Phase Transformations in the Pulsed Laser Deposition Grown TiO2

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Phase Transformations in the PLD Grown TiO Thin Films as a Consequence of O Partial Pressure and Nd Doping 2

Awais Ali, Inci Ruzybayev, Emre Yassitepe, Altaf Karim, S. Ismat Shah, and Arshad Saleem Bhatti J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b01046 • Publication Date (Web): 28 Apr 2015 Downloaded from http://pubs.acs.org on April 29, 2015

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The Journal of Physical Chemistry

Phase Transformations in the PLD Grown TiO2 Thin Films as a Consequence of O2 Partial Pressure and Nd Doping A. Alia, I. Ruzybayevb, E. Yassitepec, A. Karima, S. Ismat Shahb,c and A. S. Bhattia* a

Centre for Micro and Nano Devices, Department of Physics, COMSATS Institute of Information Technology, Park Road, Islamabad 44000, Pakistan

b

Department of Physics and Astronomy, University of Delaware, Newark, Delaware 19716, United States

c

Department of Material Science and Engineering, University of Delaware, Newark, Delaware 19716, United States

ABSTRACT: In this work, we present the pronounced phase transformation in the TiO2 thin films grown in simultaneously varied oxygen (O2) partial pressure and doping concentration of neodymium (Nd). Pulsed laser deposition (PLD) was employed for the synthesis of TiO2 films in varied O2 partial pressures from 100 mtorr (O2 rich) to 0 mtorr (O2 absent) for each Nd doping concentration (1.0, 1.5 and 2.0 at.%). The phase and structural studies confirmed the systematic phase transformation from the anatase to the rutile to the metastable and finally to the amorphous phase as the O2 content was reduced and Nd doping concentration was increased. A drastic reduction in the crystallite size was also observed. XPS confirmed the reduction of Ti (Ti4+ to Ti3+), creation of O vacancies and consequently increased non-stoichiometry and short range ordered TiO2 lattice. The experimental findings were complimented with the density functional theory calculations of surface energies and crystal facet areas of the anatase and rutile phases. The Wulff equilibrium particulate shapes of the anatase and rutile phases confirmed the possibility of transformation due to increased Nd doping and reduced oxygen. Thus, the obtained results suggested that the Nd doping and O2 concentration during the growth were key parameters to precisely control the phase of the TiO2 films and its transformation from the anatase to rutile to a metastable phase and finally to amorphous phase.

Key words: TiO2, PLD, Phase transformation, Doping. *Corresponding Author: [email protected]; Tel: +92-51-90495118; Fax: +92-519247006

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I.

INTRODUCTION: TiO2 is one of the most studied oxide semiconductor in last couple of decades. TiO2

mainly exists in two polymorphs, anatase and rutile and show different physical properties due to varied oxygen content in the unit cell, which affects the c/a. The anatase phase transforms irreversibly to the rutile phase at elevated temperatures as a result of increased oxygen vacancies. The precise control on the growth parameters is of key significance, which ultimately determines the formation of phase and its transformation from the anatase to the rutile phase. The knowledge of the phase is particularly important in high temperature applications and processes, such as porous gas separation membranes and gas sensors,1-3 in which the phase transformation may take place and consequently change its properties and performance. For this reason, an understanding of the formation of the initially grown phase, phase transformation, stability and progressions involved in controlling the phases is of fundamental importance as it can result in having multiphase or even single phase microstructures that would affect the performance of the ultimate device. The growth of TiO2 films has been reported using various techniques where the preliminary crystalline TiO2 phase produced is generally anatase.4,5 From the point of view of structure, this is due to the superior easiness of the short range ordered TiO6 octahedral in assembling into long range ordered anatase structure. This is owed to a less constrained molecular structuring of the anatase phase compared to the rutile phase.6 From the thermodynamic point of view, the extra fast crystallization of the anatase may be the result of the requirement of low surface free energy irrespective of the low Gibbs free energy requirement for the rutile phase formation.7,8 In short, the high surface free energy requirement of the rutile phase crystallites favors the crystallization of TiO2 in the anatase phase. In addition, the realization of various TiO2 phases depends significantly on the synthesis parameters. The temperature and the growth time are typically considered in the kinetics of the growth processes. In temperature dependent growth conditions, irreversible transformation to the rutile phase from the pure anatase phase is widely considered to begin in air at around 600 o

C.7,9,10 However, the reported transition temperatures varied from 400 to 1200 oC due to the use

of different synthesis techniques, raw materials and processing methods.9,11-14 Due to reconstructive behavior; anatase to rutile transformation has also been shown as a time dependent process.15-17 It is broadly understood that the amount and position of defects in the 2 ACS Paragon Plus Environment

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oxygen (O) sublattice, viz., TiO2-x are the most significant factors affecting the phase transformation.18-21 The reduction of large O sublattice (lessening of structural rigidity) can enhance the simplicity of rearrangement and transformation by using large number of O vacancies.16,22,23 In different conditions, this result has been shown during heat treatment, where reducing or neutral atmospheres of low O2 partial pressure has enhanced the transformation of considerable amount of anatase to rutile.18,24,25 In certain experiments, it has also been demonstrated that the vacuum conditions could restrict the phase transformation.16 This was due to creation of oxygen vacancies that has been known to affect the phase transformation process significantly. On the other hand, the addition of dopants has been shown to enhance or impede the phase transformation by effecting the atomic rearrangement.

Generally, Nd doping is

considered an inhibitor to the phase transformation in TiO2.26 It has also been reported that incorporation of Nd produced stresses in the lattice due its larger ionic radius (34 %), which resulted in the textured growth of TiO2 along [004].27,28 In this paper, the role of O2 vacancies and Nd doping in affecting the structural, phase transformation and chemical composition is studied. Nd doping and oxygen concentrations during synthesis of TiO2 films were varied in a systematic way by using the pulsed laser deposition (PLD) system. The oxygen partial pressure was varied in a way to achieve O2 rich to no O2 atmospheres for different Nd doping concentrations during the growth.

The phase

transformations were successfully observed from the anatase to rutile to metastable and finally to amorphous TiO2 formations by suitably varying the O2 flux and Nd doping. The role of Nd concentrations and oxygen stoichiometry was studied in detail in the formation of various phases during the synthesis using X-ray diffraction, X-ray photoemission spectroscopy and Raman spectroscopy. Density functional theory (DFT) based calculations of low energy surfaces/phases complemented the experimental findings. The present results successfully demonstrated that the growth of a specific TiO2 phase could be achieved by manipulating the growth parameters. II.

EXPERIMENT AND COMPUTATION: PLD was employed to deposit TiO2 films from Nd doped TiO2 target powder prepared

with three different compositions of Nd (1.0, 1.5, and 2.0 at.%) by mixing TiO2 powder (99.999% Sigma-Aldrich) and Nd2O3 powder (99.9%, Acros) in a solid state reaction. Target pallets of 2 cm diameter each, were statically pressed, followed by sintering at 800 oC for 6 3 ACS Paragon Plus Environment


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hours. Plain microscope glass substrates (Fisher Scientific) were sequentially cleaned by sonication in solutions of detergent, isopropyl alcohol, and acetone and thoroughly rinsed in the deionized (DI) water and dried in dry nitrogen immediately prior to placing in the PLD chamber. A base pressure of 2×10-6 torr was achieved before introducing oxygen in the chamber. O2 partial pressure in the chamber was varied for the growth of each doped TiO2 thin film , i.e., 100 (O2 rich atmosphere), 10, 1 to 0 mtorr (O2 deficient atmosphere). The deposition was done with KrF excimer laser irradiating at a power of 450 mJ on a rotating target for 20 minutes. The pulse repetition rate was 15 Hz, and the substrate temperature was kept at 500 oC The phase analysis was performed by the X-ray diffraction operating at 30 kV and 30 mA, The 2θ scans were taken in the range of 20o to 80o with a step size of 0.01o in a grazing angle incident geometry. Raman spectra were collected with an excitation energy of 532 nm from Nd:YAG laser of 20 mW incident power at room temperature. Film thicknesses were determined from the field emission scanning electron microscopy (FESEM) equipped with energy dispersive X-ray (EDX) spectroscopy system. Valance states of the grown films were investigated by an X-ray photoelectron spectroscopy (XPS) system (Omicron EA125). Incident beam of non-chromatic Al X-ray (1486.5 eV) was used operating at 10 kV, 10 mA, and 100 W for both survey and high resolution scans. Pass energy was set at 50 eV for the survey scan and 25 eV for high resolution scan with a dwell time of 2 s/step for both types of scans. Measured peaks were charge corrected with respect to C-1s peak positioned at 284.6 eV. We have performed density functional calculations (DFT) using VASP28,29 code with the projected augmented wave (PAW) method.30 These calculations were carried out within the generalized gradient approximation (GGA).31 For the bulk TiO2 energy calculations, a conventional tetragonal (4/m 2/m 2/m) cell of anatase TiO2 containing 4 formula units was used. The energy cutoff of 300 eV is taken as an input to the DFT calculations along with the 8 × 8 × 8 k-point sampling Monkhorst-Pack scheme.32 We prepared the super cell of anatase TiO2 based on the above mentioned parameters and then relaxed it.

Later for the surface energy

calculations, same parameters as mentioned above, were used. The slabs of the unrelaxed (004), (101), (110), and (011) structures were obtained from the relaxed bulk anatase TiO2. To initiate the slab calculations, sets of finite number of material layers were arranged periodically by placing a vacuum layer between them. The thickness of the vacuum layer was taken to be 10 Å, which was essential to remove any type of interactions between the periodically arranged slabs. 4 ACS Paragon Plus Environment

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In the first step of the energy minimization and slab relaxation process, the volume of the whole slab including all atoms was allowed to relax. During the second step of the relaxation process, only 2 top and 2 bottom layers were set to relax. In this process a force convergence criterion of 0.02 eV/Å was implemented. The same procedure was adopted for the Nd doped TiO2 surface calculations. While performing the calculation on the Nd doped systems/surfaces, little displacements (20% of the nearest neighbor distances) in random directions to all surfaces atoms up the 2nd layer were also introduced.

After introducing the Nd dopant and random

displacements, surfaces were allowed to relax to obtain the surface energies. Further, the Wulff construction method was applied to obtain the equilibrium shape of the anatase and the Nd doped rutile crystals. In this method the particle shapes were obtained by minimizing the total surface energy.33 III.

RESULTS AND DISCUSSION: A. X-Ray Diffraction: Figure 1 (a-d) shows the XRD patterns of Nd doped TiO2 thin films synthesized at

various O2 partial pressures. For the highest O2 partial pressure, i.e., 100 mtorr, all Nd doped TiO2 grown films showed the anatase phase, as seen in Figure 1 (a). Two major reflections, i.e., (101) and (004) were observed at the expected 2θ values, which confirmed the presence of anatase phase (as per JCPD card # 02-0387). The peak intensity of the major peak (101) decreased by almost 50 % with the increase in Nd doping from 1.0 to 1.5 at% and its crystallinity degraded and its intensity dropped by almost 90 % for 2 at.% doping with respect to 1 at.% doped TiO2 film. However, the 2nd major reflection (004) showed a lesser drop in the intensity, i.e., by 25% and 30% with the increase in Nd doping to 1.5 and 2.0 at.%, respectively, with respect to 1 at.% doped films. In summary, the main effect of the increased Nd doping concentration in the TiO2 film grown at 100 mtorr of oxygen partial pressure was the reduced crystallinity. In short, the ratio of intensities of two reflections I(101)/I(004) was 2.6, 2.1 and 1.1 for 1.0, 1.5 and 2.0 at.% Nd doping. The reduced crystallinity was due to inhibition of growth at the grain boundaries with the addition of Nd, and enhanced texturing along (004) plane.

28

The

difference in the ionic radii of the dopant atom and the host lattice atom can influence the location of the substitution of the dopant atom in the host lattice, i.e., to substitute or to find an interstitial site. The addition of the dopant could also cause stress, either compressional or 5 ACS Paragon Plus Environment


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tensile. Nd has almost 34% bigger ionic radius than Ti, so its addition in the TiO2 lattice can cause the tensile stress ,28 which can lead to inhibit the grain growth. The accommodation of large ion could be facilitated along the c-axis because anatase has tetragonal structure with almost 60% larger c-axis than the other two equally spaced axes. The observation of the incomplete or partial texturing was possibly due to fast growth rate as compared to complete texturing observed earlier.28 It is worth mentioning that PLD is a very energetic process with very high growth rate of ~ 40 Å/sec compared to 0.6 Å/sec rate of growth obtained in the sputter deposition.28 The films grown at 10 mtorr of O2 partial pressure (Figure 1 (b)) were mainly consisted of rutile phase with a reduced presence of the anatase phase for all Nd concentrations. The only observed diffraction peak of the anatase phase was (101), while a good number of rutile peaks were identified as (110), (200), (111) and (210) at the expected 2θ values according to JCPD card# 1-1292. The wt. % of the anatase phase as a function of Nd concentration in each doped film was determined using Equ. (1);34

WA % =

1    I R   1 + 1.265 I A   

× 100

(1)

Where, IA and IR are the peak intensities of anatase and rutile phases, respectively. With increase in doping concentration from 1 to 2 at.%, the anatase to rutile phases ratio dropped from 10:90 (±1) to 02:98 (±1), respectively. The quantitative analysis revealed that even in the mixed phase grown TiO2, the relative ratio of the anatase phase to the rutile phase could still be tailored by simultaneous tuning the Nd and O2 contents. Similarly, TiO2 films, grown at fixed partial pressure of O2 (10 mtorr) but with varied concentrations of Nd concentration followed the same pattern as that was observed in the case of 100 mtorr growth. The films crystallinity was also reduced by almost 50%, when Nd concentration was increased from 1 to 2 at.%. The conversion of anatase to rutile with reduced O2 content from 100 to 10 mtorr and increased Nd concentrations from 1 to 2 at.% was explained on the basis of the respective crystal structures. The reconstructive transformation from the anatase to rutile phase implied the breaking and reforming of bonds involved in the transformation.35 An overall volume contraction of 8% and general contraction of the c-axis caused the observed transformation of phases,36 and increased percentage of the rutile phase relative to the anatase phase. Thus, growth of TiO2 in the reduced 6 ACS Paragon Plus Environment

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O2 with simultaneous substitution of Nd caused the crystal structure of the anatase phase to become thermodynamically unstable and transformed into a smaller cell of the rutile phase, which was thermodynamically stable. The growth of TiO2 with varied Nd doping concentrations at the fixed O2 concentration to 1 mtorr (Figure 1 (c)) even showed more interesting and unusual results. Initially for 1 at.% of Nd doped TiO2 films, (110) and (200) reflections were identified as emerged from the rutile phase. Interestingly, when Nd concentration was increased to 1.5 at.%, preferential texturing of rutile phase along (200) plane was observed. Furthermore, for 2 at.% of Nd doping, a new reflection was observed at the 2θ = 38.3o, which has not been listed in any JCPD card available for TiO2. Hence it was tentatively referred to as a new metastable phase, which has been observed earlier as well,37,38 but rarely. These phases tend to occur in the highly O deficient growth regimes and in the present case, its appearance was attributed to Nd doping in the oxygen deficient environment.

As already explained, the growth in the reduced O2 content, the

emergence of rutile was due to a smaller c/a ratio (c/a ratio for anatase is 2.52 and that of rutile is 0.64) was more favorable. The emergence of new phase was explained on the basis of the excessive stress caused by Nd addition and excessive removal of O2 (which further shrunk the already rutile crystal lattice and made it thermodynamically unstable) and transformed into a crystal structure with even smaller c/a ratio. Finally, the growth of different Nd doped TiO2 thin films were carried out in completely depleted oxygen environment (Figure 1 (d)). At 1 at.% the metastable phase was observed to be formed as a reflection again appeared at the 2θ value of 38.3o but this time with improved crystallinity as the FWHM was sharp and decreased by 40% and integrated intensity increased almost three times compared to the one observed earlier. On increasing the doping concentration to 1.5 at.%, the metastable phase was still observed but with a reduced crystallinity as the integrated intensity dropped sharply by almost 64% with increased in FWHM of 15%. Finally, the growth with 2 at% of Nd, amorphous phase of TiO2 was obtained. The systematic change in the oxygen environment with the controlled variation of Nd doping demonstrated that the phase stability of higher crystallinity and long range order was not possible. Figure 2 (a-c) displays the XRD patterns arranged to highlight the effect of Nd doping as a function of variation of O2 concentration (as labeled for each pattern in Figure 2). The phase transformation was much pronounced as shown in Figure 2 (a), where for 1 at.% doping, phase 7 ACS Paragon Plus Environment


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changed from pure anatase to mixed (anatase : rutile = 10 : 90) phases to pure rutile and then finally to metastable phase as the O2 concentration varied from 100 mtorr to 0 mtorr. As the doping concentration increased to 1.5 at%, the trend for the phase transformation remained the same but the quality of crystallinity of each phase reduced. Finally, for the maximum doping of 2 at%, the TiO2 structure further deteriorated and transformed completely to the amorphous phase. The combined effect of O2 partial pressure and Nd doping on the sizes of crystallites of the three phases was determined from the Scherrer’s equation and is plotted in Figure 3. The decrease in the crystallite size actually reflected the hindered growth with the increase in Nd doping concentration.26,28 As reported in earlier work,28 if TiO2 doped with Nd under thermodynamically equilibrium conditions, it substitutes Ti. The combined effect of doping and oxygen concentration variation on the surface morphology and the thickness of the films were further studied by the scanning electron microscopy. B. Microstructure of TiO2 films: XRD analysis has confirmed that the doping of Nd and variation in oxygen concentration both have affected the growth kinetics of the film and scanning electron microscopy confirmed that both parameters affected the thicknesses of films in a systematic way. Figure 4 (a-d) shows the cross-sectional SEM images of the four TiO2 doped with the 1 a% Nd thin films synthesized at various O2 partial pressures; (a) 100 mtorr; (b) 10 mtorr; (c) 1 mtorr; and (d) 0 mtorr. A drastic decrease in the thin film thickness was observed, even though the growth time was identical for all films under identical growth conditions. For 100 mtorr the thickness measured was almost 4.9 ± 0.2 microns and it decreased down to almost 0.35 microns when grown at 0.1 mtorr of O2 partial pressure. These results are plotted in Figure 5. The decrease in the thickness was quite significant, i.e., by almost 93%. The reduction in the thickness of the film could be understood on basis of few important crystallographic characteristics of each phase. Firstly, due to smaller c/a ratio of the rutile phase than the anatase phase, growth along the c-axis could result in smaller thickness of the film. Furthermore, the anatase phase unit cell comprises of 4 molecules per unit cell whereas rutile has 2 molecules per unit cell. Also, the basic building unit for the TiO2 consists of TiO6 octahedral and varies in both phases by arrangement. In the case of rutile crystal, each octahedron is connected to two-edge sharing and eight corner sharing adjacent octahedra, in which the edge8 ACS Paragon Plus Environment

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shared octahedra are associated along the (001) surface direction. On the other hand, in the anatase crystal there are four-edge sharing neighbors, which are associated along (100) and (010) surfaces, thus creating crisscross double chains perpendicular to the c-axis. 26 Thus making open channels parallel to the c-axis in rutile and perpendicular to the c-axis in anatase. This makes anatase more open and porous structure compared to rutile. It is well known that the rutile phase has a higher density compared to anatase phase, i.e., 4.25 and 3.98 g/cm3, respectively. Any variation in the bond lengths of the octahedral due to oxygen vacancies also played a very important role in defining the ultimate product. Reduced oxygen environment during the growth allowed the TiO2 structure to rearrange itself from comparatively oxygen rich, porous, open and larger in c/a ratio (anatase phase) to the rutile phase, which was intrinsically oxygen reduced, dense, closed and also had smaller c/a ratio. In addition, substitution of the dopant also resulted in the generation of excess oxygen vacancies.28 Thus, reduction in oxygen content and addition of dopant during the growth helped the phase transformation from anatase to rutile and consequently reduced the film thickness. A detailed study of the chemical composition was performed by the XPS to determine the simultaneous effect of oxygen concentration and Nd doping on the grown TiO2 thin films. The content of Ti and O in TiO2 films shown in Figure 5 was determined from the XPS measurements and is summarized in Table I. It was observed that growth in reduced oxygen environment led to significant increase in oxygen vacancies and presence of appreciable Ti3+ valencies in Ti rich films.

The drop in crystallite size also

contributed the reduction of thickness observed in the SEM.

Table I: Summary of XPS measurements showing the variation of Ti and O contents in 1.5 at.% Nd doped TiO2 films synthesized in various oxygen environments. Last column shows variation in the film thickness as determined from SEM micrographs.

Growth Parameters

Valance States

Phase of TiO2

Nd Conc. (at.%)

O2 Conc. (mtorr)

XRD

Raman

1.5 1.5 1.5

100 10 1

A R T.R

A R M.S

Thickness (µm)

Ti4+ (%)

Ti3+ (%)

O attached to Ti

O Vacancies

100 100 100

0 0 0

LBE (%) 87 82 75

HBE (%) 13 18 25

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4.90 1.00 0.70


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1.5

C.

0

(dist.) M.S (dist.)

M.S (dist.)

66

34

39

61

0.38

X-ray photoemission spectroscopy (XPS): XPS was employed to determine the stoichiometry and the chemical composition of the

doped films. Figure 6 shows the high resolution XPS spectra of 1.5 at.% Nd doped TiO2 thin films at various O2 partial pressures in the energy range for Ti, and two distinct peaks Ti-2p1/2 and Ti-2p3/2 were observed in all grown samples. A clear trend was observed in the shift of peak positions towards low energy with the reduction of oxygen (shown by dashed lines in Figure 6). For the film grown at 100 mtorr O2 partial pressure, the peak positions were 463.89 and 458.10 eV, respectively. It is well known that for a completely oxidized Ti, these values should be 464.21 and 458.50 eV, for Ti-2p1/2 and Ti-2p3/2, respectively. Both peak positions were at low binding energy (LBE), which meant that even at 100 mtorr O2 partial pressure, Ti was not completely oxidized. This was attributed to the addition of Nd3+ and a fast growth rate of films of almost 4 Å/sec, which resulted in the oxygen deficient growth of TiO2. The film grown at 10 mtorr also showed both characteristic peaks of Ti with a further downward shift as expected because Ti had less O2 to form TiO2. The downward shift in energy was more pronounced in the films grown at 1 mtorr and the shift was the largest in films grown in the O2 absent environment. The overall shift towards LBE was almost 0.3 eV for both peaks. Interestingly, for films grown at 0 mtorr, the XPS spectra was thoroughly analyzed to determine the contribution of new states of Ti responsible for emerging phases by using XPSPEAK as shown in Figure 7. The fits of XPS spectra confirmed the existence of two valence states, usual Ti4+ along with the emergence of new Ti3+ representing the substantial presence of highly reduced Ti states. The effect of Nd doping concentration in the O2 depleted growth is shown in the peak fittings for all (1.0, 1.5 and 2.0 at.%) films (Figure 7). The fitting results revealed that all samples had Ti4+ and Ti3+ states and the contribution of Ti3+ increased and Ti4+ decreased as the doping concentration of Nd was increased and is plotted in Figure 8. The effect of oxygen vacancies on the stoichiometry of TiO2 was also determined from the separation in the doublet Ti-2p peaks (∆ = 2p1/2 − 2p3/2). For completely oxidation (as in TiO2), Ti is in the Ti4+ state and the value of ∆ is 5.72 eV and for completely reduced Ti (as in Ti 10 ACS Paragon Plus Environment

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metal), it is in the Ti0 state, the value of ∆ is 6.05 eV. When the doping concentration of Nd was fixed at 1.5 at.% , the value of ∆ varied as 5.75, 5.81, 5.72 to 5.71 eV, in films grown in oxygen partial pressures of 100, 10, 1, and 0 mtorr, respectively. It was interesting to note that the doublet peaks showed a shift towards LBE (reduction of TiO2), while the value of difference ∆ did not show significant trend in shift. However, on performing deconvolution of Ti2p peaks into Ti4+ and Ti3+ states, the difference in ∆ became appreciable. The peak fitting for Ti4+ state showed that the value of ∆ was 5.93 eV. This further revealed the situation that at 0 mtorr (O2 absent) growth, TiO2 was converted to more metallic Ti character and contributed in the generation of the metastable phase. The ∆ value was also determined in films grown in oxygen deficient environment (0 mtorr). For 1.0 at.% Nd, ∆ was 5.86 eV, for 1.5 at.%, it increased to 5.93 eV and for the highest doping concentration of 2 at.%, it increased to 6.01 eV. This confirmed that incorporation of Nd in 3+ states in already reduced oxygen growth environment further created oxygen vacancies in the grown film. The XPS findings complimented the results of XRD and SEM and confirmed that oxygen was the key factor for defining the ultimate phase of the grown films, either by directly controlling oxygen content during growth or indirectly by adding dopants. The high resolution scans of O-1s region as shown in Figure 9 were studied in detail to reconfirm the role of oxygen stoichiometry in the grown films under varying growth conditions. For fixed 1.5 at. % Nd doped TiO2 thin film, the binding energy peak was observed at 529.5 eV, which corresponded to O-2 bonded with Ti+4 (LBE). However, there was another peak observed at 531.5 eV (HBE), which is observed in the case of presence of O vacancies. For a fixed Nd doping, the relative ratio of the two peaks showed an inverse trend with the decrease in O2 partial pressure during the growth as representatively plotted in Figure 10 for 1.5 at% Nd doping. It can be seen that with the decreased O2 content, the intensity of the LBE peak reduced and the intensity of the HBE peak became appreciable. Thus systematic transformation from pure anatase to mixed phases (rutile + anatase), to pure rutile, formation of a new (metastable) phase and finally formation of the amorphous phase was attributed to the gradual co-effect of the variation of oxygen concentration and Nd doping. Thus, the XPS findings confirmed the transformation and formation of new phases were the consequence of highly reduced Ti and generation of O vacancies caused by small amount of O2 available during the growth and by Nd doping in the TiO2 lattice. It can be concluded here that the structure of TiO2 rearranged itself to 11 ACS Paragon Plus Environment


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a short ranged order structure (as decrease in crystallite size with the increase in Nd concentration was shown in Figure 3) with less number of molecules per unit cell and made it more compact as well.

D. Raman spectroscopy: The effect of transformation of phases and formation of new phases in various TiO2 films was studied by Raman spectroscopy and Raman spectra for each doping concentration and varied oxygen partial pressure are shown in Figure 11. The anatase phase showed modes at 144, 197, 399, 515 and 639 cm-1, attributed to the Raman active modes of the anatase phase with the symmetries of Eg (144cm-1), B1g (197cm-1), A1g (399cm-1), B1g (515cm-1) and Eg (639cm-1), respectively.39 The typical Raman modes due to the rutile phase appeared at 143, 235, 447, and 612 cm-1, which were ascribed to the B1g, two-phonon scattering, Eg, and A1g modes of rutile phase, respectively.39 Additionally, the band at 144 cm-1 was the strongest for the anatase phase and the band at 143 cm-1 was the weakest for the rutile phase. Raman modes observed at smaller wave numbers (143 and 144 cm-1) were not considered as they overlapped with each other and could not depict the true picture. In addition, the symmetry of the lattice was also studied by the Raman spectroscopy. The symmetric stretching vibration of O-Ti-O bonds in TiO2 primarily produces Eg mode, and the symmetric bending vibration of O-Ti-O is related to the B1g mode, whereas the anti-symmetric bending vibration of O-Ti-O produces the A1g mode.40 The discussion on the behavior of Raman modes on doping and oxygen vacancies is carried out independently before discussing the combined effect. Figure 11 (a – d) shows the Raman spectra of Nd doped TiO2 films grown at different O2 partial pressures, namely, (a) 100 mtorr; (b) 10 mtorr; (c) 1 mtorr; and (d) 0 mtorr. The anatase and rutile phases showed active Raman modes at expected wave numbers as discussed above.. Initially, for 100 mtorr (Figure 11 (a)) grown films, all identified modes B1g, A1g and Eg belong to the anatase phase and are quite sharp. As the doping concentration increased, appreciable decrease in the peak intensities and increase in the peak broadening was observed for all Raman modes. This was a typical signature of the degraded crystallinity or increase in symmetry breaking as a consequence of doping in the films as it relaxed the Raman selection rules near the center of the Brillioun zone and made a range of q vectors available (where ∆q~1/L, where L is the length of the crystallite).41 With increase in the doping to 2 at%, the observed decrease in the integrated intensity of the 12 ACS Paragon Plus Environment

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symmetric (Eg) and anti-symmetric (A1g) mode was almost 72 % and 60 %, respectively. Also, the relative ratio of the Eg to the A1g modes decreased from 2.0 to 1.5 with the increase in doping from 1 to 2 at.%. These results indicated the generation of anti-symmetry in TiO2 with the addition of Nd. This was attributed to the possible substitution of Nd3+ in place of Ti4+ and the difference in the valance states of the two actually thought to have generated the oxygen vacancies leading to the nonstoichiometry in the host TiO2 lattice. The observed behavior of the doped films grown at 10 mtorr O2 partial pressure was quite different than observed in the films grown at 100 mtorr of O2 (Figure 11 (b)). The observed Raman signal emerged from the overlapped Raman modes of the anatase and rutile phases so deconvolution of modes was performed to extract the contribution of the each phase. At 1 at.%, mainly rutile phase (96%) with a small portion of anatase (4%) was obtained. In films doped with 1.5 at.%, anatase phase almost disappeared and 99% rutile phase was obtained. With the maximum doping of 2.0 at%, both A1g and Eg modes of rutile flattened, which clearly showed absence of symmetry to result any Raman mode. It was also concluded that nonstoichiometry contributed in the symmetry breaking as was confirmed in the XPS results. In addition, a second Raman mode at around 550 cm-1 emerged when doped with 1.5 at.% Nd or higher, which overlapped with the principal Raman mode for the rutile phase. Films grown at 1 mtorr showed (Figure 11 (c)) very weak and broad Raman modes of the rutile phase. At 1 at.%, it was distorted rutile and at 1.5 at. % a new mode appeared which peaked at around 557 cm-1 and for 2 at. % Nd doping, it shifted to lower wave numbers (Figure 11 (c)). This mode was assigned to an unidentified metastable phase. Finally, the Raman spectra in the films grown in the absence of O2 (Figure 11 (d) showed emergence of the metastable phase, which became distorted and very broad in films doped with 1.5 at. % Nd. Finally, when films doped with 2 at. %, no Raman signal was observed, which meant complete loss of symmetry or absence of long range order in the TiO2 lattice. Raman spectroscopy results have also been plotted by keeping the doping concentration fixed and varying the O2 partial pressure as shown in Figure 12 (a – c); where (a) 1.0 at. %; (b) 1.5 at.%; and (c) 2.0 at.%. The phase transformation could be observed readily. In case of 1 at.%, the transformation was from the anatase to the mixed phase (rutile : anatase = 96:04) to the pure rutile phase and finally to the metastable phase. Similar trends were also observed for the higher doping concentrations of Nd (Figure 12 (b – c). The analyses of the Raman spectroscopies 13 ACS Paragon Plus Environment


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were found in excellent agreement with the results of XRD. The transformation from the initial anatase phase to the ultimate amorphous phase was the result of reduced oxygen and increased Nd doping, and was the main cause of the reduction of the grown film thickness as seen in cross sectional SEM micrographs. Furthermore, the XPS results confirmed the increase in oxygen vacancies and the reduction of Ti with the increased Nd doping concentrations in the reduced oxygen environment during the growth.

E. Calculated surface energies and facet areas: The surface energies were calculated using the following standard formula7 for the anatase and Nd doped rutile (004), (101), (110), and the (011) low-index directions. The total energies of bulk and slab were obtained using the DFT calculations as described in the section II. γ = (Eslab −NEbulk)/2A

(2)

where γ is the computed surface energy, Eslab is the energy of the slab, Ebulk is the bulk energy per atom, N is the number of atoms in the slab, and A is the base area of the slab. Computed surface energies are tabulated in the Table 1. According to the computational data, (101), and (011) surfaces are the most stable of anatase TiO2. However the rutile (110) surface emerged to be most stable when 4.166% of Ti atoms were replaced by Nd atoms in the anatase slabs and then relaxed. To understand the anatase to rutile phase transformation with Nd doping, we have constructed Wulff shapes from the data of (004), (101), (110), and the (011) surface energies using the Wulff maker software.42,43 This software is based on the algorithm of the total surface energy minimization.34 The obtained Wulff shapes of equilibrium particles of the pure anatase (without Nd doping) and rutile (with 4.16 % Nd doping) are shown in Figure 13. Table II: Calculated surface energies with facet areas of TiO2 for the anatase and Nd doped (4.16%) rutile (004), (101), (110), and the (011) low-index directions using DFT calculations. TiO2 without Nd doping

TiO2 with 4.16% Nd doping

Anatase

Rutile


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Facet

Energy

Facet Area

Energy

Facet Area

J/m2

nm2

J/m2

nm2

(004)

0.40

7.20

0.79

2.0

(101)

0.35

133.60

0.59

49.60

(110)

0.81

0.0

0.30

250.0

From these equilibrium particles Wulff shapes, we calculated the areas of different surface facets, these values are given in Table II. In the anatase phase, the most dominant facets are (101) and (011). Upon Nd doping and random displacements (20% of the nearest neighbor distance) in the surface atoms in the anatase TiO2, these facets were reduced in their areas through the process of surface reconstruction, which happened during the slab relaxation and energy minimization.

The surface reconstruction process ultimately invoked the phase

transformation from the anatase to the rutile and (110) facet become the dominant phase over the other facets. To elaborate more on these findings, it can be said that the Wulff shapes were based on surface energies of the relaxed surfaces (with and without Nd dopant). It was found that the significant changes in the surface areas of different facets were correlated with the phase transformation as shown in Figure 13. It established the argument that Nd doping led to local strains in the TiO2 surfaces responsible for the phase transformation.

In short the Wulff

constructions described the Wulff shapes of two types of TiO2 nanoparticles in thermal equilibrium, i.e., one was TiO2 nanoparticle without Nd doping in the anatase phase and the second was Nd doped TiO2 nanoparticle in the rutile phase. The obtained shapes showed that the facet areas were different in the two phases and these differences in facet areas were indirectly complimented the experimental findings. Further, it was clearly observed in the Wulff shapes that the (101) plane was dominating in the anatase phase and (110) plane did not exist in the anatase phase, whereas (101) plane was not dominating in the Nd doped rutile, anymore, and (110) plane became dominant in the Nd doped rutile phase as shown in the Figure 13, which was exactly the findings of the experimental results. In the light of experimental results explained above, the role of the doped Nd was twofold, first introduced stress and second created O vacancies, which led to the phase transformation. The synthesis of TiO2 films in oxygen reduced 15 ACS Paragon Plus Environment


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environment was expected to have enhanced the rate of phase transformation and finally in amorphous phase. In present calculations, Nd was doped in the top layer, of course the surface energies may differ to some extent (but not significantly) if the Nd atom was doped in the 2nd or 3rd layers and if the slab was thick enough. In our calculations the thickness of the slab was sufficiently thick enough to conclude that the surface energies would not be significantly different if the atom is in the 2nd layer. The DFT calculations were kept simple and were performed for Nd doping, which confirmed the phase transformation. However, inclusion of oxygen vacancies would have enabled to determine the rate of phase transformation without affecting the outcome of present results.

IV.

CONCLUSION: Nd doped TiO2 films were grown by pulse laser deposition in varying O2 partial pressures

in a wide range from 100 mtorr (O2 rich) to 0 mtorr (O2 absent). Three concentrations of Nd were used as 1.0, 1.5 and 2.0 at.%. XRD and Raman spectroscopy results complimented each other and revealed that the phase transformation from the anatase to rutile to the metastable and finally to amorphous can be tailored by tuning the O2 content and Nd doping concentration. Reduction in the crystallite size was also observed by XRD and Raman analyses. SEM showed that the reduction in film thickness was related to the crystallite size reduction and phase transformation. XPS confirmed the reduction of Ti (Ti4+ to Ti3+), emergence of the O vacancies and consequently generation of non-stoichiometry in TiO2 lattice. So, the structure of TiO2 rearranged itself to a short range ordered structure. Thus, it is concluded that a proper control of oxygen content during the growth and suitable variation in the dopant concentrations can lead to the growth of any desired phase from anatase to amorphous in the PLD growth technique. The experimental findings were supported by the DFT calculations of surface energies for the

anatase and Nd doped rutile (004), (101), (110), and the (011) low-index planes. From the data of surface energies, the area of different facets was determined by constructing Wulff equilibrium particle shapes of the anatase and rutile phases. These equilibrium particle shapes of Wulff constructions clearly show the phase transformation upon Nd doping.

V.

ACKNOWLEDGMENTS: 16 ACS Paragon Plus Environment

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The research work was funded by Higher Education Commission (HEC) of Pakistan though NRPU # 1770 and IRSIP program and indigenous scholarship to A. Ali for his M.S. leading to Ph.D. studies (under pin No. 074-3560Ps4-134). VI.

REFERENCES:

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(24) Gamboa, J. A.; Pasquevich, D. M. Effect of Additives on Photocatalytic Activity of Titanium Dioxide Powders Synthesized by Thermal Plasma. J. Am. Ceram. Soc. 1992, 75, 2934-2938. (25) Syarif, D. G.; Miyashita, A.; Yamaki, T.; Sumita, T.; Choi, Y.; Itoh, H. Preparation of Anatase and Rutile Thin Films by Controlling Oxygen Partial Pressure. Appl. Surf. Sci. 2002, 193, 287-292. (26) Hanaor, D. A. H.; Sorrell, C. C. Review of the Anatase to Rutile Phase Transformation. J. Mater Sci. 2011, 46, 855-874. (27) Li, W.; Frenkel, A. I.; Woicik, J. C.; Ni, C.; Shah, S. Ismat. Dopant Location Identification in Nd3+-Doped TiO2 Nanoparticles Phy. Rev. B 2005, 72, 155315-155316. (28) Ali, A.; Yassitepe, E.; Ruzybayev, I. O.; Shah, Ismat S.; Bhatti, A. S. Improvement of (004) Texturing by Slow Growth of Nd Doped TiO2 Films. J. App. Phy. 2012, 112, 113505 (1-6). (29) Kresse, G.; Hafner, J. Ab Initio Molecular Dynamics for Liquid Metals. Phys. Rev. B. 1993, 47, 558-561. (30) Kresse, G.; Furthmuler, J. Efficiency of Ab-Initio Total Energy Calculations for Metals and Semiconductors Using A Plane-Wave Basis Set. Comput. Mater. Sci. 1996, 6, 15-50. (31) Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector Augmented-Wave Method. Phys. Rev. B 1999, 59, 1758-1775. (32) Perdew, John P.; Burke, Kieron; Ernzerhof, Matthias. Generalized Gradient Approximation Made Simple. Phys. Rev. Letters 1996, 77, 3865-3868. (33) Monkhorst, H. J.; Pack, J. D. Special Points for Brillouin-Zone Integrations. Phys. Rev. B. 1976, 13, 5188-5192. (34) Wulff, G.; Kristallogr, Z. Velocity of Growth and Dissolution of Crystal Faces. Mineral 1901, 34, 449-530. (35) Batzill, M.; Morales, E. H.; Diebold, U. Influence of Nitrogen Doping in the Defect Formation and Surface Properties of TiO2 Rutile and Anatase. Phys. Rev. Lett. 2006, 96, 26103-26107. (36) Craido, J.; Real, C. Mechanism of the Inhibiting Effect of Phosphate on the Anatase-Rutile Transformation Induced by Thermal and Mechanical Treatment of TiO2. J. Chem. Soc. Faraday Trans. 1983, 79, 2765-2771.



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(37) Holmberg, B. Disorder and Order in Solid Solutions of Oxygen in Alpha-Titanium. Acta Chem. Scand. 1962, 16, 1245-1250. (38) Rumaiz, Abdul K.; Ali, Bakhtyar; Ceylana, Abdullah; Boggs, M.; Beebe, T.; Shah, S. Ismat. Experimental Studies on Vacancy Induced Ferromagnetism in Undoped TiO2. Solid State Comm. 2007, 144, 334-338. (39) Rossella, F.; Galinetto, P.; Mozzati, M. C.; Malavasi, L.; Fernandez, Y. D.; Drera, G.; Sangaletti, L. TiO2 Thin Films for Spintronics Application: A Raman Study. J. Raman Spectrosc. 2009, 41, 558-565. (40) Tian, F.; Zhang, Y.; Zhang, J.; Pan, C. Raman Spectroscopy: A New Approach to Measure the Percentage of Anatase TiO2 Exposed (001) Facets. J. Phys. Chem. C 2012, 16, 75157519. (41) Campbell, I. H.; Fauchet, P. M. The Effects of Microcrystal Size and Shape on the One Phonon Raman Spectra of Crystalline Semiconductors. Solid State Commun. 1986, 58, 739741. (42) Karim, A.; Fosse, S.; Persson, K. A. Surface Structure and Equilibrium Particle Shape of the LiMn2O4 Spinel From First-Principles Calculations. Phys. Rev. B 2013, 87, 7532275326. (43) Zucker, R.; Chatain, V. D.; Dahmen, U.; Hagege, S.; Carter, W. C. New Software Tools for the Calculation and Display of Isolated and Attached Interfacial-Energy Minimizing Particle Shapes. J. Mat. Sci. 2012, 47, 8290-8302.


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Figure Captions: Figure 1: Plot of XRD patterns of Nd:TiO2 films to show the effect of variation of Nd concentration from 1 to 2 at.% at fixed O2 partial pressures (a) 100 mtorr, (b) 10 mtorr, (c) 1 mtorr and (d) 0 mtorr. Figure 2: Plot of XRD pattern of Nd:TiO2 films to show the effect of varying O2 partial pressures from 100 to 0 mtorr at fixed doping concentration of Nd (a) 1.0 at%, (b) 1.5 at% and (c) 2.0 at%. Figure 3: Plot of the crystallite sizes of different growth planes of TiO2 calculated by Scherrer’s equation with varying Nd doping concentration. Solid lines are used to guide the eye and to highlight the drop in crystallite size for the each plane. Figure 4: The cross-sectional SEM micrographs of 1 at.% Nd doped TiO2 thin films at different oxygen partial pressures (a) 100mtorr, (b) 10mtorr, (c) 1mtorr and (d) 0mtorr. Figure 5: The plot of TiO2 film thickness with the variation of O2 partial pressure at fixed 1 at.%.Nd doping concentration. Figure 6: High resolution XPS spectra of the Ti-2p region obtained from the 1.5 at.% Nd doped TiO2 thin films grown at different O2 partial pressure from 100 mtorr to 0 mtorr. The bottom spectrum also displays the Gaussian fitting used to determine the contribution of Ti4+ and T3+. Figure 7: High resolution XPS spectra of the Ti-2p region for TiO2 thin films grown at 0 mtorr O2 partial pressure with varying concentration of Nd (1.0, 1.5 and 2.0 at.%). The low and high binding energy labeled peaks is due to Ti-2p3/2 and Ti-2p1/2, respectively. Peak fitting shows the presence of multiple states (Ti4+ and Ti3+) of Ti. Figure 8: The plot of the variation of integrated intensities of Ti-2p1/2 (left) and Ti-2p3/2 (right) with increasing concentration from 1 to 2 at.% Nd in TiO2 films grown at 0 mtorr O2 partial pressure.



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Figure 9: High resolution XPS spectra of O-1s region of TiO2 thin films grown at fixed 1.5 at.% Nd doping concentration and varying O2 partial pressure from 100 mtorr to 0 mtorr. Peak fitting represents the contribution of low and high binding energy peaks due to O vacancies and O attached to Ti, respectively. Figure 10: The plot of the variation of integrated intensities of HBE (left) and LBE (right) of O1s region with increasing concentration of Nd doping in TiO2 films grown at 0 mtorr. Figure 11: Raman spectra of Nd:TiO2 films to show the effect of variation of Nd concentration from 1 to 2 at.% at fixed O2 partial pressures (a) 100 mtorr, (b) 10 mtorr, (c) 1 mtorr and (d) 0 mtorr. Figure 12: The Raman spectra of Nd:TiO2 films to show the effect of varying O2 partial pressures from 100 to 0 mtorr at fixed doping concentration of Nd (a) 1.0 at%, (b) 1.5 at% and (c) 2.0 at%. Figure 13: The TiO2 anatase and Nd doped rutile faceted equilibrium particle shapes based on the Wulff construction


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Figure 1:



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Figure 2:


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Figure 3: 60

Crystallite Size (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

101 004 110 ms

45

30

15

0

1.0

1.5

Nd at. % in TiO2


2.0


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Figure 4: (a) 100 mtorr

(b) 10 mtorr

(c) 1 mtorr

(d) 0 mtorr

4.9

1.0

0.7

0.38


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Figure 5:

10000

Film Thickness (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

1000

100 0.1

1

10

O2 partial presure (mtorr)


100


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Figure 6:


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Figure 8:


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Figure 9: O1s

1.5 At. % of Nd

100 mtorr

Intensity (arb. u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

10 mtorr

1 mtorr

LBE 0 mtorr

HBE

534

531

528

Binding Energy (eV)


525


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Figure 10:


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Figure 11:



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Figure 12:


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Figure 13:



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Phase Transformations in the Pulsed Laser Deposition Grown TiO2

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