Influence of reactive sputter deposition conditions on crystallizationof zirconium oxide thin films

Guneet SethiMaterials Research Institute, The Pennsylvania State University, University Park, Pennsylvania 16802

Paul SunalMEMS and Nanotechnology Exchange, 1895 Preston White Drive, Suite 100, Reston, Virginia 20191

Mark W. Horn and Michael T. Lanagana�

Materials Research Institute, The Pennsylvania State University, University Park, Pennsylvania 16802

�Received 30 July 2008; accepted 20 March 2009; published 1 May 2009�

Zirconium oxide thin films were prepared through reactive magnetron sputtering with a zirconiumtarget using pulsed-dc and radio frequency �rf� sources. The film crystallization was studied withrespect to sputtering growth variables such as sputtering power, sputtering pressure, sourcefrequency, oxygen pressure, substrate temperature, and substrate material. The crystallization wasstudied through x-ray diffraction �XRD� 2� scans and was quantified with peak full width at halfmaximum and crystallite size. Crystallization of the films was found to occur over a broad range ofsputter deposition parameters, while the amorphous phase was produced only at high sputteringpressure and low sputtering power. With a decrease in sputtering pressure or power, the crystallitesize decreased. Energy dispersive x-ray spectroscopy, electron microscopy, and XRD analysisrevealed that at very low pressures, these films are polyphase assemblages of cubic phases ofoxygen deficient zirconium oxides such as ZrO and Zr2O. When the sputtering oxygen content ofthese films is increased above 25%, monoclinic-ZrO2 phase is stabilized in the films and thedeposition rate decreases. However, in the case of rf sputtering, an additional peak corresponding totetragonal phase of ZrO2 is observed. The sputtering parameters were related to physical parameterssuch as sputtering mode, ion energy, and substrate temperature, which influence crystallinity.

© 2009 American Vacuum Society. �DOI: 10.1116/1.3119669�

I. INTRODUCTION

There is an immediate need for fast-response capacitorswith decreased volume, weight, and cost for high-voltagepulsed-power applications and power distribution system.1

Energy density dominates this research challenge and is re-lated to the dielectric constant and breakdown strength of thecapacitor dielectric. Single-cation metal-oxide thin filmssuch as zirconium oxide films are strong candidates for thecapacitor dielectric in high energy density capacitor.2 Thesefilms are commonly deposited through sputtering. Over thepast decade, different researchers prepared zirconia thin filmsthrough magnetron sputtering. Kuo and Chien,3 Amor et al.,4

and Ning-Lin et al.5 used a ceramic target to produce amor-phous ZrO2 films as a gate dielectric in silicon complemen-tary metal-oxide-semiconductor applications. There havebeen other investigations using reactive sputtering of a zir-conium target6–8 to prepare polycrystalline ZrO2 films foroptical coatings. The breakdown strength of these oxidefilms is dependent on several factors including crystal struc-ture and microstructure of the capacitor dielectric.9–14 It isknown that the crystal structure, microstructure, and dielec-tric properties of the films are strongly influenced by thedeposition conditions. In the case of single crystal materials,the crystal phase with highest cation coordination has the

a�Author to whom correspondence should be addressed; electronic mail:

mxl46@psu.edu

577 J. Vac. Sci. Technol. A 27„3…, May/Jun 2009 0734-2101/2009

highest breakdown strength.2 In the case of polycrystallinematerials, materials with smaller grain size have higherbreakdown strength.15 Additionally, amorphous materialshave higher breakdown strength than their crystallinecounterparts.16 This study aims at understanding the effectsof the reactive sputtering conditions on the crystallization ofzirconium oxide thin films. This includes the formation ofspecific crystal structures such as monoclinic, tetragonal, orcubic and crystallite size. The reactive sputtering parametersexamined in this article include sputtering power, powersource type, sputtering pressure, oxygen partial pressure,substrate type, and substrate temperature. In order to obtaincommercially viable amorphous ZrO2 film high energy den-sity capacitors, an understanding of crystallization and filmformation is essential.

II. EXPERIMENTAL DETAILS

Zirconium oxide thin films were prepared using reactivemagnetron sputtering and deposited on �100� silicon andglass substrates in a stainless steel vacuum chamber. The75 mm diameter zirconium target used for sputtering is99.5% pure. Before deposition, the vacuum chamber wasevacuated to a pressure of 1�10−4 Pa using a turbomolecu-lar pump. The target was cleaned by presputtering it for5 min at source power of 50 W with 99.99% pure Ar at0.7 Pa. The substrate was shielded during cleaning of the

target. At the start of the deposition, 99.99% pure O2 was

577/27„3…/577/7/$25.00 ©2009 American Vacuum Society

578 Sethi et al.: Influence of reactive sputter deposition conditions 578

introduced into the chamber, controlled by a 20 SCCM�SCCM denotes cubic centimeters per minute at STP� stan-dard mass flow controller. The O2 fraction is given by

�pO2=

Ptot − PAr

Ptot� , �1�

where pAr is the Ar base pressure and ptot is the measuredtotal pressure. The chamber pressure was measured using afactory-calibrated MKS Baratron capacitance manometer.The manometer was controlled through an MKS 270B highaccuracy signal conditioner with an accuracy of 0.014 Pa.The Ar flow rate was fixed at 42 SCCM for total sputteringpressures greater than 0.7 Pa. It was reduced to produce totalsputtering pressures lesser than 0.7 Pa. Two types of powersources, pulsed direct current �PD� at 150 kHz and radiofrequency �rf� at 13.56 MHz, were used. The target voltagewas measured using an oscilloscope. The substrates wereplaced on a base copper plate and the substrate temperaturewas monitored during deposition using a shielded chromel-alumel thermocouple placed on the copper. The temperaturerecorded by the thermocouple is a representative of the sub-strate temperature because of high thermal conductivity ofthe copper plate and its close contact with the substrate. Thesubstrate to target distance was fixed at 89 mm. Table Ishows the growth parameters �with their specific values� thatwere varied in order to study their effect on crystallization.The deposition rate of the zirconia thin films on the siliconsubstrate was calculated by dividing the deposited thicknesswith the deposition time.

The crystal structure of the films on glass substrates wascharacterized with an x-ray diffractometer �Scintag XDS2000� using Cu K� radiation in grazing incidence. The scanrate, step size, and acquisition time were 0.4 deg /min, 0.02°,and 1.5 s, respectively. Film crystallization was quantifiedthrough full width at half maximum �FWHM� of x-ray peaksand crystallite size �X� analysis using the Scherrer equation�Eq. �2�� by neglecting the microstrain17

X =k�

B cos���, �2�

where B is FWHM after instrument correction, k is a shapefactor �=0.85�, � is the x-ray wavelength, and � is the Braggdiffraction angle. The instrument correction was performedby using an instrument standard NIST Si 640. Most of theprepared films in this investigation have monoclinic crystal

18

TABLE I. Growth parameters studied for their effects on crystallization inthis investigation.

Growth variable Values

Pressure 0.2–13.3 PaPower 100–400 WSource frequency PD �at 60–215 kHz� and rfOxygen fraction 7%–75%Substrate temperature −100 and 25 °C

structure of zirconia. Hence the FWHM and crystallite size

J. Vac. Sci. Technol. A, Vol. 27, No. 3, May/Jun 2009

analysis presented here was conducted on the strongest peak

of the phase �1̄11� at 28.2° corresponding to a d spacing of3.1647 Å. An average of two x-ray diffractions �XRDs� wasrecorded for each sample under a particular deposition con-dition.

The film structure was studied using a JEOL 6700F fieldemission-scanning electron microscope �SEM� at 10 keV ina secondary electron mode. The stoichiometry of the filmswas characterized by using energy dispersive x-ray spectros-copy �EDS� on a Hitachi S-3500N at 15 keV. The EDSequipment was calibrated for oxide analysis using a commer-cially available 100-�m-thick stoichiometric ZrO2 film. Allthe films were characterized for thickness using a TencorAlpha-Step 500 surface profiler.

III. RESULTS

A. Effect of film thickness on the crystallite size

Figure 1 shows the variation in the crystallite size withfilm thickness for zirconia films deposited at 400 W and0.7 Pa with deposition times ranging from 10 to 70 min.There is a slight increase in the crystallite size from105 to 120 Å with thickness ranging from 200 to 500 nm.There is no overall variation in the crystallite size for filmthickness greater than 500 nm. All the films experienced acompressive stress �measured on a Tencor Flexis 2320� in therange of 875 and 985 MPa. This compressive film stress isknown to affect crystallization including the crystal structurein these films.4,19

B. Effect of source type on crystallization

Figure 2 compares the XRD patterns of zirconia filmsprepared with PD and rf sputtering at different powers. Themonoclinic-ZrO2 phase was observed in films prepared byboth PD and rf sputtering. However, films prepared by rfsputtering show the presence of an additional minor phase of

20

FIG. 1. Effect of zirconia film thickness on the crystallite size.

tetragonal ZrO2. In addition, films prepared by rf sputtering

579 Sethi et al.: Influence of reactive sputter deposition conditions 579

appear to have a larger crystallite size at the same power andpressure. This difference is more evident for films made atlower powers.

C. Effect of sputtering power and pressure oncrystallization

Figure 3 shows the effect of sputtering power and pres-sure on the crystallization and deposition rate of the zirconiafilms. As the power increases, a slight increase in the crys-tallite size is observed for both rf and PD sputtering, as seenin Fig. 3�a�. We noted, however, that the FWHM calculationsat 100 W are more error prone due to thinner film thickness,as seen in Fig. 3�c�. In general, the crystallite size increaseswith increasing pressure except for films that are prepared atlower powers and higher pressures �Fig. 3�b��. The rf sput-tered films prepared at 100 W, 6.0 Pa and 150 W, 10.0 Paare very thin �thickness of 50 nm� and show very broadcrystalline peaks with FWHM �5.0, indicating reduced crys-tallite size. The lack of crystallization in these films suggestsa “critical thickness” of 50 nm, below which the films aremainly amorphous. This aspect will be covered in detail inthe discussion. Additionally, we observe that the depositionrate increases with an increase in power. A drastic increase inthe deposition rate was observed for the films deposited be-low 0.7 Pa at 400 W.

Figure 4 shows the thin film and cross-sectional morphol-ogy of two films possessing two extremes of crystallite sizeand deposition rate, as seen in Fig. 3. Figure 4�a� shows afilm deposited at a rate of 18 Å /min and crystallite size85�25 Å fabricated using rf sputtering at 100 W and6.0 Pa. Figure 4�b� shows a film deposited at a rate of96 Å /min and crystallite size of 300�5 Å fabricatedusing PD sputtering at 400 W and 6.0 Pa. Both microstruc-tures consist of 50–75 nm wide columns composed ofgrains. The film morphology follows the structure zonemodel.21 A porous structure consisting of tapered crystallitesseparated by voids at high sputtering pressures and low pow-

FIG. 2. �Color online� Effect of source type on the XRD pattern for sputter-ing at 200 and 400 W at 6.0 Pa. Phases m and t are the monoclinic ZrO2 andtetragonal ZrO2, respectively.

ers is observed in Fig. 4�a�. The film would have to be in the

JVST A - Vacuum, Surfaces, and Films

sputtering zone 2 /3 to achieve a continuous crystal fromsubstrate to film surface. However, a transition structure�zone T� consisting of densely packed fibrous columns at lowsputtering pressures and high powers is observed in Fig.4�b�. The columns here are composed of broken verticallyelongated crystals. Due to the porous structure in the sampleshown in Fig. 4�a�, the rms roughness of these films was high�11 nm for a 50-nm-thick film�. This was measured on aDigital Instruments Dimension 3100 atomic force micro-scope in a tapping mode with a Si cantilever with tip radius

FIG. 3. �Color online� Effect of sputtering �a� power, �b� pressure on crys-tallization, and �c� deposition rate for both PD and rf sputtering with 16%sputtering O2.

of 15–20 nm at a scan rate of 1 Hz. In the following sec-

580 Sethi et al.: Influence of reactive sputter deposition conditions 580

tions, the two films described above are used as baselines toevaluate the effects of different sputtering parameters oncrystallization.

At sputtering pressures below 0.7 Pa at 400 W, the mono-clinic phase is absent. The deposition rate at these pressuresis extremely high ��15 Å /s�. Figure 5 compares the XRDpatterns for sputtering pressures of 0.2, 0.3, and 6.0 Pa at400 W. As evident from Figs. 5�a�, 5�b�, and 5�d� the mono-clinic phase is not present at pressures of 0.2 and 0.3 Pa. Inorder to study the crystal structure of films prepared at thesepressures in detail, XRD was conducted in powder diffrac-tion mode up to 2�=120° and is shown in Figs. 5�c� and5�e�. Many of the peaks in these XRD patterns match the

(a) (b)

100nm 100nm

100nm100nm

film top

film

Si substrate

film top

film

FIG. 4. Top-down �top� and cross-sectional �down� SEM of thin film pre-pared at �a� 150 W, 6.0 Pa with rf sputtering and �b� 400 W, 0.7 Pa with PDsputtering.

FIG. 5. �Color online� XRD patterns of ZrO2 films deposited at 400 W withPD at sputtering pressures of �a� 6.0 Pa, ��b� and �c�� 0.3 Pa, and ��d� and�e�� 0.2 Pa. The XRDs in �a�, �b�, and �d� are conducted in grazing inci-dence, while �c� and �e� were conducted in powder diffraction modes. Phase1 cubic Zr2O, phases 2 and 3 are cubic ZrO, and phase 4 is monoclinic

ZrO2.

J. Vac. Sci. Technol. A, Vol. 27, No. 3, May/Jun 2009

cubic-Zr2O phase22 and cubic-ZrO phase,23 suggesting thatthese films are mixtures of oxygen deficient zirconium oxidephases. EDS at 15 keV confirmed that the films prepared at0.2 and 0.3 Pa are oxygen deficient. The Zr:O ratio in thesefilms was 1.3 instead of the desired 0.5. The opaque nature ofthe films produced at these pressures supports these argu-ments. However, the films prepared at higher pressures suchas 6.0 Pa shown in Fig. 5�a� were more stoichiometric with aZr:O ratio of 0.6. �EDS x-ray map not shown here�. Due tothe reduced thickness of the films at these pressures �Fig.3�c��, the Si substrate also contributed to the EDS spectrumeven at accelerating voltages of as low as 5 keV. Hence theabsolute Zr:O ratio could not be measured for the films pre-pared at 0.6 Pa.

D. Effect of oxygen partial pressure on crystallization

In order to improve the stoichiometry of the high deposi-tion rate PD-sputtered films at 400 W and 0.2 Pa �O2 partialpressure of 0.04 Pa� mentioned in the previous section,higher oxygen partial pressures were employed during sput-tering. Figure 6 shows the effect of sputtering oxygen partialpressure on the crystallization and deposition rate. It is ob-served that a slight increase or decrease in the oxygen partialpressure by 0.02 Pa �25% or 7% O2� from 0.04 Pa �16% O2�can drastically change the crystal structure. The crystal struc-ture of the films at 0.02 Pa is hexagonal-Zr3O structure24 andat 0.06 Pa is monoclinic ZrO2. The deposition rate alsochanges drastically from over 15 Å /s to lesser than 5 Å /swith an increase in sputtering oxygen partial pressure. Above0.06 Pa, the deposition rate further reduces. This is attributedto the reduction in the Ar partial pressure, which is mainlyresponsible for sputtering Zr atoms from the target. However,there is no change in the crystallite size of the monocliniczirconia thin films grown at oxygen partial pressures above

FIG. 6. �Color online� Effect of sputtering oxygen content on the crystalli-zation and deposition rate for PD sputtering at 400 W and 0.2 Pa.

0.06 Pa.

581 Sethi et al.: Influence of reactive sputter deposition conditions 581

E. Effect of substrate temperature and substratematerial on crystallization

To investigate the effect of substrate temperature on zir-conia film crystallization, a liquid N2 substrate-cooling sys-tem was developed, which can reduce the substrate tempera-ture to −100 °C. Table II shows the effect of sputtering atlow temperatures on crystallization. The crystallite size atlow temperatures is smaller than that at room temperature forhigh powers. This decrease in the crystallite size is not no-ticeable at lower powers. In addition, the lower substratetemperature also results in significantly thinner films.

The deposition rates dropped from 117 to 37.8 Å /min at400 W and from 25.2 to 15.6 Å /s at 200 W in PD sputteringat 0.7 Pa. Cooling of the substrate results in cooling of thegas discharge and the target surface. This will result in asignificant decrease in the active energy of the zirconium andoxygen atoms at the target surface causing a decrease in thesputtering and the deposition yields. We also observe inTable II that the deposition on silicon �Si� substrates at lowertemperatures shows that the crystallite size is reduced. Sili-con has a high thermal conductivity of 139 W /m K in com-parison with SiO2 glass substrates25 �1–2 W /m K�, whichcould be responsible for this change in the crystallite size.26

The basic film morphology on the two different substrates isexpected to be similar because of the presence of native ox-ide �SiO2� on top of the Si substrate.

IV. DISCUSSION

The thermodynamics and kinetics of crystallization ofZrO2 thin films formed during sputtering are believed to becontrolled by mainly three factors. These factors are sputter-ing mode, ion bombardment, and substrate surface tempera-ture. In this section, each of these factors will be discussedwith respect to the film morphology results.

The sputtering pressure, power, and oxygen content ulti-mately control the growth mode, which defines the crystalli-zation and stoichiometry in the film. Two different sputteringmodes—“elemental” and “poisoned”—exist.27 Figure 7shows the curve for different growth conditions depicting thetransition from elemental to poisoned mode, which is a dis-continuous function with respect to the oxygen flow rate. Theobserved step transition in the ambient oxygen in Fig. 7 isdependent on both sputtering pressure and power.28 Sputter-

TABLE II. Effect of substrate temperature and type on the crystallization ofZrO2 thin films with PD sputtering at 200 and 400 W at a pressure of 0.7 Pa.

SubstrateTemperature

�°C�Power�W�

Crystallite size��

Glass 25 200 99�5Glass 25 400 271�5Glass −100 200 121�8Glass −100 400 212�6Si −100 200 79�3Si −100 400 87�4

ing conditions at oxygen flows prior to this step lead to non-

JVST A - Vacuum, Surfaces, and Films

stoichiometric film formation and high deposition rate con-ditions, as seen in Figs. 3 and 6. Films that are deposited atsputtering pressures of less than 0.7 Pa or sputtering O2 con-tent lesser than 16% correspond to this growth mode andpossess mixtures of nonstoichiometric zirconium oxidephases. Sputter conditions at oxygen flows greater than theelemental/poisoned mode transition lead to target poisoningand lower deposition rates �Fig. 3�. These films are depositedat higher pressures and higher O2 contents. Control of thesputtering conditions near the transition allow for both stoi-chiometric and high deposition rate of ZrO2. Creating repro-ducible films near this sharp transition requires a strict anddynamic control of the oxygen flow rate, which was possiblein this investigation.28

In the poisoned sputtering mode, the energetic particles�positive ions, negative ions, electrons, and neutrals� are alsoresponsible for the crystal structure evolution during sputter-ing. Since the movement of these energetic particles willdepend on the direction of the electric field, the film will seedifferent ion and electron energies during the positive andnegative cycle. In the case of group 4 oxides, it has beenobserved that these energetic particles are often oxygenions.29 These medium and the high energy ions �with rela-tively large fluxes� may be sputtered from the target as O2−

and O−, or produced by electron attachment at the bulkplasma side of the cathode sheath and accelerated over thecathode sheath.30 Estimation of their energies at the substrateis simple in the case of rf sputtering due to the symmetricaland simple nature of the waveform. The pressure �p� andtarget voltage �Vt� is empirically related to this ion energy asEi�Vt / p and ion flux, Ji� p in rf sputtering.31–34 In the caseof PD sputtering, the signal waveform is more complex andthere is an additional contribution from the positive over-shoot to the overall energy. An increase in pressure reducesthe ion energy and increases the ion flux, both of which havebeen found to increase crystallization. It has also been sug-gested that high ion energy provides a path for the film to be

FIG. 7. �Color online� Variation in the O2 content with its flow rate usingmass flow control of O2 gas showing the step transition from elementalmode to poisoned mode �at Ar base pressure of 0.7 Pa with a flow rate of8.4 SCCM� for ZrO2 thin film deposition.

at thermodynamic equilibrium, which is the monoclinic

582 Sethi et al.: Influence of reactive sputter deposition conditions 582

phase.35 This finding is consistent with the current studywhere an increase in the crystallite size is observed at higherpressures. However, at very low powers, this trend is re-versed, suggesting that an additional mechanism controls thecrystallization for thin film thickness.36 Figure 8 shows theeffect of rf sputtering power at 6.0 Pa on films of the samethickness �50 nm� at this possible critical thickness limit.The target currents and voltages are 0.8 A�131 V, 1.3 A�153 V, 1.8 A�166 V, and 2.3 A�172 V at 100, 200,300, and 400 W, respectively. It is seen in the figure that asthe target voltage and power increase, the tetragonal phase inthe film stabilizes in addition to the monoclinic phase. Thefilms prepared at 100 W are mostly amorphous, while thefilms prepared at 400 W are polycrystalline with both mono-clinic and tetragonal zirconia phases. Thus, in addition to thecritical thickness, it is surmised that the ion energy also playsa critical role in the crystallization. Lousa and Gimeno37 ob-served a threshold of momentum per atom �or equivalent ionenergy� to be 1600 �eV amu�1/2 to stabilize cubic phase inBN films. However, very high ion bombardment energy candamage the formation of particular crystal structure and cre-ate a textured or random structure as observed by Usov etal.38 Ar+ ion bombardment damage of atomic planes orientedin the �111� direction in yttria stabilized-ZrO2 at ion energyof 150 keV using ion beam assisted deposition was observedin that investigation.

The very act of sputtering can cause self-heating of thesubstrates and thus raise the surface temperature of the grow-ing film, the third factor mentioned previously. This heatingcould increase the adatom mobility on the substrate and pos-sibly increase crystallization. The overall substrate tempera-ture by self-heating can be as high as 200 °C within 60 min,as shown in Fig. 9�a�. However, the localized temperature atthe surface could be much higher than this value. Thisgreatly increases the adatom mobility and thus the thermo-dynamics become more favorable for the formation of high

FIG. 8. �Color online� Effect of rf sputtering conditions at 6.0 Pa, �a� 100 W,�b� 200 W, �c� 300 W, and �d� on XRD pattern of same thickness zirconiathin films. Phases �1� and �2� are monoclinic ZrO2 and tetragonal ZrO2,respectively. All films have nominal thickness of 50 nm.

temperature symmetrical crystal phases. This local tempera-

J. Vac. Sci. Technol. A, Vol. 27, No. 3, May/Jun 2009

ture rise, in conjunction with high defect density, surfaceenergy, and microstructure, could possibly explain the for-mation of cubic and tetragonal phases observed in this studyunder certain conditions.39 Stabilization of tetragonal struc-ture over monoclinic in ZrO2 films is widely regarded as aresult of a surface energy contribution.40 This certain degreeof control of the crystallographic effects in ZrO2 such asamorphization and preferential growth of some crystallo-graphic planes and phases, which is possible because of ionbombardment during sputtering, is not available in othergrowth processes such as chemical vapor deposition andatomic layer deposition. Furthermore, sputtering has an op-timum control of the microstructure of the films.

In order to determine the effects of substrate temperaturerise during sputtering on the crystallization, crystallizationwas studied as a function of substrate temperature in ZrO2

thin films. A 50-nm-thick film prepared at 150 W and 6.0 Paby rf sputtering was postdeposition annealed at various tem-peratures ranging from 200 to 1000 °C. Figure 9�b� showsthe crystallite size as a function of the annealing temperature

for the �1̄11� monoclinic ZrO2 peak. The peak broadeningdefined by FWHM decreases, while the peak intensity and

FIG. 9. �Color online� �a� Temperature evolution during PD and rf sputteringof ZrO2 thin films. �b� Effect of the annealing temperature on the crystallitesize corresponding to the strong XRD peak in monoclinic-ZrO2 phase as afunction of the annealing temperature.

crystallite size increase with the annealing temperature. Ad-

583 Sethi et al.: Influence of reactive sputter deposition conditions 583

ditionally, the peak shifts to higher Bragg angle and d spac-ing. Significant change in the crystallite size takes place onlyafter temperatures of 200 °C is attained. These results sug-gest that only the local temperature and not the overall sub-strate temperature affects the crystallization.39

V. CONCLUSIONS

ZrO2 thin films were deposited by reactive magnetronsputtering of a Zr target using PD and rf sources. Sputteringpower, pressure, oxygen content, substrate surface tempera-ture as well as substrate material have been shown to affectthe crystallization of ZrO2 thin films for both methods. Thecrystallite size increased with both power and pressure. Theeffect of pressure on crystallization was higher than the ef-fect of power. At very low pressures, the films were crystal-lized to a mixture of nonstoichiometric zirconium oxidephases. Increase in the sputtering oxygen content improvedthe Zr:O stoichiometry for ZrO2 and stabilized the mono-clinic phase. However, high oxygen content in the sputteringchamber did not affect the crystallite size. The crystallizationphases were slightly different in both power sources with rfsputtering showing inclination to stabilize the tetragonalphase. The two most important parameters for controllingcrystallization of ZrO2 films were the sputtering mode andion energy. The thickness of the films was also found to becritical; zirconia films under 50 nm were mainly amorphous.These amorphous films were obtained at high pressures andlow powers with a deposition rate of 10–15 Å /min.

ACKNOWLEDGMENTS

This research was conducted under MURI grant from Of-fice of Naval Research under Contract No. N000140510541.The authors would like to thank Nichole Wonderling for herhelp in performing the x-ray diffraction analysis.

1M. V. Fazio and H. C. Kirbie, Proc. IEEE 92, 1197 �2004�.2J. W. McPherson, J. Kim, A. Shanware, H. Mogul, and J. Rodriguez,IEEE Trans. Electron Devices 50, 1771 �2003�.

3D. H. Kuo and C. H. Chien, Thin Solid Films 429, 40 �2003�.4S. B. Amor, B. Rogier, G. Baud, M. Jacquet, and M. Nardin, Mater. Sci.Eng., B 57, 28 �1998�.

5Z. Ning-Lin, S. Zhi-Tang, S. Qin-Wo, and L. Cheng-Lu, Chin. Phys. Lett.20, 273 �2003�.

6P. Gao, L. J. Meng, M. P. dos Santos, V. Teixeira, and M. Andritschky,Thin Solid Films 377–378, 557 �2000�.

7M. S. Wong, W. J. Chia, P. Yashar, J. M. Schneider, W. D. Sproul, and S.

JVST A - Vacuum, Surfaces, and Films

A. Barnett, Surf. Coat. Technol. 86–87, 381 �1996�.8P. Gao, L. J. Meng, M. P. dos Santos, V. Teixeira, and M. Andritschky,Vacuum 56, 143 �2000�.

9Y. Tanaka, N. Ohnuma, K. Katsunami, and Y. Ohki, IEEE Trans. Electr.Insul. 26, 258 �1991�.

10T. W. Giants, IEEE Trans. Dielectr. Electr. Insul. 1, 991 �1994�.11R. Nath and M. M. Perlman, IEEE Trans. Electr. Insul. 24, 409 �1989�.12D. Triyoso et al., J. Electrochem. Soc. 151, F220 �2004�.13A. V. Wagner, G. W. Johnson, and T. W. Barbee, Mater. Res. Soc. Symp.

Proc. 574, 219 �1999�.14K. Wu, T. Okamoto, and Y. Suzuoki, J. Appl. Phys. 98, 114102 �2005�.15Y. Ye, S. Z. Zhang, F. Dogan, E. Schamiloglu, J. Gaudet, P. Castro, M.

Roybal, M. Joler, and C. Christodulou, in Digest of Technical Papers—2003 14th International Pulsed Power Conference, edited by M. Giessel-mann and A. Neuber �IEEE, Piscataway, NJ, 2003�, pp. 719–722.

16J. J. O’Dwyer, The Theory of Dielectric Breakdown of Solids �OxfordClarendon, London, 1964�.

17B. D. Cullity, Elements of X-Ray Diffraction �Addison-Wesley, Reading,MA, 1978�.

18H. F. McMurdie, M. C. Morris, E. H. Evans, B. Paretzkin, and W. Wong-Ng, Powder Diffr. 1, 40 �1986�.

19A. Mehner, H. Klümper-Westkamp, F. Hoffmann, and P. Mayr, ThinSolid Films 308–309, 363 �1997�.

20J. Málek, L. Benes, and T. Mitsuhashi, Powder Diffr. 12, 96 �1997�.21J. A. Thornton, J. Vac. Sci. Technol. 11, 666 �1974�.22V. I. Khitrova and V. V. Klechkovskaya, Kristallografiya 30, 126 �1985�.23N. Schönberg, Acta Chem. Scand. �1947-1973� 8, 627 �1954�.24B. Holmberg and T. Dagerhamn, Acta Chem. Scand. �1947-1973� 15,

919 �1961�.25W. F. Gale and T. C. Totemeier, Smithells Metals Reference Book

�Butterworth-Heinemann, Burlington, MA, 2004�.26R. F. Messier, Ph.D. thesis, The Pennsylvania State University, 1973.27W. D. Sproul, Vacuum 51, 641 �1998�.28W. D. Sproul, D. J. Christie, and D. C. Carter, Thin Solid Films 491, 1

�2005�.29J. M. Ngaruiya, O. Kappertz, S. H. Mohamed, and M. Wuttig, Appl.

Phys. Lett. 85, 748 �2004�.30S. Mráz and J. M. Schneider, Appl. Phys. Lett. 89, 051502 �2006�.31A. D. Glew, R. Saha, J. S. Kim, and M. A. Cappelli, Surf. Coat. Technol.

114, 224 �1999�.32J. H. Lee, D. S. Kim, Y. H. Lee, and B. Farouk, J. Electrochem. Soc. 143,

1451 �1996�.33D. L. Smith, Thin-Film Deposition: Principles and Practice �McGraw-

Hill, New York, NY, 1995�.34H. Bartzsch, P. Frach, and K. Goedicke, Surf. Coat. Technol. 132, 244

�2000�.35Stanislav Mraz and Jochen M. Schneider, J. Appl. Phys. 100, 023503

�2006�.36S. Kremmer, H. Wurmbauer, C. Teichert, G. Tallarida, S. Spiga, C. Wi-

emer, and M. Fanciulli, J. Appl. Phys. 97, 074315 �2005�.37A. Lousa and S. Gimeno, Diamond Relat. Mater. 10, 1347 �2001�.38I. O. Usov, P. N. Arendt, J. R. Groves, L. Stan, and R. F. DePaula, Nucl.

Instrum. Methods Phys. Res. B 240, 661 �2005�.39Fletcher Jones, J. Vac. Sci. Technol. A 6, 3088 �1988�.40R. G. Garvie, J. Phys. Chem. 69, 1238 �1965�.