Thin Solid Films 629 (2017) 97–102

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Oxygen diffusion in TiO2 films studied by electron and ionRutherford backscattering

G.G. Marmitta,*, S.K. Nandib, D.K. Venkatachalamb, R.G. Ellimanb, M. Vosb, c, P.L. GrandeaaInstituto de Física, Universidade Federal do Rio Grande do Sul, Porto Alegre, RS, BrazilbDepartment of Electronic Materials Engineering, Research School of Physics and Engineering, The Australian National University, Canberra, AustraliacAtomic and Molecular Physics Laboratories, Research School of Physics and Engineering, Australian National University, Canberra, Australia

A R T I C L E I N F O

Article history:Received 24 October 2016Received in revised form 23 February 2017Accepted 12 March 2017Available online 18 March 2017

Keywords:DiffusionRBSElectron-RBS

A B S T R A C T

The diffusivity of oxygen in thin, sputter-deposited TiO2 films, as can be used in RRAMs, is measured usingelectron and ion backscattering techniques. The as-grown sample consisted of two layers (Ti16O2 and Ti18O2)and was annealed between 500 ◦C and 900 ◦C. The depth profiles of 18O, as measured with both techniques,were similar. The extent of diffusion was much larger than expected from the literature data for O diffusionin single-crystal rutile, suggesting that defects in the sputter-deposited film play an essential role in thediffusion process.

© 2017 Elsevier B.V. All rights reserved.

1. Introduction

Oxygen diffusion in rutile single crystals has been investigatedextensively in the past [1–6]. TiO2 has many technological appli-cations but many of these exploit an amorphous or polycrystallinephase [7]. To our knowledge no data exist for oxygen diffusionunder these conditions. In particular, TiO2 is a prototypical memris-tive material [8], used for resistive random access memory (RRAM).The resistive switching is a consequence of the formation (anddestruction) of oxygen-deficient filaments in the TiO2 layer byelectric-field induced diffusion of oxygen [8]. Details of these pro-cesses are still poorly understood, but it is clear that O diffusion isan essential ingredient. Here we study O self-diffusion during theannealing of sputter-deposited thin films.

Different experimental techniques have been used in the litera-ture to determined the diffusivity and activation energy for oxygenself-diffusion in TiO2, which include mass spectrometry and nuclearreaction analysis by 18O labeling. They were successfully used forthe cases where the diffusion length is several hundreds of nm.For much shorter diffusion lengths a new technique was proposedrecently [9], using the difference in energy of keV electrons scatteredelastically from either 18O or 16O. This technique, electron Rutherfordbackscattering spectrometry (ERBS) has depth-sensitivity based on

* Corresponding author.E-mail address: gabriel.marmitt@ufrgs.br (G. Marmitt).

the electron inelastic mean free path (IMFP). Therefore electronsscattered from 18O (or 16O) atoms close to the surface are more likelyto contribute to the elastic peak than those scattered from O fur-ther below. In contrast ion scattering provides information about thedepth distribution based on the energy loss per unit length traveled.Here we compare the results from both approaches.

2. ERBS technique

Fast electrons moving through matter interact with nuclei (subse-quently referred to as atoms) and electrons. Elastic collision betweena fast electron and an atom transfers momentum q, and hence energyto the scattering atom. If the atom is initially at rest it acquires akinetic energy q2/2Mi, where Mi is the mass of atom i. However, ifthe impacted atom has a momentum p before the collision (e.g. dueto thermal vibrations), then the recoil energy transferred to this atomis Doppler broadened and given by

Eirec =(p + q)2

2Mi− p

2

2Mi=

q2

2Mi+

p • qMi

. (1)

The electron loses this amount of energy. If the target containsatoms with different masses Mi, then the elastic peak will split upinto several components with energies related to their mass. Underour experimental conditions almost all detected electrons have lin-ear trajectories with a single large-angle backscattering collisionoccurring at depth x. In a previous work about multiple scattering in

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98 G. Marmitt et al. / Thin Solid Films 629 (2017) 97–102

ERBS experiments, we have shown that such approximation allowsfor the determination of the stoichiometry of TiO2 films within a1–2% error margin [10]. For isotropic materials the width due tothermal motion of each component is given by

si =

√43

ĒirecĒikin. (2)

with Ēirec the mean recoil energy (q2/2Mi) and Ēikin the mean kinetic

energy of atom i [11].The electron may also scatter inelastically, i.e. create electronic

excitations but then the energy loss is generally so large that it willnot contribute to the elastic peak. It is commonly assumed that theelectron does not change its direction in inelastic events and hencethe shape of a trajectory is determined by elastic collisions. Theprobability that a trajectory contributes to the elastic peak is deter-mined by the likelihood that no inelastic scattering occurred. If I0 isthe number of electrons backscattered from each depth, then theircontribution to the elastic peak will vary as

I(x) = I0e−l(x)/k, (3)

where l(x) = x(1/cos(h1) + 1/cos(h2)) is the length traveled and kis the IMFP of electrons with energy E0 and h1,2 the incoming andoutgoing angles.

When a substrate is covered homogeneously by an overlayer thenthe ratio of the overlayer and substrate signal strength is a mea-sure of the overlayer thickness [12]. Similarly, when the materialhas a concentration depth profile 0i(x) for some element i, the peakintensity of element I as measured by ERBS is proportional to

Hi = s i∫ ∞

0I(x)0i(x)dx, (4)

where s i is the differential cross section for backscattering from ele-ment i. Because only the ratio of peak intensities is measured wecannot uniquely determine 0i(x), so further assumptions about thedepth profile are necessary. A typical assumption is that atomic dis-tributions can be described as diffusion profiles as calculated fromthe diffusion equation

∂0i

∂t= D∇20i, (5)

where D is the diffusion coefficient. The oxygen isotopes have sim-ilar chemical properties, thus the diffusion coefficient is consideredindependent of variations of only one isotope concentration. It hasalready been shown [9] that solutions with different diffusion coeffi-cients lead to different measured peak intensity ratios.

3. Experimental procedure

In order to study O diffusion in TiO2, two sets of samples wereprepared. Set A consisted of pieces of a Si wafer on which 60 nmof Ti16O2 was first deposited followed by 20 nm of Ti18O2. Set Bhad the isotope composition inverted: 60 nm of Ti18O2 followed by20 nm of Ti16O2. We also grew samples with a single, isotopically-pure TiO2 layer with 80 nm thickness. The reactive sputter deposi-tions were performed with a Ti metal target and oxygen containingambient, which was powered by a pulsed DC source at 5 kHz andpulse length of 0.4 ls. The system used was calibrated for a deposi-tion rate of 0.67 Å/s, which for our samples leads to total depositiontimes of about 1000 s. The samples’ temperature during depositionremained below 200 ◦C. The Ti18O2 layers were grown using 99.9%pure 18O2 gas during deposition.

Fig. 1. (a) NRP spectra obtained using the 151 keV resonance in the 18O(p,a)15Nreaction for both sets. The lines are simulated intensities based on the 18O2 depthprofile shown in (b).

Nuclear Reaction Profiling (NRP) was used to measure the 18Odepth profile using the 18O (p, a) 15N reaction which has a resonanceat 151 keV (Fig. 1). The resonance has a width of 100 eV correspond-ing to a depth resolution of ≈0.5 nm close to the surface [13]. Asshown in Fig. 1, the as deposited sample from set A has 18O onlyin the top layer before the thermal treatments. In contrast, set Bshows contamination of 15% 18O in the top Ti16O2 layer, probably asa consequence of diffusion occurring during sputtering deposition.

Experimentally it was verified that at room temperature the iso-tope composition at the surface did not change with time. In orderto induce a controlled diffusion, rapid thermal annealing (RTA) wasused. During this process oxygen may exchange between the sur-face and the atmosphere, even under a flow of inert gasses such asAr. To prevent this exchange a capping layer of Si3N4, was depositedby PECVD at about 300 ◦C. Ellipsometry indicated a thickness of thecapping layer of ≈33 nm, in agreement with ERBS measurements.

Samples from each set were exposed to temperatures between500 and 900 ◦C for 5 min in an Ar atmosphere. Some samples fromset B were annealed at fixed temperature of 650 ◦C for time peri-ods between 5 and 100 min. After annealing, the capping layer wasremoved by chemical etching with solution of 6:1 BHF (6 40% NH4F:1 49% HF), for duration between 3 and 9 min. Etching was stoppedwhen visual inspection made it clear that the capping layer wasremoved. Under high annealing temperatures Si3N4 changes its den-sity and becomes more resistant to chemical etching, hence longeretching times were required. The etching process is sufficientlyselective that TiO2 removal is not an issue, as was confirmed byRutherford Backscattering Spectrometry (RBS).

The ERBS system is described extensively elsewhere [14]. In short,electrons with 40 keV energy scatter over 135◦ and are energy-analyzed with a resolution of 0.3 eV. Different geometries can beused to change the depth probed. We used a perpendicular incoming(h1 = 0◦, h2 = 45◦) and glancing exit geometry (h1 = 35◦, h2 = 80◦)for bulk and surface sensitive measurements, respectively, where h1and h2 are the incidence and exit angle with respect to the surfacenormal. A similar approach was previously used to determine theoxygen concentration profile for Hf18O2/Hf16O2 structures [9].

RBS used 2 MeV He+ ions impinging along the surface normaland a scattered angle of 110◦ (i.e. detected at 70◦ to the sample nor-mal). The O peak differentiation is only straight-forward for set Asamples, where the energy loss in the first Ti18O2 layer helps sepa-rate the 16O and 18O peaks. Moreover, the O peaks are superimposedon the signal from the (single crystal) Si substrate. The samples were

G. Marmitt et al. / Thin Solid Films 629 (2017) 97–102 99

Fig. 2. The top panel shows ERBS spectra of the as deposited samples from both sets.The arrows indicate the mean, calculated energy loss for 16O and 18O. The spectrataken at glancing out geometry were normalized to equal Ti peak area. The lower panelshows that after thermal annealing, the O peaks shifts, indicating that some diffusionoccurred.

rotated to minimize the influence of channeling, and the Si substratesignal was subtracted to proceed with the analysis. Quantitativeanalysis of the O concentration profile was still difficult, so RBS spec-tra were analyzed based on the oxygen profiles as derived fromERBS.

4. Results and discussion

ERBS measurements of the as deposited samples in the glanc-ing geometry are shown in Fig. 2. Under this condition, only thetop layer of either Ti16O2 and Ti18O2 contributes to the spectrum.Indeed the observed O peak position depends on the isotope mass.The extra width of the O peak in the ERBS spectra due to Dopplerbroadening (see Eqs. (1), (2)) is resolved in the experiment. The peakwidth used for the fit was calculated from Eq. (2) and the samewidth was used before and after annealing. Then, we determined theoxide film composition from samples with thick, isotopically pure,as deposited titanium oxide layers. The sputter deposited oxide filmshad the expected stoichiometry of TiO2. Subsequently, we fitted theO peaks of the sandwich structures using two components (16O and18O) with the parameters as determined from the above mentionedsamples. Changes in the shape of spectra could be entirely explainedby variations in the intensity ratio of the 16O and 18O signal.

Concentration profiles were calculated by solving Fick’s secondlaw to track the 18O movement, as shown in Fig. 3a. Because thestoichiometry of the oxide is TiO2 the sum of both isotopes alwaysequal 2, hence we may write it as Ti18O16x O2−x and from one isotope

Fig. 3. In (a), diffusion profiles of samples from set A, obtained by solving the Fick’ssecond law for different values of Dt. The ERBS peak intensity is given by the integra-tion of the concentration profile attenuated by the length traveled by the electrons.The resultant normalized peak intensities for 16O and 18O are shown in (b). Thus, ERBSspectra of such samples show contributions of both 18O and 16O.

concentration profile calculate the other. These calculations assumedthat the Si substrate and the Si3N4 capping layer were impenetra-ble barriers for the oxygen. The final diffusion profile does dependon the initial concentration of 18O in each layer. The result of theNRP measurements were used to determine the starting point of theFick-diffusion profile calculation.

Once the 16O and 18O concentration profiles are given, theexpected ERBS peak intensities for each isotope can be calculated byEq. (4). Due to the beam attenuation, ERBS is more sensitive to thesample surface. As 18O diffuses deeper, its elastic peak will becomeless intense. Conversely, a higher 16O concentration is then foundnear the surface and its elastic peak becomes more intense. Thecomplete ERBS spectrum has a superposition of elastic peaks of thetwo isotopes in slightly different energy positions and the relativecontribution of both isotopes can be determined by fitting the spec-trum. Analyzing the ratio of the 16O and 18O elastic peak intensitiesthen provides a direct measure of the extent of diffusion. The resul-tant peak intensities for selected concentration profiles are shown inFig. 3b.

From this analysis it is possible to determine the Dt values thatbest match the measured profiles for both ERBS geometries, asdiscussed before [12]. Extracted diffusion lengths, (sqrt(Dt)), and dif-fusion coefficients, D, are plotted in Fig. 4, as a function of annealingtemperature. The diffusion coefficients are consistent for each of thetwo sample sets and are observed to satisfy an Arrhenius relationwith an activation energy EA ≈ 1.05 eV.

Set A was also measured by RBS to cross check the diffusion coef-ficients obtained by ERBS. Fig. 5 shows the RBS spectra for differentannealing temperatures, together with simulations obtained using

100 G. Marmitt et al. / Thin Solid Films 629 (2017) 97–102

Fig. 4. Arrhenius plot for both sets. From the fitted line one can obtain the activationenergy for the oxygen diffusion in TiO2.

the PowerMEIS software [15]. In a similar way to the ERBS analy-sis, the sample was described as a series of films of Ti18O162−xOx withx obtained from the Fick’s diffusion profile. Quantitative analysis ofthe O concentration profile was still difficult, so the diffusion profilesused in these RBS simulations are those that described best the ERBSspectra. The resulting description of the RBS experiment is very good.The ERBS and RBS approaches thus describe the diffusion processconsistently.

Even for single crystal rutile there is a large spread in reportedactivation energies [1-3,5,16,17], varying from 0.35 to 2.42 eV. Thework of Arita et al. [4] follows a similar approach, where longer ther-mal annealings with t > 20 h and T > 1000 ◦C were conductedon TiO2 samples. For the diffusion coefficients measured from set Bsamples, agreement is seen for similar temperatures — between 800and 1000 ◦C. However, the activation energies measured are greatlydifferent (dotted lines’ slope on Fig. 6, top pane). The experimentsfor crystalline rutile at T > 1000 ◦C produced an activation energyof about 3 eV, which is much larger than the one obtained here forthin films of TiO2 at lower temperatures. EA = 3 eV is consistentwith the diffusion of charged vacancies (V2+O ) as described by ab-initio calculations from Iddir et al. [18], which estimate the vacancyforming energy and the migration energy barrier. The former is esti-mated at 2 eV for the given conditions, and the latter depends on thecrystalline configuration, with energy barriers in different migration

Fig. 5. The top panel shows the RBS spectrum (black points) of set A, as depositedsample, the red lines are simulations of the TiO2 contribution. The difference betweenthe experimental and simulated spectra (i.e. the Si contribution) is in blue. The bot-tom panel shows the oxygen region in detail (Si contribution subtracted) for differentannealing temperatures.

directions between 0.69 and 1.10 eV. These migration barriers val-ues are consistent with the diffusion activation energy obtained here,which suggest that diffusion without thermal vacancy generationcan describe our 5 min experiments but not experiments using muchlonger annealing time.

In order to investigate the discrepancy observed in the activa-tion energy, further experiments were performed to determine theregularity of the diffusion. This is rather important, since previousmeasurements had annealing duration of a few hours instead of min-utes. By keeping the annealing temperature constant at 650 ◦C andvarying the annealing time between 5 and 100 min, diffusion lengthswere measured on samples from set B. As can be observed in Fig. 6,the time dependence of the diffusion length does not follow theusual square root law expected for regular diffusion. On the con-trary, two diffusion regimes may be identified. On a time scale ofminutes a rapid diffusion occurs, compatible with the low activationenergy measured in our samples. Then a much slower diffusion pro-cess takes place, which can be described by a larger EA value of ≈3eV from ref. [4].

A simplistic way to describe these experiments assumes that thesputter-deposited films have initially a high defect density. Anneal-ing at a certain temperature will lead to their break-up resulting indiffusion of point defects. The break up of defects will be accompa-nied by diffusion of interstitials and vacancies. Excess point defectconcentrations will increase the oxygen diffusivity, as is well knownfrom radiation enhanced diffusion experiments. After a certain timeat a given temperature the concentration of point defects will drop,

G. Marmitt et al. / Thin Solid Films 629 (2017) 97–102 101

Fig. 6. On the top pane, the oxygen diffusion coefficients measured after annealing for5 min are compared to values obtained by Arita et al. [4] for annealing with t > 20 h.On the bottom pane, extrapolation of the diffusion length is compared to measure-ments of set B samples, annealed at fixed temperature but different annealing times.The solid line corresponds to the time extrapolation using EA = 1.05 eV whereas thedashed-line represents an extrapolation using EA = 3 eV and an initial diffusion lengthn according to

√Dt + n2.

as only the more stable defects remain and the diffusivity will bereduced. Increasing the annealing temperature causes more defectsto break up and introduces again excess point defects, increasing thediffusivity temporarily.

5. Conclusion

For the oxygen self-diffusion measurements, the electron- andion-Rutherford backscattering techniques were employed. Partic-ularly, ERBS was shown to have sensitivity and depth resolutionrequired to determine the oxygen diffusion on a length scale ofseveral 10s of nm – the length scale on which resistive switchingoccurs. This technique also has the advantage that it avoids possiblebeam damage induced diffusion during measurements. The profilesobtained by this procedure were then checked by conventional RBSmeasurements with 2 MeV ions. Due to the samples structure, spec-tra with good peak separation of the 16 and 18 isotopes were onlyfound in one of the sample sets. Simulations, based on the profilesmeasured by ERBS, agree well with those measured by RBS.

Sputtered deposited samples containing Ti18O2 and Ti16O2 filmswere also measured with NRP, which is sensitive only to the 18O iso-tope and is not able to probe for the total oxide stoichiometry. There-fore, the technique is useful to evaluate samples’ initial conditions,however is not sufficient for a complete diffusion analysis.

The observed diffusion was found to be irregular, dependingstrongly on the annealing history, not just the temperature. Hencethe interpretation of the experiments in terms of activation energiesis questionable. However, if one interprets our 5 min isochronalannealing experiments in terms of an activation energy, one derivesa value compatible with calculated oxygen vacancy migration bar-riers in TiO2. Also, the sputtering deposition of TiO2 used in oursamples is known to produce films with large number of defects, con-dition which can enhance the initial self-diffusion rate. For longerannealing times the activation energy reported on literature is muchhigher than the 1.05 eV measured in the present study, indicating aslower diffusion process in which oxygen vacancy forming is neces-sary. Our isothermal experiments are also consistent with a transientfast diffusion process followed by a slower regime. The use of TiO2 asoxide layer in RRAM devices is desired precisely because of this fastdiffusion observed in our sputtered deposited films, even at lowerannealing temperatures.

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Oxygen diffusion in TiO2 films studied by electron and ion Rutherford backscattering1. Introduction2. ERBS technique3. Experimental procedure4. Results and discussion5. ConclusionReferences