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Materials Chemistry and Physics 126 (2011) 515–523

Contents lists available at ScienceDirect

Materials Chemistry and Physics

journa l homepage: www.e lsev ier .com/ locate /matchemphys

nfluence of hydrogen annealing on the properties of hafnium oxide thin films

.F. Al-Kuhaili ∗, S.M.A. Durrani, I.A. Bakhtiari, M.A. Dastageer, M.B. Mekkihysics Department, King Fahd University of Petroleum and Minerals, Dhahran 31261, Saudi Arabia

r t i c l e i n f o

rticle history:eceived 13 August 2010eceived in revised form 18 October 2010ccepted 10 January 2011

ACS:8.37.Ps1.05.Gc1.15.Fg1.40.Tv1.70.Jb

a b s t r a c t

Thin films of hafnium oxide were deposited by electron beam evaporation, and were subsequentlyannealed in hydrogen. X-ray diffraction, X-ray photoelectron spectroscopy, atomic force microscopy,photoluminescence, spectrophotometry, and current–voltage measurements were performed to inves-tigate the structural, chemical, optical, and electrical properties of the films. As-deposited films wereamorphous and nearly stoichiometric. Annealing led to crystallization of the films, and reduction of stoi-chiometry. Photoluminescence measurements revealed the presence of oxygen-related defects. Optically,the films were transparent with a wide band gap, and this was not affected by hydrogen annealing. More-over, the films were suitable as anti-reflection coatings on silicon. The electrical resistivity of the filmswas significantly reduced as a result of annealing.

© 2011 Elsevier B.V. All rights reserved.

eywords:afnium oxide-Beam evaporationydrogen annealingtructural propertieshemical propertiesptical properties

lectrical properties

. Introduction

Hafnium oxide (HfO2) is a material that is characterized byxcellent chemical, thermal, and mechanical stability, and pos-esses a high melting point [1]. It is a wide band gap semiconductorhat is transparent from the deep ultraviolet to the mid infraredpectral ranges [2]. Moreover, it has a relatively high refrac-ive index and shows the highest laser damage threshold [3].hus, HfO2 has been used extensively in optical coatings, such asnterference filters [4] and anti-reflection coatings [5]. For micro-lectronic applications, HfO2 is characterized by a high dielectriconstant, along with good thermodynamic and mechanical stabil-ty with silicon [6]. Therefore, HfO2 has emerged as the leadingandidate for the replacement of silicon oxide as the gate in metal-xide-semiconductor field effect transistors (MOSFETs) [7]. Otherpplications have been reported for hafnium oxide, including gas

ensors [8] and protective coatings [4].

Thermally evaporated films consist of cylindrical columns withoids in between [9]. Such a microstructure leads to deteriora-ion of the properties of the films, such as optical homogeneity,

∗ Corresponding author. Tel.: +966 3 860 3747; fax: +966 3 860 2293.E-mail address: kuhaili@kfupm.edu.sa (M.F. Al-Kuhaili).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.01.036

environmental stability and adhesion [9]. Solutions to this prob-lem include using deposition methods with high particle energy,bombardment with energetic ions during deposition, and sub-strate heating [9–11]. Another solution is post-deposition thermalannealing, which leads to densification of the films, and thus higherpacking density [9–11]. Another advantage of thermal annealing isthe crystallization of the films, and consequently a reduction in thedensity of structural defects.

Annealing in hydrogen has been widely used in the microelec-tronics industry to remove native oxides and passivate danglingbonds [12]. Recently, numerous advantages of annealing in hydro-gen have emerged. For example, it has been established thatannealing zinc oxide thin films in hydrogen led to a reduction ofthe resistivity of the films by several orders of magnitude with-out reduction in transmittance [13]. Moreover, hydrogen annealingwas suggested as a practical method for controlled n-type doping ofzinc oxide [14]. Annealing silicon in hydrogen was used to fabricatemicrodisks with quality factors in excess of 3 × 105 [12]. Annealingof wide band gap semiconductors, such as gallium nitride and sili-

con carbide, in hydrogen led to significant improvement in surfaceand interface quality [15]. Finally, hydrogen annealing was used toremove oxides from elemental nanowires [16].

Previous studies have mainly focused on annealing HfO2 thinfilms in nitrogen [6,17–27] or oxygen [24,28–30]. Several studies

5 emistry and Physics 126 (2011) 515–523

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0104]. The most intense peak in the patterns was the (1 1 1) peakat 2� = 31.6◦. This peak was used to estimate the crystallite sizesusing the Scherrer formula, and these are given in Table 1. Thefilms deposited on fused silica showed a similar behavior, except

Table 1Summary of the structural properties of the films.

Temperature Crystallite size Roughness Maximumheight

16 M.F. Al-Kuhaili et al. / Materials Ch

ere reported on the annealing of ultrathin (thickness < 5 nm) HfO2lms in a forming gas containing hydrogen (90–95% N2 + 5–10% H2)24,25,28,31]. In addition, there are two studies on the annealingf ultrathin HfO2 films in pure hydrogen. One study reported therap density at the HfO2/Si interface in films with a thickness of–12 nm deposited by atomic layer chemical vapor deposition andnnealed in hydrogen at 400 or 550 ◦C [21]. Another study reportedhe chemical composition of the HfO2/Si interlayer in films with ahickness of 5–10 nm deposited by chemical vapor deposition andnnealed in hydrogen at 400 ◦C [22].

In this paper, we report the influence of hydrogen annealingn the structural, chemical, optical, and electrical properties ofafnium oxide thin films deposited by reactive electron beam evap-ration.

. Experimental details

Hafnium oxide thin films were deposited by electron beam (e-beam) evapo-ation using hafnium oxide (HfO2) pellets (Balzers, 99.9% purity). The films wererepared in a Leybold L560 box coater. The material was slowly out-gassed beforevaporation. The system was pumped to a base pressure of 5 × 10−4 Pa. Films wereeposited in an oxygen atmosphere. The partial pressure of oxygen during deposi-ion was 5 × 10−2 Pa. Two types of substrates were used: single-side polished (1 0 0)-doped n-type silicon substrates with a resistivity of 1–100 � cm, and fused sil-ca (SiO2) substrates that were single- or double-side polished. The substrates wereltrasonically cleaned in acetone and methanol, rinsed with de-ionized water, andried with nitrogen, prior to deposition. The films were deposited on unheated sub-trates. The substrates were rotating during deposition, and the source-to-substrateistance was 40 cm. The evaporation rate (0.5 nm s−1) and thickness of the filmsere controlled by a quartz crystal thickness monitor. After deposition, the filmsere annealed in a horizontal tube furnace. The furnace was first evacuated down

o 10−5 Pa and was then kept under flowing hydrogen gas of 99% purity at a fixedow rate of 100 sccm (standard cubic centimeter per minute). The samples werennealed at temperatures of 400, 600, and 800 ◦C for 4 h and were cooled down tooom temperature under the same flow rate inside the furnace before removinghem from the furnace.

The thickness of the as-deposited films was measured using a surface profilome-er (Ambios XP-2), and was found to be 490 nm. The thickness was also foundptically. The structure of the films was investigated by X-ray diffraction (XRD)sing a Shimadzu XRD-6000 diffractometer, employing Cu K� (1.54 A) radiation.RD was performed for films deposited on silicon and fused silica. The surfaceorphology of the films was examined by tapping mode atomic force microscopy

AFM) (Veeco Innova diSPM). The sample surface was probed with a silicon tip of0 nm radius oscillating at its resonant frequency of 300 kHz. The scan area was�m × 2 �m, and the scan rate was 1 Hz. AFM images were obtained for filmseposited on silicon substrates. The chemical composition of the films was studiedsing X-ray photoelectron spectroscopy (XPS), and was performed in a VG Scien-ific ESCALAB MKII spectrometer equipped with an Al K� (1486.6 eV) X-ray source.rior to the XPS analysis, the samples were transferred in air to the XPS analy-is chamber. The C 1s peak of hydrocarbon contamination, at a binding energy of84.5 eV, was used as an energy reference. During the XPS analysis, the samplesere maintained at ambient temperature at a pressure of 5 × 10−7 Pa. XPS was per-

ormed on samples deposited on silicon, since the charging effect will be less forhese substrates. Photoluminescence measurements were performed at room tem-erature using a Shimadzu RF5301PC spectrofluorometer whose excitation source

s a broad band 150 W xenon lamp. Samples deposited on fused silica were usedor the photoluminescence studies. Normal-incidence transmittance was measuredver the wavelength range 200–1000 nm using a Jasco V-570 double beam spec-rophotometer. Normal-incidence reflectance was measured over the wavelengthange 600–1400 nm using the same instrument. Transmittance was measured forlms deposited on transparent fused silica substrates, and reflectance was measured

or films deposited on silicon and single-side polished fused silica substrates. Theesistivity of the films was determined by measuring the current passing throughhe sample at a fixed bias voltage. For the current measurement, two thin film goldlectrodes (thickness 50 nm, separation 0.5 mm) were deposited on the surface ofhe films. Current–voltage characteristics were measured using an automated sys-em employing a Keithley 238 source measure unit. Resistivity was measured usingamples deposited on fused silica.

. Results and analysis

.1. Structural properties

Hafnium oxide is a material that exhibits several polymorphs,ncluding the monoclinic, tetragonal, orthorhombic, and cubic

Fig. 1. XRD patterns of the films deposited on silicon substrates, and subsequentlyannealed in hydrogen at the indicated temperatures.

structures. Under normal conditions (room temperature and atmo-spheric pressure), the equilibrium phase is the monoclinic [20].Phase transitions to other phases take place at very high tem-peratures (>1700 ◦C) [20]. Under some special conditions, such asthin film deposition, polymorphs can exist at much lower tem-peratures [20]. This was verified in several studies [6,29,32]. Thecrystallinity and existence of polymorphs depend on the substratecrystallinity and temperature, deposition technique, and anneal-ing atmosphere and temperature. Moreover, it was shown thatthere is a critical minimum thickness for the onset of crystallization[6].

Fig. 1 shows XRD patterns of hafnium oxide thin films depositedon silicon and annealed at different temperatures in hydrogen. TheXRD pattern of the as-deposited films showed a large and broadbackground with no diffraction peaks. Such a pattern is character-istic of amorphous films. At an annealing temperature of 400 ◦C,the films had an XRD pattern with small peaks superimposed onthe amorphous background, indicating the onset of crystallization.The crystallinity was substantially improved with the increase inannealing temperature, as was exhibited by the higher intensityand narrower widths of the XRD peaks. This indicates that graingrowth is thermally activated [19]. The multi peaks in the XRDpatterns reflect the polycrystalline nature of the films. All peakswere ascribed to the monoclinic phase [JCPDS card no. 00-034-

Ta (◦C) Si (nm) SiO2 (nm) Rrms (nm) H (nm)

As-deposited – – 1.31 9.74400 6.0 – 0.96 7.67600 10.8 10.0 1.36 8.11800 11.2 11.9 1.42 12.98

M.F. Al-Kuhaili et al. / Materials Chemistry and Physics 126 (2011) 515–523 517

fnium

tsfi

ostwmcrTFatfirTh8ap(

TS

Fig. 2. Two- and three-dimensional AFM images of as-deposited ha

hat the films annealed at 400 ◦C were still amorphous. This clearlyhows the influence of the substrate on the crystallinity of thelms.

Fig. 2 shows AFM images of as-deposited and annealed hafniumxide thin films. The morphology of the as-deposited filmshows a uniform columnar microstructure. The morphology ofhe annealed films shows a distorted columnar microstructure,ith the column consisting of spherical grains. The average root-ean-square roughness (Rrms) and the maximum height of the

olumns (H) are given in Table 1. The as-deposited films had lowoughness, which decreased as the films were annealed at 400 ◦C.here are three effects taking place upon annealing at 400 ◦C.irst, the intensity of columnar microstructure was decreased,s revealed by the reduction of the maximum height. Second,he films became more compact, with lower thickness. Third, thelms became polycrystalline. The first two effects tend to reduceoughness, whereas the third effect tends to increase roughness.he reduction of roughness suggests that the first two effects

ad a dominant role. When the films were annealed at 600 or00 ◦C, the roughness increased. Since the thicknesses of thennealed films were comparable, the increase in roughness wasossibly due the increase in crystallite size, as revealed by XRDTable 1).

able 2ummary of the chemical properties of the films.

Ta (◦C) Binding energies (eV)

Hf 4f7/2 Hf 4f5/2 Hf 4d5/2 Hf 4f3/2 O 1s (A)

As 16.6 18.1 212.9 223.5 530.2400 16.4 18.0 213.0 223.7 530.0600 16.4 18.0 212.9 223.6 530.1800 16.2 17.8 212.3 223.7 530.2

a ı4f is the energy separation between the Hf 4f7/2 and 4f5/2 levels.b ı4d is the energy separation between the Hf 4d5/2 and Hf 4d3/2 levels.c � is the energy separation between the Hf 4f7/2 and O 1s levels.d r is the ration of the B component to the total O 1s peak.

oxide thin films (a), and films annealed at 800 ◦C in hydrogen (b).

3.2. Chemical properties

The chemical state of the films was investigated using XPS. Widesurvey scans of the films revealed that the only elements presentwere the constituent elements (hafnium and oxygen) in addition toadventitious carbon. No silicon (from the substrates) was detected,indicating that the substrates were completely obscured by the rel-atively thick films. In addition to the wide survey scans, detailedhigh-resolution spectra were obtained in the Hf 4f, Hf 4d, and O1s core level regions. The spectra are shown in Fig. 3, and the XPSresults are summarized in Table 2.

The Hf 4f spectrum consists of two peaks (4f7/2 and 4f5/2) thatare due to spin-orbit splitting. In order to resolve these peaks, theHf 4f spectrum was de-convoluted into two components using aGaussian/Lorenzian mixed function employing Shirley backgroundcorrection, as shown in Fig. 3. The binding energy (BE) of the Hf4f7/2 peak was 16.4 ± 0.2 eV. This value is in close agreement withthe value of 16.2 ± 0.4 eV, reported for bulk HfO2 [1]. Moreover,

the spin-orbit splitting (ı4f) was 1.5–1.6 eV, and thus is in closeagreement with the reported values of 1.4 eV [5] and 1.7 eV [33].Several XPS studies on ultrathin HfO2 films reported that the BE ofthe Hf 4f7/2 peak was almost 1.0 eV higher than our values. Thereare two reasons for this difference. The first reason is the value of

Energy separations (eV) Ratios

O 1s (B) ı4fa ı4d

b �c rd [O/Hf]

531.8 1.5 10.6 513.6 0.236 2.06532.1 1.6 10.6 513.6 0.234 1.94531.5 1.6 10.6 513.7 0.221 1.91532.0 1.6 10.5 514. 0 0.140 1.91

518 M.F. Al-Kuhaili et al. / Materials Chemistr

F(

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ig. 3. XPS spectra of hafnium oxide thin films in the Hf 4f region (a), Hf 4d regionb), and O 1s region (c).

he BE of the carbon C 1s peak that is used as the energy reference.econd, the shifting of the Hf 4f7/2 to higher BE values is caused byhe formation of the Hf–O bond in the vicinity of silicon [17,29].

The oxygen O 1s peak was similarly resolved into two com-onents: a low-BE component (A) at a BE of 530.1 ± 0.1 eV, andhigh-BE component (B) at a BE of 531.5–532.1 eV. The lower-BE

omponent corresponds to oxygen in HfO2, where the reported val-es are 529.7 eV [5], 530.5 eV [29], and 530.9 eV [28]. The energyeparation (�) between the low-BE O 1s component and the Hff7/2 peak is 513.8 ± 0.2 eV. The value of � in pure HfO2 powder

s 513.7 ± 0.1 eV [34]. The high-BE component was attributed tohysi-sorbed water or OH groups [5,28]. Thus, this component maye related to the porosity of the films. The ratio (r) of this compo-ent to the total O 1s peak, as calculated from the areas under theeaks, is given in Table 2. The strength of the high-BE componentrogressively decreased as the annealing temperature increased.

The atomic concentration of the elements was calculated fromhe normalized areas of the Hf 4d5/2 and O 1s peaks, taking intoccount the atomic sensitivity factor of each element [35]. The Hfd5/2 peak was used because it is well resolved and separated fromhe Hf 4d3/2 peak. The results of these calculations are shown byhe [O/Hf] ratios given in Table 2. It should be noted that the [O/Hf]

atio, as determined by XPS, is over-estimated. This is due to theresence of hafnium hydroxide [1]. In our case, this is supported byhe high-BE O 1s component. Moreover, XPS is a surface techniquehat probes only a few top monolayers, and thus cannot reveal theO/Hf] ratio in the bulk of the film. Indeed, the photoluminescence

y and Physics 126 (2011) 515–523

results suggest the existence of oxygen vacancies. Nevertheless, the[O/Hf] ratio (Table 2) showed the expected trend, i.e. it decreasedupon annealing in hydrogen. This behavior is consistent with thevariation in the values of the separation � between the Hf 4f7/2 andO 1s peaks.

3.3. Optical properties

3.3.1. PhotoluminescencePhotoluminescence (PL) was measured for films deposited

on fused silica substrates. For these substrates, the as-depositedfilms and those annealed at 400 ◦C were amorphous, whereasthe films annealed at 600 or 800 ◦C were polycrystalline. PLspectra were collected in two spectral regions. The first regionwas in the ultraviolet (UV) range, with emission wavelengths�em = 325–400 nm (3.10–3.81 eV). The excitation wavelength was�ex = 220 nm (5.64 eV), and the instrumental bandwidth was 3 nm.The second region was in the visible and near-infrared rangeswith �em = 650–750 nm (1.65–1.95 eV). The excitation wavelengthwas �ex = 410 nm (3.02 eV), and the instrumental band width was10 nm. The PL spectra are shown in Fig. 4.

In the UV range, the spectra consisted of a main emission bandcentered at 3.37 eV, with shoulders at 3.51 eV and 3.30 eV. Theintensity of the PL spectra increased as the annealing temperatureincreased. For the highest annealing temperature (800 ◦C), the PLspectrum was split into two peaks at 3.30 eV and 3.42 eV. Such asplit in the luminescence spectrum for the films annealed at thehighest temperature was observed in HfO2 films deposited by RFsputtering and atomic layer deposition [26]. The origin of the emis-sion band at 3.37 eV was attributed to oxygen vacancies [36]. Thesevacancies form an energy level below the conduction band, andelectronic transitions from the valence band to the “vacancy band”give rise to the observed emission spectrum [36]. Other authorsobserved an emission band at 3.5 eV, and attributed it to interstitialoxygen rather than oxygen vacancies [26]. In that case, the emis-sion band shifted to 4.1 eV upon annealing the films at 900 ◦C innitrogen or oxygen [26].

In the visible and near-infrared ranges, an emission band wasobserved at 1.77 eV, with a shoulder at 1.70 eV. The PL intensityincreased as the annealing temperature was increased, and shifteddown by 0.02 eV upon annealing at 800 ◦C. Similar bands wereobserved in e-beam evaporated films, and were attributed to oxy-gen vacancies [30].

3.4. Optical constants

The transmittance (T) spectra of the films are shown in Fig. 5. Ingeneral, the transmittance of the films decreased as the annealingtemperature increased, with the films annealed at 400 ◦C showingthe largest reduction in transmittance. With the exceptions of thefilms annealed at 400 ◦C, the samples were transparent down toa wavelength of 250 nm. Similar to photoluminescence, transmit-tance was measured for films deposited on fused silica substrates.

In the visible and near-infrared wavelength ranges, thetransmittance spectra were fitted using the equation for the trans-mittance of a thin film on a transparent substrate [37,38]:

T = 16ns(n2 + k2)ˇA + Bˇ2 + 2ˇ[C cos(4�nd/�) + D sin(4�nd/�)]

(1)

with

A = [(n + 1)2 + k2][(n + ns) + k2]

B = [(n − 1)2 + k2][(n − ns) + k2]

C = −(n2 − 1 + k2)(n2 − n2s + k2) + 4k2ns

D = 2kns(n2 − 1 + k2) + 2k(n2 − n2s + k2)

M.F. Al-Kuhaili et al. / Materials Chemistry and Physics 126 (2011) 515–523 519

F ed hau

wcdˇsiC

n

w

tttp

Fas

ig. 4. Room-temperature photoluminescence spectra of as-deposited and annealltraviolet region.

here n is the refractive index of the film, k is the extinctionoefficient of the film, ns is the refractive index of the substrate,is the thickness of the film, � is the wavelength of light, and= exp(−4�kd/�). In order to fit the experimental transmittance

pectra using Eq. (1), models for the dispersion of n and k must bemplemented. The refractive index of the films was modeled by aauchy dispersion equation [10]:

(�) = no + A1

�2+ A2

�4(2)

here no, A1, and A2 are constants.The extinction coefficient of the films was modeled by an equa-

ion that takes into account the major absorption mechanisms inhis spectral region: the Urbach tail, defect absorption, multipho-on absorption, and light scattering [39]. The �-dependence of theserocesses is complicated, but if the total absorption coefficient due

ig. 5. Normal-incidence transmittance spectra of as-deposited and hydrogen-nnealed hafnium oxide thin films. The transmittance spectrum of a bare silicaubstrate is also indicated.

fnium oxide thin films: (a) in the visible and near-infrared regions, and (b) in the

to these processes is small, k can be expressed as:

k(�) = ko + A3

�� (3)

where ko, A3, and � are constants. For all films, � was found to be2 ± 0.2.

The experimental transmittance spectra were fitted using Eq. (1)with Eqs. (2) and (3) as models for the optical constants. The fittingparameters were no, A1, A2, A3, ko, and d. The substrate refractiveindex was taken from reference [40]. The calculated transmittancespectra, employing the above models for n and k successfully fit-ted the measured spectra throughout the visible and near-infraredwavelength ranges with a correlation that was better than 99%,as shown in Fig. 6a. The best-fit parameters are shown in Table 3,and they were used to calculate the optical constants of the films,which are shown by the dispersion curves of Fig. 6b (for n) andFig. 6c (for k). The best-fit values for the thickness were within5% of those obtained from the surface profilometer. Our valuesfor the refractive index of the as-deposited films decreased from1.846 (� = 500 nm) to 1.834 (� = 600 nm). These values are close tothe value of 1.86 reported for the refractive index of HfO2 films at� = 550 nm [9,41]. There is a wide scatter in the reported values ofthe refractive index of HfO2 films. Values as large as 2.16 [2] and aslow as 1.41 [42] were reported for the refractive index of HfO2 filmsat � = 633 nm. This significant variation reflects the critical depen-dence of the refractive index on the deposition conditions and theanalytical techniques used to derive them.

The refractive index is related to the packing density of the films.Evaporated thin films have low packing density as a result of colum-nar growth, which introduces voids in their microstructure [43,44].The lower value of the packing density can contribute to lower-ing the refractive index [45]. Moreover, in the present method ofcalculation, the refractive index was derived from transmittancemeasurements. Transmittance decreases due to the increase of

absorption associated with structural and chemical defects that arepresent in the films. The density of the structural defects is propor-tional to the degree of structural disorder and lack of crystallinity,whereas the chemical defects are related to the oxygen vacanciesthat are associated with the sub-stoichiometry of the films.

520 M.F. Al-Kuhaili et al. / Materials Chemistry and Physics 126 (2011) 515–523

Table 3Summary of the best-fit parameters used in fitting the experimental transmittance spectra of hafnium oxide thin films.

Ta (◦C) no A1 (nm2) A2 (nm4) A3 (nm2) ko d (nm) Eo (eV) Ed (eV)

3 8 111 −3

273134329

sohs[ckltTmdt

F((

As-dep. 1.818 3.06 × 10 9.87 × 10 –400 1.807 2.48 × 104 −1.69 × 109 5600 1.778 2.34 × 104 −1.21 × 109

800 1.767 2.34 × 104 −1.18 × 109

The refractive index of the as-deposited films (1.84 at 550 nm,ee Fig. 6b) was less than the bulk value of 2.08 [42]. This low valuef the film refractive index indicates that the as-deposited filmsad relatively low packing density. Lowering of the packing den-ity is caused by the incorporation of oxygen during film growth44], which may create voids that absorb moisture [42]. Moreover,ollisions of the evaporated species with O2 molecules reduce theirinetic energy before reaching the substrate, and this will result inower packing density [42]. When the films were annealed at 400 ◦C,heir refractive index increased significantly, as shown in Fig. 6b.

he increase in the refractive index for the films annealed at 400 ◦Cay be attributed to densification, and chemical and structural

efects. The densification of the films is verified by the reduction ofhickness upon annealing. Annealing provides thermal energy that

ig. 6. (a) Goodness of fit of transmittance using the models presented by Eqs.1)–(3). (b) Dispersion curves of the refractive indices of hafnium oxide thin films.c) Dispersion curves of the extinction coefficients of hafnium oxide thin films.

2.01 × 10 497 10.6 24.12.51 × 10−3 474 8.8 20.49.94 × 10−4 463 8.2 18.21.50 × 10−3 456 8.2 17.6

increases the mobility of the atoms of the films, thereby increas-ing the packing density of the films [18]. This packing will reducethe voids within the film and increase absorption. The increase ofdefects is also verified by the significant increase of the extinctioncoefficient, which is related to the absorption coefficient (˛) by:k = ˛�/4�. In this spectral region, absorption is caused by band-tail absorption and structural defects rather than by inter-bandelectronic transitions that prevail in the fundamental absorptionregion. A similar behavior was observed in ZnO films annealed inair at 350 ◦C [46], and was attributed to the decrease in the trans-mittance of the films. However, the mobility provided by thermalenergy was not sufficient to initiate crystallization, and thus thefilms remained amorphous.

When the films were annealed at 600 or 800 ◦C, there was a sig-nificant reduction in the refractive index. Such a reduction in therefractive index as the films were transformed from an amorphousto a crystalline state was observed in e-beam deposited HfO2 films[30,47]. In our case, the films annealed at 600 or 800 ◦C showedlower values of the thickness and lower stoichiometry. Yet, theirrefractive indices were almost equal to the value of the refractiveindex before annealing. This reduction supports the assumptionthat the increase in the refractive index in the films annealed at400 ◦C was mainly caused by an increase in the density of structuraldefects. The films annealed at 600 or 800 ◦C were polycrystalline,and thus had a lower density of structural defects. The decrease ofthe refractive indices of the films annealed at 600 or 800 ◦C com-pared to the as-deposited films can be attributed to a decrease inoptical absorption in the visible region by the annealed films. Inthe visible region, the refractive index can be modeled by a singleoscillator model [48]:

n2 = 1 + EdEo

E2o − E2

(4)

where E is the photon energy (E = hc/�, where h is Planck’s constantand c is the speed of light in vacuum), Eo is the average excitationenergy for electronic transitions, and Ed is the dispersion energywhich is a measure of the strength of inter-band optical transi-tions [49]. Our values of the refractive indices were fitted using theabove equation. The values of Eo and Ed obtained from such fits areincluded in Table 3. They clearly show a higher value of Ed for theas-deposited films, and thus a higher value of the refractive indexof these films at longer wavelengths.

3.4.1. Band gapThe transmittance of a film can also be written as [50]:

T = (1 − R)2 exp(−˛d)1 − [R2 exp(−2˛d)]

(5)

where R is the reflectance of the film. In the fundamental absorptionregion (˛ > 104 cm−1), the reflectance of our films did not exceed0.18. Thus, the second term in the denominator of equation (5) [R2

exp(–2˛d)] is negligible (<0.01). Therefore, Eq. (5) can be rewrittenas T ≈ (1 − R)2e−˛d. From which,

˛ = 1d

ln

((1 − R)2

T

)(6)

M.F. Al-Kuhaili et al. / Materials Chemistry and Physics 126 (2011) 515–523 521

Ffi

Trasod[

i

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wTcrfida

twCttrgsdtucbr[gng

Table 5Refractive index values in the infrared region.

√

TE

ig. 7. Absorption coefficient of as-deposited and annealed hafnium oxide thinlms.

he absorption coefficient was calculated using Eq. (6), and theesults are shown in Fig. 7. The absorption coefficients of themorphous films (as-deposited films and films annealed at 400 ◦C)howed a monotonic increase with photon energy, whereas thosef polycrystalline films (annealed at 600 or 800 ◦C) showed a shoul-er which may be related to peculiarities of the monoclinic HfO251].

In the fundamental absorption region, the absorption coefficients related to the band gap (Eg) by the relation [52]:

= ˛o

E(E − Eg)� (7)

here ˛o is a constant with values between 105 and 106 cm−1 [52].he constant � depends on the type of transitions involved: � = ½orresponds to a direct transition, and � = 2 corresponds to an indi-ect transition. The absorption coefficient (Eq. (6) and Fig. 7) wastted using Eq. (7), taking ˛o and Eg as fitting parameters. Bothirect and indirect transitions were tried. The best-fit parametersre given in Table 4, along with the correlation (R) of the results.

The value of the band gap of bulk monoclinic HfO2 was reportedo be 5.55 eV [53]. The values of the band gap of HfO2 thin filmsere critically evaluated by Aarik et al. [51] and Cheynet et al. [54].heynet et al. reported that the electronic structure of HfO2 close tohe band gap is rather complicated, leading to direct and indirectransitions that are close in energy [54]. Theoretically, the indi-ect transition corresponds to a → B transition, and has a bandap energy of 5.65 eV [54]. This is followed by a direct B → B tran-ition and has a band gap energy of 5.9 eV [54]. Thus, there is aifference of 0.25 eV between these transitions. It was argued thathese theoretical values were higher than the experimental val-es due to the negligence of excitonic effects in the theoreticalalculation [54]. Experimentally, the reported values of the directand gap were 5.64 eV [41], 5.68 eV [2], and 5.68–5.72 eV [51]. The

eported experimental values of the indirect band gap were 5.3 eV54], 5.54 eV [4], and 5.50–5.72 eV [26]. Other values of the bandap in the range 5.37–6.0 eV were reported without specifying theature (direct or indirect) of the band gap. The values of the bandap are not only affected by the assumed nature of the gap, but

able 4nergy gap values of hafnium oxide thin films.

Ta Direct

(◦C) ˛o × 105 cm−1 Eg (eV) RAs-dep. 2.41 5.63 0.880400 4.42 5.61 0.927600 3.10 5.63 0.937800 3.46 5.63 0.941

� (nm) ns ns n (As-dep.) n (400 ◦C) n (600 ◦C) n (800 ◦C)

1310 3.505 1.872 1.82 1.82 1.79 1.781550 3.476 1.864 1.82 1.82 1.79 1.78

also by the experimental method (UV spectroscopy or ellipsome-try) used in its determination [41,54]. Park et al. showed that theindirect band gap of HfO2 films increased upon annealing in nitro-gen [26]. Other studies reported the same trend without specifyingthe nature of the gap [6,18,23,27]. This increase was attributed tothe polycrystallization of the films [23]. On the other hand, Wanget al. reported a decrease in the band gap upon annealing in air[27]. In summary, the band gap of the films depends on the modelused in its determination, the experimental technique, and theannealing temperature and atmosphere. Our value of the directband gap of HfO2 thin films (5.63 eV) did not show any variationwith the annealing temperature, and was comparable to the valuesreported in the literature [2,41,51]. The indirect band gap showedconsiderable variation with annealing temperature, with the valueof the as-deposited films close to the value reported in Ref. [26].Thus, annealing in hydrogen did not affect the direct band gap, anddecreased the indirect band gap. The difference between the twoband gaps (0.17 eV) for the as-deposited films was comparable tothe theoretical value (0.25 eV).

3.4.2. Anti-reflection coatingsThe reflectance spectra of the films on silicon and fused silica

substrates are shown in Fig. 8. The reflectance spectra of the filmson fused silica (Fig. 8a) were typical of transparent high refractiveindex dielectric films on a transparent substrate in that the minimaof the reflectance spectra of the films matched that of the substrate.The reflectance spectra of the films on silicon substrates showed amuch lower reflectance than that of the substrate, indicating theirpotential use as anti-reflection (AR) coatings on silicon. For a thinfilm to function as an AR coating, its refractive index (n) must berelated to the refractive index of the substrate (ns) as n = √

ns. Forfiber-optic communications, the wavelengths of interest are 1310and 1550 nm [5]. The corresponding refractive indices of silicon are3.505 (� = 1310 nm) and 3.476 (� = 1550 nm) [55]. Our refractiveindex values (calculated using the Cauchy dispersion formula withthe parameters given in Table 3) are compared to those of siliconin Table 5. The values indicate the suitability of the as-depositedfilms and those annealed at 400 ◦C as AR coatings on silicon. This isclearly shown in Fig. 8b, where the reflectance can approach zerofor these films for certain wavelengths.

3.5. Electrical properties

The resistivity was determined for films deposited on fused

silica substrates. As-deposited films showed a large resistivity of3.4 × 106 � m, indicative of insulating films. This large value ofresistivity can be understood in light of the wide band gap of HfO2and the stoichiometric nature of the as-deposited films. Annealingthe films in hydrogen at 400 ◦C resulted in a significant reduction of

Indirect

˛o × 105 cm−1 Eg (eV) R4.72 5.46 0.9993.43 5.12 0.9964.04 5.34 0.9914.09 5.30 0.988

522 M.F. Al-Kuhaili et al. / Materials Chemistry and Physics 126 (2011) 515–523

fnium

rib8o

ciAdasbeaoT6ctfislcs

4

cfi4at

Fig. 8. Normal-incidence reflectance spectra of as-deposited and annealed ha

esistivity by three orders of magnitude to 8.8 × 103 � m. Anneal-ng the films at 600 ◦C resulted in a further decrease of resistivityy an order of magnitude to 9.3 × 102 � m. The films annealed at00 ◦C had a resistivity of 8.5 × 102 � m, which was similar to thatf films annealed at 600 ◦C.

Highly disordered inter-granular regions provide sites for thehemi-sorption of oxygen, which segregates into these sites dur-ng deposition, and gives rise to surface acceptor (trap) states [56].nnealing in a reducing atmosphere (such as hydrogen) leads to aecrease of resistivity due to desorption of oxygen, resulting in thennihilation of oxygen acceptor states [57]. The resulting vacanttates will act as donor states lying at the bottom of the conductionand [58], and thus increase the electron concentration [59]. Thisxplains the decrease of resistivity when the films were annealedt 400 ◦C, and is further supported by the less stoichiometric naturef the films annealed at 400 ◦C compared to the as-deposited films.he further decrease of resistivity when the films were annealed at00 ◦C can be understood on the basis of crystallinity. Improvedrystallinity was shown to reduce resistivity, due to the reduc-ion of charge carrier scattering at the grain boundaries [60]. Thelms annealed at 600 ◦C or 800 ◦C were polycrystalline and of theame chemical composition. Thus, their resistivities were simi-ar. Because of their reduced resistivity, stable microstructure, andhemical composition, these films are potential candidates for gasensor applications.

. Conclusions

Hafnium oxide thin films were deposited on unheated sili-

on and fused silica substrates using electron beam e-beam. Thelms were post-annealed in hydrogen in the temperature range00–800 ◦C. XRD results indicated that the as-deposited films weremorphous. Annealing led to the crystallization of the films inhe monoclinic phase. The onset temperature for crystallization

oxide thin films: (a) on fused silica substrates, and (b) on silicon substrates.

was 400 ◦C for films deposited on silicon, and 600 ◦C for filmson fused silica. The morphology of the films showed a colum-nar microstructure. As-deposited films were nearly stoichiometric.However, the annealed films were sub-stoichiometric. Photolumi-nescence measurements revealed the presence of emission bandsin the ultraviolet and near-infrared ranges, both of which wererelated to oxygen defects. The films were transparent down toa wavelength of 250 nm, and had a direct band gap of 5.62 eV.Transparency and direct band gap were not affected by annealing.The films annealed at 400 ◦C had higher values of the refrac-tive index and extinction coefficient. This was mainly due to anincrease in the density of defects in these films. In general, therefractive index was lower than the bulk value, and this wasattributed to the columnar growth of e-beam evaporated films.The as-deposited and 400 ◦C-annealed films exhibited an excellentanti-reflection behavior on silicon. Finally, the electrical resistiv-ity of the films was decreased by several orders of magnitudeupon annealing. In summary, hydrogen annealing did not affectthe optical properties of the films. Rather, annealing improvedcrystallinity and electrical conductivity, at the expense of stoi-chiometry. In particular, the films annealed at 400 ◦C showed thebest performance from an application point of view, due to theirlower resistivity (compared to as-deposited films), higher refrac-tive index, wide band gap, and anti-reflection behavior. However,they may not be suitable for applications that require crystallinefilms.

Acknowledgements

The support to this work by the Physics Department of King FahdUniversity of Petroleum and Minerals is acknowledged. We wouldalso like to thank professor N. Tabet for his assistance with the XPSand AFM measurements.

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