Reactive magnetron sputtering of Nb-doped TiO2 films: Relationships between structure, composition and electrical properties

    Reactive magnetron sputtering of Nb-doped TiO 2 films: between structure, composition and electrical properties

Relationships

Stefan Seeger, Klaus Ellmer, Michael Weise, Daniela Gogova, Daniel Abou-Ras, Rainald Mientus PII: DOI: Reference:

S0040-6090(15)01211-0 doi: 10.1016/j.tsf.2015.11.058 TSF 34840

To appear in:

Thin Solid Films

Received date: Revised date: Accepted date:

30 May 2015 18 November 2015 21 November 2015

Please cite this article as: Stefan Seeger, Klaus Ellmer, Michael Weise, Daniela Gogova, Daniel Abou-Ras, Rainald Mientus, Reactive magnetron sputtering of Nb-doped TiO2 films: Relationships between structure, composition and electrical properties, Thin Solid Films (2015), doi: 10.1016/j.tsf.2015.11.058

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ACCEPTED MANUSCRIPT Reactive magnetron sputtering of Nb-doped TiO2 films: relationships between structure, composition and electrical properties

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Stefan Seeger1, Klaus Ellmer2, Michael Weise1, Daniela Gogova3,4, Daniel Abou-Ras2, Rainald Mientus1 1

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Optotransmitter-Umweltschutz-Technologie e.V., Köpenicker Str. 325, 12555 Berlin, Germany

2

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Helmholtz-Zentrum Berlin für Materialien und Energie, Hahn-Meitner-Platz 1, 14109 Berlin, Germany

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Central Lab of Solar Energy and New Energy Sources at the Bulg. Acad. Sci., Blvd. Tzarigradsko shose 72, Sofia, Bulgaria 4

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Leibniz Institute for Crystal Growth, Max-Born-Str.2, 12489 Berlin, Germany

Abstract

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Niobium-doped TiO2 films as highly transparent conducting oxides for electrical contacts were investigated. As-deposited films were amorphous and exhibited high

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resistivities ranging from 10 to 105 cm. A slight oxygen deficiency in as-deposited

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films was essential to gain low resistivities (10-3 cm) and low optical absorption

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coefficients (α550nm < 2x103 cm-1) in the annealed films. Therefore, we controlled the oxygen stoichiometry during the film deposition by adjusting the magnetron discharge voltage, while the oxygen gas flow was kept constant. The Hall mobility of

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degenerately doped films (electron concentration > 1020 cm-3) increased with decreasing substrate temperature owing to metal-like phonon scattering in these samples.

Corresponding author: [email protected] 1

ACCEPTED MANUSCRIPT 1. Introduction In recent years, highly niobium-doped titanium oxide (TiO2:Nb) films have been

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studied widely as an alternative to the rather expensive, transparent conductive material

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In2O3:Sn [1-8]. Many efforts have been concentrated on deposition of polycrystalline

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niobium-doped anatase (TNO) films on glass and polyimide substrates by sputtering processes, i.e., flexible electronics [9-19]. Although sputtered polycrystalline TNO films exhibit high optical transparency of larger than 80 % at 550 nm and reach

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resistivities as low as ρ < 10-3 Ωcm, their carrier mobilities (µe < 10 cm2V-1s-1) [9, 10, 16, 17, 20, 21] are still lower than that of conventional TCOs, as, e.g., µe > 50 cm2V-1s-1

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of In2O3:Sn or ZnO:Al. Therefore, a route to reduce the resistivity below 10-3 Ωcm is to

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increase the carrier mobility in polycrystalline TNO films [20].

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Various strategies have been used to increase the mobility µe,

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(i) by controlling the crystalline phases and their preferential orientations in polycrystalline films [4],

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(ii) by increasing the average grain sizes [22], and

(iii) by exact adjustment of the [Nb]/[Ti] ratio and the [O]/[metal] ratio [6, 20].

The crystallographic orientation of sputtered polycrystalline TNO can be controlled by using single-crystalline substrates, e.g., (100) LaAl2O3 or (100) SrTiO3 [20, 23]. Epitaxial TNO films are characterized by low resistivities (< 10-3 Ωcm) and high carrier mobilities (> 20 cm²V-1s-1) at room temperature, although their average grain sizes were only about 30 nm [20, 23]. In contrast, polycrystalline TNO films with large average grain sizes (>20 µm) were obtained, when amorphous TNO films were deposited on 2

ACCEPTED MANUSCRIPT unheated substrates and subsequently annealed at 400 °C in vacuum or in a reducing atmosphere [9, 20, 22, 24]. In spite of the large average grain sizes in these highly-

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doped polycrystalline TNO films, their mobilities at room temperature (µe < 10 cm2V-1

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s-1) is smaller than those of epitaxial TNO films.

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Yamada et al. [20] investigated the direct growth of polycrystalline TNO films on LaAlO3 and glass substrates by reactive sputtering from a compound target (Ti0.94 Nb0.06

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O1.95). The sputtering of a thin TNO seed layer (<50 nm) in an oxidizing atmosphere on glass substrates at low temperature (250 °C) prior to the deposition of polycrystalline

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TNO films at 400 °C induces the growth of the anatase phase. In spite of the preferential growth of the anatase phase, the considerable size of the crystallites (1 µm) and the high

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carrier concentration (about 1021 cm-3), the resistivity at room temperature of these

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directly sputtered polycrystalline TNO films was still too high (>1 10-3 Ωcm) due to the

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rather low mobility (<5 cm2V-1s-1). It was reported that the precise control of the chemical composition of TNO films, e.g., the niobium content and the [O]/[metal] ratio,

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eventually is important to reduce the resistivity [20, 23].

Furubayashi et al. [23] found that the electrical transport mechanism depends strongly on the niobium concentration (x) in Ti1- x Nb x O2-y films. At low niobium concentrations x < 0.01 (resulting in ne < 2x1020 cm-3), scattering at grain boundaries determines the carrier mobility, while at x > 0.1 (ne >2x1021 cm-3) scattering at ionized and neutral impurities are dominant in epitaxial TNO films. Yamada et al. [20] sputtered epitaxial films on (100) LaAlO3 substrates at temperatures of about 400 °C from a ceramic target (Ti1-xO2-yNbx) in a mixture of argon and oxygen at various oxygen contents. The lowest resistivity (3.6x10-4 Ω cm) and highest mobility (13 cm2 V-1 s-1) was reached at an 3

ACCEPTED MANUSCRIPT oxygen fraction of 0.35 % in the sputter gas. Furthermore, these authors have found that sputtering with lower (<0.15%) and higher (>0.75%) oxygen fraction results in TNO

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films with a much higher resistivities (>1x10-3 Ω cm). Sputtering at low oxygen partial

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pressures result in films with low mobilities, which can be attributed to the

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accompanied phase transformation from anatase to rutile [20]. The high resistivity of oxygen-rich films was attributed to the formation of acceptor-like defects, such as oxygen interstitials (Oi) and/or titanium vacancies (VTi), which compensate the

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electrons from the niobium dopant [20].

xNbxO2

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In the present work, we aimed at improving the electron mobility of polycrystalline Ti1films. We prepared these films with x in the range 0.03 ≤ x≤ 0.06 [25] by

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reactive sputtering from a Ti-Nb alloy (3.2 at% Nb) target in an argon-oxygen

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atmosphere on glass substrates and subsequently annealed them in vacuum at 450 °C,

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since by this approach (i) the doping concentrations optimum for both, high transparency and low resistivity, was obtained, (ii) the annealing of amorphous films leads to polycrystalline film with average grain sizes (d >20 µm) [9, 20, 22, 24], and

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(iii) we were able to vary the integral film composition. Furthermore, reactive magnetron sputtering from metallic targets allows for higher deposition rates compared with sputtering from titanium oxide targets, because of their high thermal conductivities and mechanical stability [16].

The deposition rate and the film composition depend nonlinearly on the flow of the reactive gas, which can cause instabilities in the sputtering process, i.e., leading to a well-known hysteresis behavior [26]. Therefore, a process control is essential for a reproducible deposition of compound films [16, 27]. Moreover, the composition of the 4

ACCEPTED MANUSCRIPT as-deposited TNO films determines the structural and electrical properties after thermal annealing at 450 °C, particularly, the formation of the anatase phase and the electron

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mobility. The oxygen-to-metal ratio in the amorphous TNO films was controlled by

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adjusting the discharge voltage at a fixed oxygen gas flow. Therefore, the TNO films

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were deposited at different discharge voltages and investigated by Hall and conductivity measurements, Raman spectroscopy, Rutherford backscattering spectrometry (RBS) as well as by electron back scattering diffraction (EBSD). Polycrystalline TNO films

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(anatase) with a low resistivity (ρ =1.1·10-3 Ω cm), a carrier concentration of ne =

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6.5·1020 cm-3 and an electron mobility μe = 8 cm² V-1 s-1 combined with high optical transmittance of about T > 90 % in the VIS-IR spectrum were obtained.

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2. Experimental details

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Deposition series of amorphous TNO film were prepared by reactive dc magnetron

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sputtering using a circular Ti:Nb (3.2 at%) target (75 mm in diameter) in an argon/oxygen atmosphere at total sputtering pressures of 1 Pa (FAr = 30 sccm) at various oxygen flows (FO2 1.2 sccm to 1.5 sccm). The experiments were performed in a

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commercial sputtering system (Leybold Z400) with a base pressure of better than 10-3 Pa and equipped with a load-lock. The target-to-substrate distance was about 60 mm and the substrates were not intentionally heated during the film growth. The gas flows were controlled by mass flow controllers and the total pressure was measured with a gas independent capacitance vacuummeter (Baratron). The thin amorphous TNO films (thickness about 100 nm measured with profilometer) were deposited onto 0.5 mm thin borosilicate glass substrates (Schott, D263T) and glassy carbon substrates. Afterwards, these TNO films were annealed in vacuum at a base pressure better than 1·10-4 Pa at a temperature of about 450 °C for 10 min. The heating ramp was about 50 K/min. 5

ACCEPTED MANUSCRIPT The chemical composition in the amorphous TNO films deposited on glassy carbon substrates was measured by Rutherford backscattering spectroscopy (RBS) with 1.7

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MeV He ions at a scattering angle of 170°. The RBS spectra were analyzed using the

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software SIMNRA [28]. The resistivity of the as-deposited TNO films was measured by

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a four-point probe method. The structural properties of the films were analyzed using X-ray diffraction (XRD) (Cu-anode) with a silicon stripe detector (D2 Phaser with a Lynxeye detector, Bruker AXS) in Bragg-Brentano geometry. The diffraction peaks

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were fitted using Lorentzian peak-shaped double peaks, i.e., yielding the position and

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width of the two CuKα peaks. Additionally, Raman spectroscopy measurements were carried out to determine the crystallinity of TNO films. For this purpose a Horiba JobinYvon LabRAM HR800 high-resolution confocal µ-Raman system was employed at

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room temperature. The 488 nm line of an argon-ion laser was used for excitation. The carrier transport properties of the thin polycrystalline TNO films, i.e., the resistivity and

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Hall coefficient, were determined using the van der Pauw method. Ohmic contacts were deposited at the corners of the annealed polycrystalline TNO films which were prepared

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onto square glass substrates (10x10 mm2).

The microstructure of the polycrystalline TNO films was analyzed by electron back scattering diffraction (EBSD). These measurements were performed using a Zeiss Ultra Plus scanning electron microscope equipped with a NordlysNano EBSD detector and the Oxford Instruments AZtec acquisition and evaluation software. The applied voltage, beam current, and step size were 15 kV, 6 nA, and 200 nm.

3. Results and discussion 3.1 Sputter deposition 6

ACCEPTED MANUSCRIPT At a constant dc power of a magnetron discharge with a Ti target a distinct increase of the plasma impedance, i.e., the discharge voltage, characterizes the transition from the

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metallic to the oxidic mode, caused by a significant change of the secondary-electron

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emission coefficient [29]. In Fig. 1, the black open diamonds show the discharge

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voltage in dependence on the oxygen flows for constant dc sputtering power (100 W) at 1 Pa total pressure. The two vertical dotted lines mark the discharge voltages in the metallic mode at 235 V and the oxidic mode at 323 V. The common way to stabilize a

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reactive magnetron discharge is to control the oxygen flow with a feedback system

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while monitoring a characteristic plasma emission line [27] or the plasma impedance [16]. In the present work, we used the sputtering power to stabilize the reactive sputter process while monitoring discharge voltage at various fixed oxygen flows. To study the

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sensitivity of this feedback system, we deposited TNO films at several discharge voltages between 280 V and 300 V, while the deposition process was stabilized by fine

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tuning the sputtering power (± 2 W). Since the discharge voltage is a measure of the oxidation of the surface of the erosion groove, it was possible to conduct the depositions

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at different working points.

Fig. 1 shows the deposition rates (see the closed symbols) for reactive gas sputtering of TNO films at four oxygen flows in dependence on the discharge voltage (see the open circles). Additionally, the two magenta squares mark the deposition rates for sputtering in the oxidic mode ≈ 2 nm min-1as well as in the metallic mode ≈27 nm min-1. Only in a narrow region area of discharge voltages, the annealed TNO films exhibited low resistivities. The grey vertical bar refers to the corresponding deposition rates of about ≈20 nm/min, i.e., about 80 % of the metal deposition rate.

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3.2 Annealing

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The adjustment of the discharge voltage at a fixed oxygen flow strongly affected the

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resistivity of the amorphous TNO films and allowed us to tune the resistivity over a

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wide range (several orders of magnitude). Since the absolute values of the discharge voltage shifted slightly during different experimental series we normalized the discharge voltage Vdis(norm) for each deposition series to the value at which the

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annealed films exhibited the lowest resistivity. Vdis(norm) smaller than 1 refers to

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sputtering under oxygen deficiency (metallic mode), and Vdis(norm) larger than 1 stands for slight oxygen excess (oxidic mode). Fig. 2 shows the resistivity of the TNO films before (dashed lines) as well as after the thermal annealing (solid lines) in dependence

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on Vdis(norm). The resistivities of the as-deposited films increased monotonously with higher discharge voltage. In case the target surface captured more oxygen the sputtering

increased.

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mode changed from metallic to oxidic and the resistivity of the as deposited TNO films

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However, during sputtering from the oxidized target surface, the plasma impedance recorded during the deposition process was less reliable than in the metallic mode. Hence, the resistivities of the amorphous TNO films were spread across a large range. After thermal annealing the resistivity decreased considerably (see the solid lines in Fig. 2) and exhibited a local minimum between the metallic and the oxidic sputtering mode. That is, in the case the as-deposited films showed a resistivity of about 50 to 200 Ω cm we obtained the lowest resistivities (of down to 1x10-3 Ω cm) after thermal annealing.

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ACCEPTED MANUSCRIPT Fig. 3 shows the spectral transmittance and the reflectance of a TNO film sputtered at Vdis(norm) = 1 before and after the annealing at 450 °C in vacuum. The thickness of the

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as-deposited film was 130 nm and its resistivity about 100 Ωcm. The transmittance of

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the film in the visible spectrum increased significantly after annealing (red solid lines).

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The inset in Fig. 3 depicts the absorption coefficients at 550 nm (α550nm) of TNO films in dependence on Vdis(norm). The black open squares mark the values for the asdeposited films, and solid red squares the α550nm of the annealed films, respectively. The

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film deposition at Vdis(norm) ≤0.99 led to high optical absorption coefficients, 5x103

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cm-1 < α < 3x104 cm-1,in accordance with the low resistivity of these as-deposited films (see Fig. 2).

With increasing Vdis(norm), i.e., capturing of more oxygen onto the target surface, α550nm

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of the as-deposited films decreased, while the resistivity increased. At Vdis(norm) >1.007 the as-deposited films were oxide-like, i.e., they exhibited low absorption

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coefficients α < 1 102 cm-1 and high resistivities (>1.6 103 Ωcm). After the annealing, only the TNO films sputtered around Vdis(norm) ≈1 (see shaded region in Fig. 3)

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exhibited a strong decrease of the absorption coefficient (α550 nm < 2000 cm-1) and, at the same time, these films showed low resistivities (ρ ≈ 1.3 10-3 Ωcm). The optical absorption of the films sputtered at Vdis(norm) <0.98 only slightly decreased after annealing, while the absorption coefficient of the sputtered films at Vdis(norm) >1.007 increased.

3.3 Film compositions The chemical composition of the as-deposited TNO films was investigated by RBS. An RBS spectrum is shown exemplarily in Fig. 4a. In order to achieve a high accuracy for 9

ACCEPTED MANUSCRIPT the atomic concentrations the TNO films were deposited onto glassy carbon substrates which do not produce a significant background signal for the elements in the film (O,

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Ti, and Nb). Figs. 5a-c show the atomic concentrations of titanium, niobium and oxygen

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in dependence on Vdis. The accuracy of the atomic concentration is mostly determined

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by the count numbers in the peaks of the different element leading to statistical accuracies of better than 1.6 % (Nb), 1 % (O), and 0.5 % (Ti), depicted as error bars. As expected, the change of the sputtering mode from metallic to oxidic is accompanied by

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an increase of the oxygen concentration by about 0.7 % (relative) which is resembled by

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the corresponding increase of the resistivities. The opposite trend is observed for the titanium concentration which is decreasing for more oxidic deposition condition by about 14.7 % (relatively) because of reduced titanium sputtering yield drops in case the

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target surface becomes oxidized. Just the reverse is observed for the concentration of niobium which increases significantly by about 10 % (relative) in case the sputtering

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mode changed to oxidic. Though, the concentration of niobium, intently added as a dopant, increased by 10 %, the resistivity of these polycrystalline TNO films was still

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quite high (≈5x10-2 Ωcm). This could be caused by the formation of acceptor-like defects, such as oxygen interstitials (Oi) and/or titanium vacancies (VTi), which compensate the electrons of the niobium dopant [20, 30]. Furthermore, the total amount of niobium is less (2.9 at%) than that in the metallic target (3.2 at%), which could be to the target manufacturing process (see Fig. 4b).

The RBS analysis confirmed that Vdis(norm) < 1 corresponds to TNO films with a slight oxygen deficiency (metallic mode) whereas Vdis(norm) > 1 indicate a weak oxygen excess (oxidized target) (see Fig. 4c). This means, our films are nearly stoichiometric within the accuracy limits of the RBS analysis. On the other hand, a quite significant 10

ACCEPTED MANUSCRIPT oxygen excess was also reported by other authors (see ref. [31, 32]). Since their composition data were also obtained from RBS analysis, the correct consideration of the

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background for oxygen signals is critically for a proper calculation of the oxygen

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content of the films.

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3.4 Structural properties

Fig. 6 depicts the XRD patterns of the annealed TNO films. The dashed vertical lines

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mark the diffraction peaks of anatase according to the reference powder diffraction file

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(JCPDS file No. 021-1272) for Cu-Kα1 X-ray radiation. The detected diffraction peaks of our films matched well with the reference data in case the TNO films were sputtered at Vdis(norm) > 1. The more the sputtering mode was adjusted towards the metallic

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mode, the more the diffraction peaks of 101, 200 and 211 moved towards lower diffraction angles. Only the position of the 004 diffraction peak did not vary. The films

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prepared at Vdis(norm) << 1 were still amorphous after annealing.

To take a closer look at the structural properties of these films, the pronounced 101

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diffraction peaks were analyzed with respect to their peak positions and peak widths, where the latter are related to the sizes of coherently diffracting domains. Fig. 7a shows the 101 peak position in dependence on Vdis(norm). The horizontal line marks the position of the 101 diffraction peak according to the anatase reference powder diffraction file. The 101 peak shifts to a lower diffraction angles indicating an expansion of this axis. Furubayashi et al. [6, 23] attributed the lattice expansion in TNO films, i.e., the larger length of the a- and c-axes of the tetragonal unit cell, to the niobium incorporation into the lattice, because the Nb5+ ionic radius (78 pm) is slightly larger than that of Ti4+ (74.5 pm) [16]. Since the niobium concentration of our sputtered 11

ACCEPTED MANUSCRIPT TNO films slightly increase with higher Vdis (see the RBS data in Figs. 4), the shifts of the diffraction peaks at Vdis (norm)<1 might be ascribed to changes in the structural

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properties, i.e., to preferred orientation of the crystallites, and lattice strain. At

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Vdis(norm) ≤ 0.993, the diffraction patterns revealed insufficient crystallization after

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annealing. These films exhibited increased resistivities and higher optical absorption coefficient, see Figs 2 and 3.

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We used the Scherrer equation [33] to determine the coherently diffracting domain sizes (CDS) of the annealed TNO films. Fig. 7b shows the calculated CDS of these films and

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the profilometric film thicknesses in dependence on Vdis(norm). The thickness of the annealed films is plotted over Vdis(norm) as blue solid line. The annealed TNO films

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with lowest resistivity, i.e., Vdis(norm) ≈ 1, exhibited coherent diffracting domain sizes

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which were comparable to the film thicknesses. At Vdis(norm) < 1 the films grew even

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thicker due to the higher deposition rate (see Fig. 1), but these annealed films exhibited smaller average grain sizes.

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Additionally, we used Raman spectroscopy to investigate as-deposited as well as annealed TNO films. Fig. 8a shows the Raman spectra of a deposition series, which was performed at various discharge voltages, i.e., Vdis(norm) at FO2 =1.4 sccm and 1 Pa. The as-deposited samples of both deposition series exhibited no Raman bands. Therefore, they were considered as amorphous which was in agreement with the X-ray diffraction analysis. The annealed TNO films exhibited five Raman bands which were assigned to the Raman modes of the anatase phase. The two Raman modes Eg (at 144 cm-1 and 197) and B1g (399 cm-1) are the O-Ti-O bending type vibrations and the other two modes, A1g* (516 cm-1) and Eg (639 cm-1), are the Ti-O bond stretching type vibrations [34]. 12

ACCEPTED MANUSCRIPT Since Raman bands of rutile [35] were not detected, we assigned the Raman band at 144 cm-1 solely to the Eg mode of anatase. For comparison, a Raman spectrum of bulk

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anatase (R060277) [36] is shown as a dotted blue line (see Figs. 8a and b). The Raman

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bands of the annealed TNO films show Lorentzian peak shapes comparable with the

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Raman spectrum of the anatase reference.

The band frequency and peak shape can be strongly influenced by residual stress [13],

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crystallite size [37], changes of composition [38-40], and temperature [41]. With the decrease of the Vdis(norm), i.e., lower niobium and oxygen concentration in the

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sputtered TNO films (see Figs. 4d and f), the two Eg modes (144 cm-1 and 197 cm-1) showed a blue shift compared with the anatase reference, while the B1g modes (399 cm) moved towards lower frequencies independently of the sputtering condition. In

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1

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contrast, the Raman bands of the A1g* mode at 516 cm-1 which represents the Ti-O

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stretching type vibrations [34] exhibited no frequency shift.

Lü et al. [40] investigated the pressure-induced phase transition of TiO2 nanoparticles

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by in situ Raman spectroscopy and observe a blue shift of the anatase Eg modes during compression, while the frequency band B1g shifts towards lower energy. They attribute the blue shift of the Eg mode to a shortening of the O-Ti-O bond, i.e., a larger displacement of the oxygen atoms than that of titanium atoms during the compression of the TiO6-octahedra. Nb-doping of the anatase nano-particle increases the frequency of the Eg mode as well as their band width [40]. Castro et al. [13] annealed amorphous TNO films in vacuum and under hydrogen atmosphere for 1 hour at 500 °C and observe a blue shift of the Eg mode in the Raman spectra. They presume that the Nb-dopants generate stress in the TiO2 lattice due to the different ionic radii of Ti4+ and Nb5+. 13

ACCEPTED MANUSCRIPT However, since the size of the crystallites, i.e., CDS (see Fig. 7a), of our TNO films were comparable to the film thickness (~ 100 nm), we excluded that particle size effects

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[42] induced the blue shift of the Raman bands. Moreover, the Raman studies of our

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samples were realized at the same laser power hence we assumed similar temperatures

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for all measurements.

Instead, we correlated the Raman band frequencies and their band width with the carrier

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concentration, see Fig. 9. While the frequencies of the Eg modes significantly rose from 143 cm-1 to 151 cm-1 with increasing carrier concentration, their FWHM did not vary.

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We attributed the blue shift of the Eg mode to the distortion of the anatase structure due to the incorporation of niobium. In other words, if more niobium is activated, i.e.

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niobium replaced titanium at lesser oxygen content, the higher the distortion of the

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anatase structure. Since Eg frequency bands showed neither a broadening nor distinct

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new features, we tentatively excluded phonon-confinement effects as well as plasmonphonon interactions as reason for the observed blue shift of the Eg modes.

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3.5 Electrical transport properties

Fig. 10 depicts the carrier concentrations (ne) and the mobilities (µH) in dependence on Vdis(norm) of the annealed TNO films measured at room temperature. At Vdis(norm) ≈ 1 the films exhibited low resistivity due to a high carrier concentrations and high mobilities. While the mobilities showed a bell shaped curve around Vdis(norm) ≈ 1, the carrier concentration increased when the depositions were driven into the metallic mode. After the annealing those films exhibited a minor crystallography quality and low mobilities. At Vdis(norm) >1 the carrier concentration and the mobilities of the TNO films decreased. To investigate the electronic transport in the TNO films, temperature14

ACCEPTED MANUSCRIPT dependent Hall and conductivity measurements have been performed, displayed in Fig.11. The samples prepared at Vdis(norm) <1 showed a thermally activated

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conductivity which is characteristic for semiconductors. If the samples were sputtered at

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Vdis(norm) ≥1, they showed an increasing conductivity with decreasing temperature.

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This behaviour is typical for a degenerate semiconductor and resembles that of metals. The carrier concentrations in those degenerated semiconducting films were independent of the temperature due to the vanishing donor ionization energy. On the other hand, the

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Hall mobilities rose with decreasing temperature caused by the reduced phonon

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scattering, which is also typical for a metal. The samples with a thermally activated conductivity exhibited too small Hall voltages, thus not allowing an extraction of carrier concentration and Hall mobility. This behaviour points to the well-known effect of

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electrical potential barriers at grain boundaries which strongly reduce the lateral current

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transport, i.e., the effective Hall mobility (see for instance [43]).

The T-dependence of the Hall mobility can be fitted by the equation [44]: µii  opt µii   opt

(eq. 1 ),

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 T  

where µii is the temperature-independent part of the total mobility, for instance due to ionized impurity scattering in degenerately-doped films. µopt is the temperaturedependent mobility, which describes the temperature dependence of the polar optical phonon scattering. The mobility µopt due to polar optical phonons is given as

µopt 

e 2m  LO *

  exp D   T 

(eq. 2),

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ACCEPTED MANUSCRIPT since the temperature T is low in comparison to the Debye temperature θD, where α is the dimensionless polaron coupling constant, m* is the effective mass and ωLO is the

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frequency of a longitudinal optical (LO) phonon [45]. The polaron coupling constant

1  1 1  m*c 2    137     s  2k B  D

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

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depends on the Debye temperature and is described by the relation

(eq.3),

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where ε∞ is the high-frequency dielectric constant and εS is the static dielectric constant. Fig. 12 shows the mobilities of the samples versus the temperature which exhibited the

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behaviour of a degenerate semiconductor (Vdis(norm) ≥ 1). The fitting curves of µ(T) are shown as red lines and the corresponding fitting parameters, µopt, µii and θD are listed in

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table 1. The Hall mobilities at room temperature increased if both the temperature-

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independent mobilities µii and the temperature-dependent mobility µopt increased. The

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fits revealed Debye temperatures extending from 340 to 513 K which were smaller than θD of microcrystalline TiO2 powder (θD =653 K) [46].

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The spatially-resolved microstructure of thin polycrystalline TNO films was analysed by EBSD. Figs. 13a and b shows the EBSD maps giving the orientation distributions of two polycrystalline TNO films sputtered under a) oxygen deficit Vdis(norm) ≤ 1) and b) under oxygen excess (Vdis(norm) > 1), and Figs. 13c and d give the corresponding EBSD pattern-quality maps. Grain boundaries are visible by dark lines in the patternquality maps since at these positions, the measured EBSD patterns are superposition of those of the neighboring grains, which results in low pattern quality. The compressive stress is visible in the EBSD maps by the slight change in local orientation within apparently contiguous crystalline regions (see Figs. 13a and b). However, larger grains 16

ACCEPTED MANUSCRIPT are found for the TNO films sputtered under slightly oxygen deficient conditions (about 30 µm vs. 10 µm). Also, for the TNO films grown under oxygen-rich conditions, a

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larger density of extended structural defects (probably stacking faults) is visible. We

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attribute these differences in microstructure to the differences in resistivity. Pore et al.

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[47] prepare amorphous TNO films by using atomic layer deposition and studied their crystallization in dependence on the composition and the post-deposition annealing. They also report grains (up to 50 µm) and a cross-like feature within the individual

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crystalline regions in the EBSD data of the annealed TNO films and attribute this

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microstructure to the explosive crystallization, i.e., the crystallization of an amorphous thin film occurs very rapidly after the first crystallization event, which is assumed to

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4. Conclusions

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start from anatase seeds in the centers of the large anatase crystallites.

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The [O]/[metal] ratio of the as-deposited TNO films predetermined the crystallographic structure and the electrical properties after annealing. Therefore, the precise control of the oxygen stoichiometry by adjusting the discharge voltage was essential during the

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reactive magnetron sputtering from a TiNb alloy target. After annealing in vacuum TNO films were highly doped (ne >2x1020 cm-3), transparent (α550 nm< 2x103 cm-1), and conductive (ρ ≈ 1.3x10-3 Ωcm). The EBSD analysis revealed large lateral crystallite sizes exceeding 20 µm with planar structural defects. X-ray diffraction revealed the anatase crystal structure and a CDS comparable with the films thickness. The Raman Eg modes showed a blue shift with an increase of the electron concentration whereas RBS analysis of these films revealed a decrease of the niobium concentration (<1 at%). Therefore, we tentatively ascribe that blue shift of Raman Eg modes to a distortion of the anatase structure. Furthermore, the Hall mobilities of highly doped TNO films 17

ACCEPTED MANUSCRIPT increase at low temperature which is typical for a metal. We assume that the electron mobility at room temperature (µe ≤ 8 cm² V-1s-1) is limited owing to phonon scattering,

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ionized impurity scattering, and probably also due to the presence of planar defects in

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the TNO films.

Figure and table captions

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Fig. 1: Deposition rates (closed symbols) for sputtering of amorphous TiO2:Nb at four oxygen flows (open symbols) at the adjusted discharge voltage. The magenta squares

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denote the deposition rate in the metallic and oxidic mode. The black diamonds plots the characteristic of the discharge voltage in dependence on the oxygen flow (right

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axis). Deposition parameters: p = 1 Pa, FO2(Pdc) =1.2 sccm (111 ± 2 W) (green), 1.3

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sccm (116± 2 W) (black), 1.4 sccm (126 ± 2 W) (blue), and 1.5 sccm (136 ± 2 W) (red).

Fig. 2: Resistivity of TiO2: Nb films as a function of the discharge voltage (a) and in

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dependence on the normalized discharge voltage (b), for different oxygen partial pressures. Deposition parameters: p = 1 Pa, FO2(Pdc) = 1.2 sccm (111 ± 2 W), 1.3 sccm (116 ± 2 W), 1.4 sccm (126 ± 2 W) and 1.5 sccm (136 ± 2 W) and p =1.5 Pa, 1.4 sccm (118 ± 2 W).

Fig. 3: Spectral transmittance and reflectance of a 130 nm thin TiO2: Nb film asdeposited, 100 Ωcm, (dotted black line) and after annealing, 1.1x10-3 Ωcm, (red solid line). The inset graph shows the absorption coefficient at 550 nm wavelength in 18

ACCEPTED MANUSCRIPT dependence of Vdis(norm) for the as-deposited and annealed films, respectively.

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Deposition parameters: p= 1 Pa, FO2 (Pdc)= 1.4 sccm (126 ± 2 W).

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Fig. 4: An RBS spectrum of a TNO-film on a glassy carbon substrate is shown,

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exemplarily (a). The [Nb]/([Ti]+[Nb]) ratios (b) and the [O]/[metal] ratios (c) of TNOfilms are depicted as a function of the discharge voltage. The horizontal dashed lines are

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calculated for a TiO2/Nb2O3 mixture, corresponding to the metal ratio of the target. Deposition parameters: p=1 Pa: FO2(Pdc)=1.3 sccm (116 ± 2 W), 1.4 sccm (126 ± 2 W)

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and 1.5 sccm (136 ± 2 W) and p=1.5 Pa, 1.4 sccm (118 ± 2 W).

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Fig. 5: Atomic concentrations of niobium (a), titanium (b), and oxygen (c) of asdeposited TNO-films in dependence on the discharge voltage. The horizontal dashed

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lines are calculated for a TiO2/Nb2O3 mixture, corresponding to the metal ratio of the target. Deposition parameters: p=1 Pa: FO2(Pdc)=1.3 sccm (116 ± 2 W), 1.4 sccm (126 ±

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2 W) and 1.5 sccm (136 ± 2 W) and p=1.5 Pa, 1.4 sccm (118 ± 2 W).

Fig. 6: X-ray diffraction patterns of the TNO films, prepared at different normalized discharge voltages. Depicted are four regions around the most intense reflections from the (101), (004), (200) and (211) planes. The blue vertical lines mark the positions of these peaks for anatase powder (CuKα1, JCPDS file No. 021-1272). Deposition parameters of the TNO film: p= 1 Pa, FO2(Pdc)=1.4 sccm (126 ± 2 W).

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ACCEPTED MANUSCRIPT Fig. 7: The positions of the 101 diffraction peaks (a) and the calculated coherently diffracting domain size (CDS) (b) of the annealed TNO films in dependence on the

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normalized discharge voltage. The dashed horizontal line marks the position of the

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(101) peak according to the reference pattern (JCPDS file No. 021-1272). Deposition

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parameters: p = 1 Pa, FO2(Pdc) = 1.2 sccm (111 ± 2 W), 1.3 sccm (116 ± 2 W), 1.4 sccm (126 ± 2 W) and 1.5 sccm (136 ± 2 W) and p =1.5 Pa, 1.4 sccm (118 ± 2 W).

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Fig. 8: Raman spectra measured at room temperature. The black dashed curve marks the

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annealed TNO film with low resistivity, deposited at Vdis(norm)= 1.000, and the asdeposited TNO film is shown as a green dashed line. A Raman spectrum of an anatase

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crystal is shown as dotted blue line for comparison [36]. Vertical blue lines mark the

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frequency modes of anatase, Eg mode (639, 197 and 144 cm-1), the Bg1 mode (399 cm-1), and the band at 516 cm-1 (A1g*) [34], a). An enlarged view of the low frequency mode Eg

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(144 cm-1) is shown in b).

Fig. 9: Raman shift and the corresponding peak width (FWHM) of the Eg mode (144 cm-1) versus the measured charge carrier concentration. The dashed horizontal lines mark the frequency of the Eg mode and the band width of pristine anatase [34, 48]. Deposition parameters: p=1 Pa, FO2(Pdc) = 1.4 sccm (126 ± 2 W) (blue triangles), FO2(Pdc) = 1.3 sccm (116 ± 2 W) (black triangles), and p =1.5 Pa, FO2(Pdc) = 1.4 sccm (118 ± 2 W) (blue circles).

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ACCEPTED MANUSCRIPT Fig. 10: Carrier concentration ne (a) and carrier mobilities µHall (b) at room temperature in dependence on the normalized discharge voltage for five deposition series.

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Deposition parameters: p= 1 Pa, FO2(Pdc) = 1.2 sccm (111 ± 2 W), 1.3 sccm (116 ± 2

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W), 1.4 sccm (126 ± 2 W) and 1.5 sccm (136 ± 2 W) and p=1.5 Pa, FO2(Pdc) = 1.4 sccm

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(118 ± 2 W).

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Fig.11: Transport properties (carrier concentrations (ne), conductivities (σ) and Hall mobilities (µHall)) of TNO films in dependence on the measured temperature. Deposition

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parameters: p= 1 Pa, FO2(Pdc)=1.4 sccm (126 ± 2 W).

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Fig.12: Hall mobilities (µHall)) of TNO films in dependence on the measured

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temperature. Deposition parameters: p= 1 Pa, FO2(Pdc)=1.4 sccm (126 ± 2 W). The

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fitting curves are shown as continuous red lines.

Fig. 13: EBSD maps of two annealed TNO films which were sputtered at Vdis(norm) ≈1 (left hand side) and Vdis(norm) >1 (right hand side) at a total pressure of 1 Pa, FO2(Pdc) =1.4 sccm (126 ± 2 W). a) and b) show the local orientation distribution and c), d) depict the pattern-quality maps, extracted from the same set of EBSD data.

Table 1: Transport properties carrier concentrations (ne), conductivities (σ) and Hall mobilities (µHall) of TNO films at 300 K in dependence on the normalized discharge 21

ACCEPTED MANUSCRIPT voltage Vdis(norm). Deposition parameters: p= 1 Pa, FO2(Pdc)=1.4 sccm (126 W). Fit parameters according to eq.1: Debye temperature θD, temperature-independent term µii,

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temperature-dependent term µopt (300 K), and the film thickness ds.

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References

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ACCEPTED MANUSCRIPT Furubayashi, N. Yamada, Y. Hirose, Y. Yamamoto, M. Otani, T. Hitosugi, T. Shimada and T. Hasegawa, Transport properties of d-electron-based transparent conducting oxide: Anatase Ti1-xNbxO2, J. Appl. Phys., 101 (2007) 093705. [24] K. Tonooka, T.-W. Chiu and N. Kikuchi, Preparation of transparent conductive TiO2:Nb thin films by pulsed laser deposition, Appl. Surf. Sci., 255 (2009) 9695-9698. [25] T. Hitosugi, N. Yamada, S. Nakao, Y. Hirose and T. Hasegawa, Properties of TiO2based transparent conducting oxides, phys. stat. sol. (a), 207 (2010) 1529-1537. [26] S. Berg and T. Nyberg, Fundamental understanding and modeling of reactive sputtering processes, Thin Solid Films, 476 (2005) 215-230. [27] M. Neubert, S. Cornelius, J. Fiedler, T. Gebel, H. Liepack, A. Kolitsch and M. Vinnichenko, Overcoming challenges to the formation of high-quality polycrystalline TiO2:Ta transparent conducting films by magnetron sputtering, J. Appl. Phys., 114 (2013) 083707. [28] M. Mayer, Ion beam analysis of rough thin films, Nucl. Instr. Meth. Phys. Res. B, 194 (2002) 177-186. [29] D. Depla, J. Haemers and R. De Gryse, Discharge voltage measurements during reactive sputtering of oxides, Thin Solid Films, 515 (2006) 468-471. [30] H.-Y. Lee and J. Robertson, Doping and compensation in Nb-doped anatase and rutile TiO2, J. Appl. Phys., 113 (2013) 213706. [31] K. Zakrzewska, Nonstoichiometry in TiO1-y Studied by Ion Beam Methods and Photoelectron Spectroscopy, Adv. Mater. Sci. Eng., 2012 (2012) 1-14. [32] S. K. Mukherjee, H. W. Becker, A. P. Cadiz Bedini, A. Nebatti, C. Notthoff, D. Rogalla, S. Schipporeit, A. Soleimani-Esfahani and D. Mergel, Structural and electrical properties of Nb-doped TiO2 films sputtered with plasma emission control, Thin Solid Films, 568 (2014) 94-101. [33] P. Scherrer, Bestimmung der inneren Struktur und der Größe von Kolloidteilchen mittels Röntgenstrahlen, Nachr. Königl. Gesell. Wiss. Göttingen, 2 (1912) 98-100. [34] T. Ohsaka, F. Izumi and Y. Fujiki, Raman spectrum of anatase, TiO2, J. Raman Spectrosc., 7 (1978) 321-324. [35] S. P. S. Porto, P. A. Fleury and T. C. Damen, Raman Spectra of TiO2, MgF2, ZnF2, FeF2, and MnF2, Phys. Rev. , 154 (1967) 522-526. [36] B. Lafuente, R. T. Downs, H. Yang and N. Stone, The power of data base: the RRUFF project in: T. Armbruster and R. M. Danisi (Eds.) Highlights in Mineralogical Chrystallography, De Gruyter, 2015, pp. 1 -30. [37] K.-R. Zhu, M.-S. Zhang, Q. Chen and Z. Yin, Size and phonon-confinement effects on low-frequency Raman mode of anatase TiO2 nanocrystal, Phys. Lett. A, 340 (2005) 220-227. [38] E. Uyanga, A. Gibaud, P. Daniel, D. Sangaa, G. Sevjidsuren, P. Altantsog, T. Beuvier, C. H. Lee and A. M. Balagurov, Structural and vibrational investigations of Nb-doped TiO2 thin films, Matter Res. Bull., 60 (2014) 222-231. [39] B. Choudhury, M. Dey and A. Choudhury, Defect generation, d-d transition, and band gap reduction in Cu-doped TiO2 nanoparticles, International Nano Letters, 3 (2013) 25. [40] X. Lü, W. Yang, Z. Quan, T. Lin, L. Bai, L. Wang, F. Huang and Y. Zhao, Enhanced Electron Transport in Nb-Doped TiO2 Nanoparticles via Pressure-Induced Phase Transitions, J. Am. Chem. Soc., 136 (2014) 419-426.

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ACCEPTED MANUSCRIPT Sahoo, A. K. Arora and V. Sridharan, Raman Line Shapes of Optical Phonons of Different Symmetries in Anatase TiO2 Nanocrystals, J. Phys. Chem. C, 113 (2009) 16927-16933. [42] S. K. Gupta, R. Desai, P. K. Jha and A. K. Arora, Effect of Annealing Time on Synthesis of Titanium Dioxide Nanocrystals Studied by Raman and Photoluminescence Spectroscopy, AIP Conf. Proc., 1249 (2010) 125-128. [43] K. Ellmer and R. Mientus, Carrier transport in polycrystalline transparent conductive oxides: A comparative study of zinc oxide and indium oxide, Thin Solid Films, 516 (2008) 4620-4627. [44] K. Ellmer and R. Mientus, Carrier transport in polycrystalline ITO and ZnO:Al II: The influence of grain barriers and boundaries, Thin Solid Films, 516 (2008) 58295835. [45] M. Grundmann, The Physics of Semiconductors, An Introduction including Nanophysics and Applications, second ed., Springer-Verlag, Berlin Heidelberg, 2006. [46] J. Luo, X. Ding, B. Chen, J. Kong and Y. Dong, Grain Growth Kinetics and Debye Temperature of Nanometer TiO2 Powders Prepared by a Sol-gel Process, J. Mater. Sci. Technol., 10 (1994) 213-216. [47] V. Pore, M. Ritala, M. Leskelä, T. Saukkonen and M. Järn, Explosive Crystallization in Atomic Layer Deposited Mixed Titanium Oxides, Cryst. Growth Des., 9 (2009) 2974-2978. [48] H. C. Choi, Y. M. Jung and S. B. Kim, Size effects in the Raman spectra of TiO2 nanoparticles, Vib. Spectrosc., 37 (2005) 33-38.

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ACCEPTED MANUSCRIPT σ

µopt

ΘD

(nm)

(1020 cm-3)

(Scm-1)

(cm²V-1s-1)

(K)

1.020

115

0.85

24.4

1.8

7.5

340

1.014

115

2.53

142

3.5

6.5

7.6

513

1.007

115

6.46

683

6.6

12.7

13.2

356

1.000

130

9.71

763

4.9

10.7

9.5

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1.000

500

6.63

865

8.15

22.3

13.9

452

µHall

(cm²V-1s-1)

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(cm²V-1s-1)

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Table 1: Transport properties carrier concentrations (ne), conductivities (σ) and Hall mobilities (µHall) of TNO films at 300 K in dependence on the normalized discharge

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voltage Vdis(norm). Deposition parameters: p= 1 Pa, FO2(Pdc)=1.4 sccm (126 W). Fit parameters according to eq.1: Debye temperature θD, temperature-independent term µii,

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ACCEPTED MANUSCRIPT Highlights

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 slight oxygen deficient as-deposited films were highly conductive after annealing  control of oxygen stoichiometry by adjusting the discharge voltage during deposition  electron mobility at room temperature is limited due to scattering at phonons  films exhibited large average crystallite sizes with planar structural defects

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