorine radical concentration when oxygen is added to the discharge. UsingCF4 glow discharge gives a clean and uniform Ag etched pattern.

Fig. 41. RBS spectra of post-cleaned samples etched at different powers in a CF4 discharge (30 sccm, 100 mTorr and 50 s).The spectrum of the sample exposed to 50 W shows some Ag still present on the surface; whereas, there is no Ag presenton the sample exposed to 100 W. The signals of the etched samples at 50 and 100 W are multiplied by 5 and 50 times,respectively [61].


4.2. The effect of power on etch rate

The results showed a strong dependence of the etch rate on power (at constant pressure) but isindependent of flow rates of the gases. At low power, the etch rate is low. However, at a power ofabout 100 W, the etch rate rapidly increases. The increased etch rate is due to the higher plasmaheating induced by the higher power. In turn, the reaction rates strongly depend on temperatureincrease. It has been shown that plasma density increases as the power increases [70,71]. The surfacereaction rate is proportional to the plasma density. The removal rate appears to scale almost linearlyas a function of power in the range 100–175 W [61].

4.3. The effect of pressure of etch rate

Evaluation of the effect of pressure on etch rate revealed a strong chemical dependence. As thepressure is increased to about 150 mTorr, the neutral species flux arrival increases and higherremoval rate are observed. At pressures above 150 mTorr, more collisions occur in the bulk plasmaand enhance recombination of reactive radicals at the Ag surface, which results in a decreased

Fig. 42. Schematic of the removal mechanism of Ag patterned in CF4 or CF4/O2 glow discharges [61].


removal rate. Based on these findings, it follows that the removal mechanism of Ag in CF4 dischargedepends on the reactive neutral species.

5. Integration of low-K dielectrics

5.1. Thermal stability of silver on parylene-n

Gadre and Alford have investigated the integration of the low K-dielectric parylene-n (Pa-n)with Ag films of thickness �1 mm [72]. Prior to processing, the films were cleaned with acetone,de-ionized water and dried in dry nitrogen gas. Silver films of about 200 nm were deposited onto thePa-n via electron-beam evaporation and subsequently annealed in vacuum in the temperature range100–375 8C for 30 min.

XRD, RBS and SIMS were employed to characterize and analyze the Ag/Pa-n and Pa-n/Sisystems, in terms of composition, thin-film interdiffusion and texturing of the silver films [72]. Thefour-point-probe technique was used to determine the sheet resistance of both the as-deposited andannealed samples. In situ resistivity measurements were performed on the Ag/Pa-n samples duringannealing in vacuum. Adhesion analysis of Ag on Pa-n and Si on Pa-n was done by using scratch[73] and tape tests. To improve the adhesion between Ag and Pa-n, the surface of the parylene-n wastreated in oxygen plasma (50 W for 60 s), prior to deposition. The pre-treated Pa-n surfaces werethen analyzed by atomic force microscopy (AFM), RBS and the four-point-probe technique. Furtherexperimental details are given elsewhere [72].

Fig. 43 compares the RBS spectra of the as-deposited Ag (180 nm)/Pa-n (0.8 mm) structure withthat annealed at 375 8C [72]. Energy of 3.7 MeV (resonance energy for N [74]) was used for the RBS

Fig. 43. RBS spectra of as-deposited and 375 8C annealed Ag/Pa-n films. Both spectra show no diffusion of Ag whendeposited on Pa-n [72].


analysis with the intention to enhance the carbon and oxygen signals and to distinguish the carbonfrom parylene. Previous literature reports showed that silver diffuses into Pa-n at temperatures above350 8C. Within the detection limits of RBS no Ag diffusion into Pa-n was detected for the sampleannealed at 375 8C for 1 h.

Due to the sensitivity and resolution limitations of RBS, SIMS analysis was performed on theAg/Pa-n structure to determine if any diffusion of silver into the dielectric occurred at hightemperatures. Fig. 44 shows a comparison of silver counts normalized to carbon counts for bare Pa-n,as-deposited Ag/Pa-n and Ag/Pa-n annealed at 375 8C, as a function of apparent depth [72]. Theterm apparent depth was used as it was derived from actual depth measured with a profilometer.From the graph it is evident that both the as-deposited and annealed curves almost overlap,indicating no silver in-diffusion into parylene-n for temperatures up to 375 8C. As expected silvercounts in both cases are more than bare Pa-n.

Based on the SIMS results, a diffusion coefficient of 1:47 � 10 14 � 3% cm2/s was calculatedfor the as-deposited as well as annealed samples, which indicates negligible diffusion of silver inPa-n. The insignificant silver diffusion in parylene is attributed to: (1) the closed surface-structure ofparaylene and (2) large atomic size of the silver.

During annealing, Pa-n changes from a-Pa-n, which has a monoclinic crystal structure tob-Pa-n which is trigonal [74]. The XRD results showed that the crystallinity of parylene increasesduring this transformation [72]. Although the crystallinity of parylene increases with annealing itnever reaches 100% crystallization. The parylene surface can therefore be considered consisting oflong crystalline chains linked together with amorphous regions forming a closed structure [75]. Itis believed that this closed structure prohibits the diffusion of silver in parylene. SIMS dataconfirmed that the phase change of Pa-n during annealing does not enhance diffusion of silver inparylene [72].

The diffusion of copper and aluminium in Pa-n has been reported [76]. Compared to Cu (atomicsize 0.135 nm) and Al (atomic size 0.126 nm), the larger atom size of Ag (covalent radius of0.521 nm) will require more space to pass through the Pa-n structure. Annealing the Ag/Pa-n at375 8C is not sufficient to thermally activate diffusion of the silver atoms [72].

Fig. 44. Comparison of SIMS profiles of silver counts for bare Pa-n, as-deposited Ag/Pa-n and 375 8C annealed Ag/Pa-n.The linear portion of the spectrum was used to determine the diffusion coefficient of the annealed silver. The diffusion isbarely above the background level: (^) C as-deposited; (&) Ag/C as-deposited; (~) C 375 8C; (�) Ag/C 375 8C; (�) Cbare; (*) Ag/C bare [72].


5.2. Electrical stability of silver on parylene-n

Ex situ four-point-probe measurements of Ag on Pa-n were performed to obtain the sheetresistance of the Ag films [77]. The resistivity of the Ag films was then calculated from the sheetresistance and the thickness obtained from RBS. The resistivity values are given in Table 4 [72].

The variation of resistivity of Ag with temperature is best explained with XRD analysis.Fig. 45 shows a X-ray diffraction pattern of Ag on Pa-n [72]. The Ag film shows a prominenth1 1 1i peak. The intensity of the h1 1 1i peak increases up to 300 8C and then suddenly decreasesat 350 8C. X-ray diffraction analysis revealed no phase formation between the Ag and Pa-n atelevated temperatures. However, the Pa-n films decomposed when annealed at temperatures above400 8C.

5.3. Adhesion of silver and silicon to parylene-n

Gadre and Alford [72] performed pin-on-disk scratch testing to evaluate the adhesion of Pa-non silicon and the tape test for Ag/Pa-n/Si. AFM was used to determine the surface roughness.Fig. 46 shows optical images of the scratch test results of Pa-n on silicon for as-deposited andannealed samples. Fig. 46a shows that delamination of the as-deposited sample occurs at a load of8 g [72]. It is interesting to note that the annealed samples failed at a significantly lower load than

Table 4

Comparison of the resistivity for the as-deposited sample with that annealed at different temperatures

Sample r (mO cm)

As-deposited 2.73100 8C 2.55250 8C 1.95300 8C 1.87350 8C 3.48

Fig. 45. XRD diffraction patterns (under y/2y scan geometry) of Ag/Pa-n at different anneal temperatures [72].


the as-deposited sample. The sample annealed at the highest temperature (350 8C) required thesmallest load for failure (Fig. 46d).

To evaluate the adhesion of Ag on the Pa-n/Si structure, a simple tape test was performed. Priorto the deposition of the Ag film, the Pa-n surface was treated in O2 plasma. The tape test resultsshowed that the adhesion of Ag on Pa-n improved significantly with annealing temperature whencompared with the untreated Ag/Pa-n/Si structure. AFM imaging of the surfaces of the as-preparedand annealed samples was conducted after the tape test. The AFM images revealed that the surfaceroughness of the O2 plasma treated samples was �80 nm compared to �30 nm for the untreatedsamples (Fig. 46b and c). It is believed that the improved adhesion on Pa-n is due to the increasedsurface roughness cause by the O2 plasma treatment.

The data showed that the adhesion between Pa-n and Si substrate deteriorates with annealing.Pa-n is a chemically inert polymer and does not form any chemical bond with silicon, which is thereason for the poor adhesion between them. Furthermore, the presence of a thin silicon dioxidesurface layer may weaken the bonding of Pa-n with silicon [72]. Oxygen plasma treatment of theparylene surface prior to deposition of silver and subsequent annealing proved to be an effective

Fig. 46. Scratch test analysis is contrasted with AFM results for (a) as-deposited Pa-n on Si; (b) untreated, as-deposited Pa-n on Si; (c) 350 8C, 30 min; and (d) oxygen plasma treated Pa-n [72].


method to improve the adhesion between Ag and Pa-n. The oxygen plasma induced damage to theparylene surface and hence increases its roughness. The rough films are believed to increasemechanical interlocking between the two films, thereby improving the adhesion significantly.

6. Evaluation of electromigration resistance

Electromigration (EM) resistance of Ag is the most critical interconnect reliability concern thatwould determine its suitability for integrated circuit technology. EM is the drift of metal ions as aresult of either collision between the conductor electrons and/or the metal ions or high electrostaticfield force when current is passed through a metal conductor. The direction of mass transportdepends on the direction of the net force. When high current densities pass through the metal line,voids or hillocks are formed at a point of ion flux divergence (Fig. 47). Voids and hillocks deformand grow until electrical failure is completed [78].

Ho and Huntington [79], Patil and Huntington [80] reported that Ag migrates toward the anodefor the temperature range 670–877 8C. Klotsman et al. [81] and Breitling and Hummel [82] however,found that Ag ion transport is toward the cathode, and is opposite to the direction of the electronwind. Studies conducted by Hummel et al. showed that the dominating movement path is grainboundaries for temperatures in the range of 225–280 8C (activation energy of 0.95 eV) and surfacefor the temperature range 160–225 8C (activation energy of 0.3 eV) [83]. This study investigatedmass transport direction, and its dependence on current densities, capping materials and temperatureof accelerated EM test conditions.

Zeng et al. [84] investigated the influence of the encapsulation process on the electromigrationresistance. The sample configurations used consisted of bare Ag patterns (samples I, II, III and IV)and encapsulated Ag patterns (samples V, VI, VII and VII). The average of four failure times wasused to evaluate the electromigration resistance. In general, the test results for the bare Ag lines showthat in the early stage of the test, a partial depletion region first appears at the line segment, which isimmediately near the cathode pad. A partial depletion region also extends progressively towards theanode side. During the drift of this partial depletion region, significant voids and hillocks also formside by side at other areas of the line. A typical example is given in Fig. 48 [84], which shows theSEM morphology of the center area of the bare line sample IV after 3 h testing. A cluster of hillocksis located immediately on the right side of the voiding area. This confirms that mass transport isdirected towards the anode (the right-hand direction) and the formation of hillocks is theconsequence of mass depletion from voiding areas. The electromigration test results of four bare Aglines are summarized in Table 5.

For all tested Ag lines which were encapsulated by TiN(O), no partial depletion area and nosignificant number of voids and hillocks were observed even after testing for much longer time thanthat for bare Ag lines (see Table 6). For instance, the encapsulated line sample VI, which was testedfor 7 h displayed only four to five voids across the entire length of the line. The voiding area of this

Fig. 47. Schematic to illustrate metal conductor failure due to electromigration.


sample is shown in Fig. 49 [84]. These results were typical as well for other test samples. Comparedto bare Ag lines, encapsulated Ag lines exhibited much better electromigration resistance in terms ofthe time required for the formation of significant number of voids and hillocks. Based on the testsconducted on four Ag lines, the average time for the formation of considerable number of voids and

Fig. 48. Plane-view SEM micrograph indicating the surface morphology of the bare Ag test sample IV (near-cathodesegment), tested for 3 h. The line width is approximately 10 mm. The left portion of the line segment is the partial depletionregion, which is composed of a large number of voids and hillocks. It is noted that the actual electric current flowedthrough the lines from the right to the left [84].

Table 5

Electromigration test results of bare Ag lines

Sample line Actual test time (min) Partial depletion region time (min)

I 180 90II 90 25III 90 40IV 180 50

Table 6

Electromigration test results of self-encapsulated Ag lines

Sample line Actual test time (min) Surface morphology changesafter testing (voids)

V 660 11–13VI 420 4–5VII 540 8–9VIII 600 8–9


hillocks (partial depletion region) in bare Ag lines was approximately 50 min (see Table 5).However, the corresponding time for encapsulated Ag lines was at least 9 h.

The cluster of hillock adjacent to the void area (shown in Fig. 48) is consistent with theobservation of Breitling and Hummel [82] that voids and hillocks are uniformly side-by-sidedistributed along the electromigration-tested Ag strip. It is known that the electromigration failuresof interconnect lines are primarily caused by gradients which can lead to atomic flux divergencesunder the electrical driving force. There are two major types of gradients that may exist: localgradients due to microstructural or chemical inhomogeneities, and long-range temperature gradientdue to nonuniform Joule heating and heat dissipation. Based upon our observations, it appears that inthe present case the influences from both gradients are superimposed. The side-by-side distributionof voids and hillocks along the whole line length is attributed to the local inhomogenities, such asdifferent grain sizes, and variation in film thickness. Partial depletion occurring near the cathode isassumed to result from the long-range temperature gradient due to the nonuniform Joule heatingcaused by the difference in the cross-section area of the test line and pads.

The round shape of the hillocks as seen in Fig. 48 [84] implies that surface diffusion is mostlikely to be the dominant mechanism of mass transport, since the agglomeration of hillocks is theresult of minimizing surface energy configuration by the surface diffusion of Ag atoms. Forsimplification, it is assumed that all Ag grains have a hexagonal shape and transverse through theentire film thickness. If this is the case, then the ratio of the surface area to the grain boundary areacan be estimated to be �2.3 for bare Ag lines (based on the estimated average Ag grain size from theTEM micrographs). As a result the larger surface area and lower activation energy for surfacediffusion [83] leads to surface diffusion dominant electromigration. After the encapsulation process,the Ag surface is capped by a thin layer of TiN(O). The mobility of Ag atoms at the TiN(O)/Aginterface is substantially reduced. Therefore the electromigration resistance is improvedsignificantly. By examination of the surface of the encapsulated Ag lines tested for a long time,almost no round hillocks are observable. This implies that the surface diffusion has been hinderedsubstantially. The formation of voids in this case is most likely to arise from the grain-boundary

Fig. 49. Plane-view SEM micrograph of the TiN(O)-encapsulated Ag line sample VII, tested for 7 h. The line width isapproximately 10 mm. Only four to five voids were found across the entire length of the line [84].


diffusion. Comparison of Tables 5 and 6 suggests that the encapsulation process improves theelectromigration resistance of Ag metallization by at least one order of magnitude for the teststructures and conditions used [84].

7. Summary

The paper reviews bilayer and alloy techniques with Ti, Al and others to form adhesion layersand diffusion barriers. The temperature range of thermal stability is covered. During Ti and Altransport to form the encapsulating layers, the Ag films develop h1 1 1i texture. Dry etch techniqueshave been developed using O2 and CF2. Integration with low-K dielectrics such as paralene has beendemonstrated. Encapsulated Ag films have superior resistance to electromigration degradation.

7.1. Future trends

Ag metallization research has intensified significantly in recent years. Research work has beenconducted mainly in the US. Increasing numbers of companies have already begun research anddevelopment (R&D) efforts in Ag-based metallization. However in this highly competitive era theindustrial research has been propriety. Ag-based interconnects represents the future trend in the deepsub-micron regime. Ag is an attractive material for interconnect-metallization in future integratedcircuits technologies due to its low bulk resistivity and high reliability against electromigration. Agcan be deposited by plating (electroless and electrolytic), sputtering (physical vapor deposition(PVD)), laser induced reflow and chemical vapor deposition (CVD) cluster tools, chemicalmechanical polishing (CMP) polishers and low temperature etchers.

There are still many critical issues that remain to be resolved in the development of Ag-basedmetallization. For instance, to meet the throughput requirement in device manufacturing, there is aneed to demonstrate a Ag deposition technique with a high deposition rate (>250 nm/min) at lowsubstrate temperatures without simultaneously sacrificing low resistivity, good step coverage andcomplete via fill. The adequate use of thin diffusion barriers in Ag-based metallization will becritical for the 0.25 mm technology. Other materials science issues in Ag-based metallization need tobe address: (1) microstructural control (e.g. grain size and texture); (2) contamination control (C, Oand CI); (3) oxidation and corrosion control; (4) prevention of Ag diffusion in metals and dielectricmaterials; and (5) mechanical properties (e.g. stress migration, adhesion improvement). In addition,the development of Ag-based on-chip interconnects technology will also lead to new advancementsin electronic packaging technology. The implementation of Ag-based interconnects will representthe future trend in the deep submicron regime.

Acknowledgements

The work was partially supported by the National Science Foundation (NSF) to whom theauthors are greatly indebted. Support for the Center for Low Power Electronics is partially providedby the NSF, State of Arizona, Analog Devices, Analogy, Burr-Brown, Hughes Aircraft, Intel,Microchip, Motorola, National Semiconductor Rockwell, Scientific Monitoring, Texas Instrumentsand Western Design Center. The authors also would like to express their sincere thanks to theNational Research Foundation (NRF) and the University of the Western Cape in South Africa fortheir financial support.


References

[1] J.M.E. Harper, K.L. Holloway, T.Y. Kwok, US Patent No. 5,130,274 (1992).[2] The National Technology Roadmap for Semiconductors, Semiconductor Industry Association, San Jose, CA, 1994.[3] J. Li, J.W. Mayer, Y. Shacham-Diamand, E.G. Colgan, Appl. Phys. Lett. 60 (1992)2983.[4] P.A. Totta, W.E. Mutter, unpublished work, 1968.[5] T. Iijima, H. Ono, N. Ninomiya, Y. Ushiku, T. Hatanaka, A. Nishiyama, H. Iwai, Extended Abstracts of the 1993

International Conference on Solid State Devices and Materials, Makuhari, 1993, p. 183.[6] S.P. Murarka, R.J. Guttman, A.E. Kaloyeros, W.A. Lanford, Thin Solid Films 236 (1993)257.[7] J.D. McBrayer, R.M. Swanson, T.W. Sigmon, J. Electrochem. Soc. 133 (1986)1243.[8] T.E. Graedel, J. Electrochem. Soc. 139 (7) (1992)1963.[9] B. Chalmers, R. King, R. Shuttleworth, Proc. R. Soc. A 193 (1948)465.

[10] A.E.B. Presland, G.L. Price, D.L. Trimm, Prog. Surf. Sci. 3 (1973)63.[11] S.K. Sharma, J. Spitz, Thin Solid Films 65 (1980)339.[12] S.K. Sharma, J. Spitz, Thin Solid Films 66 (1980)L51.[13] P.N. Nguyen, Ph.D. thesis, Arizona State University, 2000.[14] M.E.C. Willemsen, A.E.T. Kuiper, A.H. Reader, R. Hokke, J.C. Barbour, J. Vac. Sci. Technol. B6 (1988)53.[15] S.Q. Wang, I. Raaijmakers, B.J. Burrow, S. Suthar, K.B. Kim, J. Appl. Phys. 68 (1990)5176.[16] T.L. Alford, D. Adams, T. Laursen, B.M. Ullrich, Appl. Phys. Lett. 68 (1996)3251.[17] Y.L. Zou, T.L. Alford, D. Adams, T. Laursen, K.N. Tu, R. Morton, S.S. Lau, MRS Proc. 427 (1996)355.[18] J. Li, J.W. Mayer, E.G. Colgan, J. Appl. Phys. 70 (1991)2820.[19] D. Adams, T. Laursen, T.L. Alford, J.W. Mayer, Thin Solid Films 308 (1997)450.[20] L.R. Doolittle, Nucl. Instrum. Methods Phys. Res. B9 (1985)344.[21] D. Adams, T.L. Alford, N.D. Theodore, S.W. Russell, R.L. Spreitzer, J.W. Mayer, Thin Solid Films 262 (1995)199.[22] Y.L. Zou, T.L. Alford, Y. Zeng, F. Deng, S.S. Lau, T. Laursen, A. Amali, B.M. Ullrich, J. Appl. Phys. 82 (7)

(1997)3321.[23] C.Y. Ting, M. Wittmer, S.S. Lyer, S.B. Brodsky, J. Electrochem. Soc. 131 (1984)2934.[24] Y. Zeng, L. Chen, Y.L. Zou, P.A. Nguyen, J.D. Hensen, T.L. Alford, Mater. Chem. Phys. 22 (2000)77–82.[25] Y. Zeng, Ph.D. dissertation, Arizona State University, 1999.[26] O. Madelung (Ed.), Landolt-Bornstein, New Series 111–126, Springer, Berlin, 1990, p. 136.[27] T. Okamoto, M. Shimizu, A. Ohsaki, Y. Mashiko, K. Tsukamoto, T. Matsukawa, S. Nagao, J. Appl. Phys. 62 (11)

(1987)4465.[28] S.W. Russell, J.W. Strane, J.W. Mayer, S.Q. Wang, J. Appl. Phys. 76 (1) (1994)257.[29] J.A. Leavitt, L.C. McIntyre, M.R. Weller, in: J.R. Tesmer, M. Nastasi, J.C. Barbour, C.J. Maggiore, J.W. Mayer

(Eds.), Handbook of Modern Ion Beam Materials Analysis, Materials Research Society, Pittsburgh, 1995, p. 40.[30] M. Ohring, The Materials Science of Thin Films, Academic, New York, 1992 (Chapter 8).[31] B.E. Deal, A.S. Grove, J. Appl. Phys. 36 (1965)3770.[32] A.S. Grove, Physics and Technology of Semiconductor Devices, Wiley, New York, 1967 (Chapter 2).[33] N.A. Gjostein, Diffusion, American Society for Metals, Metals Park, OH, 1973.[34] Y. Wang, T.L. Alford, Appl. Phys. Lett. 74 (1) (1999)32.[35] J.A. Dean (Ed.), Lange’s Handbook of Chemistry, vol. 14, 11th ed., McGraw-Hill, New York, 1992, pp. 6–69.[36] Y. Wang, Ph.D. dissertation, Arizona State University, 2001.[37] Y. Wang, T.L. Alford, J.W. Mayer, J. Appl. Phys. 86 (10) (1999)5410.[38] J. Crank, The Mathematics of Diffusion, second ed., Clarendon, Oxford, 1975, p. 60.[39] N. Marecal, E. Quesnel, Y. Pauleau, Thin Solid Films 241 (1994)34.[40] T.B. Maisalski (Ed.), Binary Alloy Phase Diagrams, American Society for Metals, Metals Park, OH, 1986.[41] C. Kittel, Introduction to Solid State Physics, fifth ed., Wiley, New York, 1976, p. 71.[42] K. Schroder (Ed.), CRC Handbook of Electrical Resistivities of Binary Metallic Alloys, CRC Press, Boca Raton, FL,

1983, p. 25.[43] G.F. Malgas, D. Adams, P. Nguyen, Y. Wang, T.L. Alford, J.W. Mayer, J. Appl. Phys. 90 (11) (2001)5594.[44] D. Adams, B.A. J., J.W. Mayer, T.L. Alford, Thin Solid Films 332 (1998) 237.[45] D. Adams, Ph.D. dissertation, Arizona State University, 1996.[46] Y. Zeng, Y.L. Zou, T.L. Alford, F. Deng, S.S. Lau, T. Laursen, B. Manfred Ullrich, J. Appl. Phys. 81 (1997)7773.[47] Y. Zeng, Y.L. Zou, T.L. Alford, Thin Solid Films 20 (1997)89.[48] Y. Zeng, T.L. Alford, Y.L. Zou, A. Amali, B.M. Ullrich, F. Deng, S.S. Lau, J. Appl. Phys. 83 (2) (1998)780.[49] D.B. Knorr, T.M. Lu, Appl. Phys. Lett. 54 (1989)2210.[50] L.K. Jetter, C.J. McHargue, R.O. Williams, J. Appl. Phys. 27 (1956)368.[51] B.D. Cullity, Elements of X-ray Diffraction, Addison-Wesley, Reading, 1978.[52] A.D. Rollett, D.J. Srolovitz, M.P. Anderson, Acta Metall. 37 (1989)1227.[53] I.C. Noyan, J.B. Cohen, Residual Stress: Measurement by Diffraction and Interpretation, Springer, New York, 1987.


[54] R.P. Vinci, E.M. Zielinski, J.C. Bravman, Thin Solid Films 262 (1992)142.[55] H. Shibata, M. Murota, K. Hashimoto, Jpn. J. Appl. Phys. 32 (1993)4479.[56] M. Tsukada, S. Ohfuji, J. Vac. Sci. Technol. B 11 (1993)326.[57] R. Bacon, C.S. Smith, Acta Metall. 4 (1956)337.[58] F.M. D’Heurle, J.M.E. Harper, Thin Solid Films 171 (1989)81.[59] W.D. Nix, Met. Trans. A 20 (1989)2217.[60] P. Nguyen, Y. Zeng, T.L. Alford, J. Vac. Sci. Technol. B 17 (1999)2204.[61] P. Nguyen, Y. Zeng, T.L. Alford, J. Vac. Sci. Technol. B 19 (1) (2001)161.[62] Y.L. Zou, Ph.D. dissertation, Arizona State University, 1998.[63] A. de Rooij, Eur. Space Agency 13 (1989)363.[64] M.K. Bhan, P.K. Nag, G.P. Miller, J.C. Gregory, J. Vac. Sci. Technol. A 12 (1994)699.[65] C.H. Chou, J. Philips, J. Vac. Sci. Technol. A (1991) 2727.[66] C.O. Muller, Corrosion (Houston) 47 (1991)146.[67] D.W. Rice, P. Peterson, E.B. Rigby, P.B.P. Phipps, R.J. Capell, R. Tremoureux, J. Electrochem. Soc. 128 (1981)275.[68] L.C. Feldman, J.W. Mayer, Fundamentals of Surface and Thin Film Analysis, North-Holland, New York, 1986.[69] D.R. Lide (Ed.), CRC Handbook of Chemistry and Physics, 72nd ed., Chemical Rubber, Cleveland, OH, 1991.[70] D.M. Manos, D.L. Flamm, Plasma Etching: An Introduction, Academic, Cambridge, 1989.[71] M.A. Lieberman, A.J. Lichtenberg, Principles of Plasma Discharges and Materials Processing, Wiley, New York,

1994.[72] K.S. Gadre, T.L. Alford, J. Vac. Sci. Technol. B 18 (6) (2000)2814.[73] S.W. Russell, S.A. Rafalski, R.L. Spreitzer, J. Li, M. Moinpour, F. Moghadam, T.L. Alford, Thin Solid Films 262

(1995)154.[74] W.D. Niegisch, J. Appl. Phys. 37 (1966)4041.[75] G.R. Moore, D.E. Kline, Properties and Processing of Polymers for Engineers, Prentice-Hall, New York, 1984, p. 9.[76] G.R. Yang, S. Dabral, L. You, J.F. McDonald, T.M. Lu, H. Bakhru, J. Electron. Mater. 20 (1991)571.[77] K.S. Gadre, Ph.D. dissertation, Arizona State University, 2002.[78] M. Mahadevan, R.M. Bradley, J.M. Bebierre, Europhys. Lett. 45 (1999)680.[79] P.S. Ho, H.B. Huntington, J. Phys. Chem. Solids 307 (1966)1319.[80] H.R. Patil, H.B. Huntington, J. Phys. Chem. Solids 31 (1970)463.[81] S.M. Klotsman, A.N. Timofeyev, I.S. Trakhtenberg, Phys. Metal Metallogr. 14 (1962)140.[82] H.M. Breitling, R.E. Hummel, J. Phys. Chem. Solids 33 (1972)845.[83] R.E. Hummel, H.J. Geier, Thin Solid Films 25 (1975)335.[84] Y. Zeng, L.H. Chen, Y.L. Zou, P.A. Nguyen, J.D. Hensen, T.L. Alford, Mater. Lett. 45 (2000)157–161.


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