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Investigating Optical Properties of Atomic Layer Deposited ZnO/TiO Multi-Stacked Thin Films Above Mott Critical Density x

Debabrata Saha, Pankaj Misra, Mukesh Prenbaratab Joshi, and Lalit Mohan Kukreja J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b05056 • Publication Date (Web): 24 Jul 2017 Downloaded from http://pubs.acs.org on July 28, 2017

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The Journal of Physical Chemistry

Investigating Optical Properties of Atomic Layer Deposited ZnO/TiOx Multi-stacked Thin Films Above Mott Critical Density Debabrata Saha*†, P. Misra*, M. P. Joshi, and L. M. Kukreja Laser Materials Processing Division, Raja Ramanna Centre for Advanced Technology, Indore 452 013, India

ABSTRACT: The evolution of the optical properties with dopant concentration has been investigated for a series of ZnO/TiOx multi-stacked layers having electron density exceeding the Mott critical limit of insulator-to-metal transition. These films were grown by vertically stacking multiple ZnO/TiOx bi-layers on (0001) sapphire substrates using atomic layer deposition. The films in the sparsely doped regime showed room temperature UV photoluminescence (PL) while being transparent and heavily degenerate in nature. Optical absorption spectra of these films did not exhibit any feature of excitonic resonance, indicating a possible excitonic Mott transition in the metallic limit. The low-temperature PL spectra also support this observation which shows line-shape characteristics typical for band-to-band emission. The sharp cut-off of the PL emission at the high energy edge corresponds to the Fermi level position inside the conduction band. In contrast, the broad low-energy wing is determined by the combined density of states washed out by potential fluctuations induced band tailing effect. A systematic blue shift of the

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high energy PL edge with increasing carrier density resembles the effect of band filling which has also been considered in explaining the optical absorption spectra of the films. The results of this study demonstrates that the multi-stacked dopant incorporation scheme in ALD could be highly useful to increase carrier concentration while minimizing disorder strength in the lattice which eventually results in high optical quality ZnO films with tunable electrical conductivity.

INTRODUCTION Investigating the electrical and optical properties of electronically doped ZnO thin films in the heavily doped regime is crucial for their promising applications in transparent electronics, optoelectronics and plasmonic based devices.1–6 In these semiconductors, free electron density ( n ) lies well above the Mott critical limit ( nc ) required for the onset of insulator-to-metal

transition (IMT).7,8 The value of nc in case of ZnO, as obtained from Mott criterion:

nc1/3aB = 0.25 , is found to be ~ 1x1019 cm-3.7,8 Such high carrier density can be achieved by incorporating intrinsic (oxygen vacancies and zinc interstitials) and/or extrinsic dopants (usually group III elements: Al, Ga, In etc.) in ZnO.1,3–5,8 However, dopant atoms except donating free carriers also introduce static-disorder due to their random incorporation in the host lattice which exhibits profound effects on the electrical as well as on the optical properties of the films. Dopant induced positional and energetic disorder gives rise to band tails and localized states within the band-gap which deteriorates the room temperature excitonic emission of the films.9,10 In most of the available reports on the photoluminescence (PL) of heavily doped ZnO based transparent conducting oxides (TCOs), detectable UV PL emission is reported only at lower temperatures ~ 5 K.11,12 Therefore, fabrication of dilutely doped ZnO films with enhanced


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doping-efficiency (i.e., high density of electronically active dopants) is crucial which may preserve the room temperature excitonic UV PL emission of the films while being transparent and conducting. In order to achieve such multi-functional semiconductors, controlled incorporation of dopants by inducing minimal disorder in the host lattice is extremely important which can be accomplished by employing atomic layer deposition (ALD) technique. Thin film growth in ALD proceeds through complementary and self-limiting surface chemical reactions which allow exactly the same amount of material deposition in each cycle.13 Hence, chemical compositions of multi-component materials can be controlled very accurately by simply varying the cycle ratio of the precursors involved in the ALD process. However, achieving higher doping-efficiency in ALD grown ternary metal-oxides is found to be rather difficult which can be attributed to the dense mono-layer incorporation of dopants in each cycle.3,4,14 Owing to the high dopant density on the doping plane, a significant fraction of the dopant atoms introduces static-disorder by forming electronically inactive dopant complexes which eventually results in poor doping-efficiency. The bottleneck of lower doping-efficiency in ALD grown films can be alleviated by following different dopant incorporation schemes and/or employing larger ligand size precursor molecules in the ALD process.4,14,15 In a very recent report, we have demonstrated that incorporation of sub-monolayer of group IV element Ti in ZnO can result in dopingefficiency as high as ~130% which is significantly enhanced compared to the Al doped ZnO films (~ 42 %).5 Moreover, extreme level of tunability in degree of static-disorder and carrier transport properties has also been achieved by continuously varying ZnO spacer layer thickness in the ZnO/TiOx multi-stacked structures.5 Therefore, it would be worthy enough to investigate the optical properties of ZnO/TiOx films in the lower disorder limit in search for room temperature UV light emitting TCOs. In the literature, ALD has been extensively used to grow


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electronically doped ZnO films for TCO application.3,4,15 The optical and structural properties of ZnO thin films and nanostructures grown by different deposition techniques have also been investigated by several research groups.16–19 Geng et al. have studied the room temperature PL properties of Al doped ZnO films grown by ALD.16 They have observed the deterioration of the UV PL intensity with increasing Al doping concentration. However, a detailed analysis on the PL characterization has not been presented in the paper. Chaaya et al. have also carried out room temperature PL measurements on ALD grown ZnO/Al2O3 nanolaminates.19 They have observed excitonic band edge PL along with the defect states originated PL emission in the visible spectral range. The difference in the UV/Vis PL intensity of the nanolaminates has been attributed to the structural change of the films due to the variation of the ZnO/Al2O3 bilayer thickness. However, the role of free carrier concentration on the PL and optical absorption characteristics of the films has not been clearly pointed out by the authors. To the best of our knowledge, there is so far not a single report on the detailed investigation on the temperature dependent UV PL characterization of degenerately doped ZnO films grown by ALD. In the present study, we have investigated optical properties of ZnO/TiOx multi-stacked thin films with Ti concentration in the range of ~ 0 to 10.24 at.%. The sparsely doped films having Ti concentration of ~ 0 to 0.78 at% were found to be heavily degenerate in nature which also exhibited high optical transmittance in the visible spectral range and intense UV PL emission at room temperature. Further insertion of Ti resulted in the deterioration of UV PL emission while preserving optical transparency and high electrical conductivity of the films. A detailed analysis of the optical transmittance and temperature dependent (~ 5-300 K) PL characterizations were performed to gain fundamental insights into the intricate interplay of free carrier density and dopant induced potential fluctuations on the optical properties of these films. The present work


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opens up the possibility to further extend the capabilities of thermal ALD to grow high optical quality ZnO films with tunable electrical conductivity which may find potential applications in transparent electronics, optoelectronics and plasmonic based devices. MATERIALS AND METHODS ZnO/TiOx multi-stacked films were grown on epi-polished single crystal (0001) sapphire substrates at ~ 2000C using a thermal ALD reactor (Model: TFS-200, Make: BENEQ Oy, Finland).5 The thicknesses (d) of the films as measured by spectroscopic ellipsometry were found to be in the range of ~ 180-220 nm. Diethylzinc (Zn(C2H5)2), deionised water (H2O) and titanium tetrachloride (TiCl4) were used as precursors for Zn, oxygen and Ti respectively. Pulsing time was kept at 200 ms for both Zn(C2H5)2 and H2O, whereas TiCl4 was pulsed into the reactor for a very short time of 25 ms to provide sub-saturating exposure on the growing ZnO film surface.5 N2 purging time was kept constant for 1s throughout the deposition process. A schematic representation of the pulsing and purging sequence has been given in ref. 16. Ti concentration in the films was varied systematically in a wide range from ~ 0 to 10.24 at. % which resulted in the variation of electron density from ~ 4.2x1019 cm-3 (intrinsic ZnO) to as high as ~ 3.8x1020 cm-3.5 ALD grown films are named as S (n, 1: s) in which ‘n’ cycles of ZnO followed by ‘1’ cycle of TiOx makes one complete supercycle, and ‘s’ represents total number of super-cycles in a film (see Table 1). The number of super-cycles ‘s’ is varied from 4 (S1) to 167 (S10) to achieve total number of ALD cycles around 1000 for all the films. A more detailed description on the growth of ZnO/TiOx multi-stacked layers can be found in our earlier reports.5,20,21 In the present paper, we are reporting the optical properties of the ZnO/TiOx samples on which electrical measurements have already been carried out in ref. 5. As can be seen in ref. 5 (Figure 1 (c) and Table 1), ZnO/TiOx samples are divided in two zones, namely zone-I


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which comprising of samples with Ti concentration 0-0.78 at.% and zone-II beyond 0.78 at.(%) of Ti. The samples in zone-I are in the lower disorder limit, whereas disorder strength is found to be higher in zone-II.5 Room temperature PL measurements were carried out using a 20 mW HeCd laser operating at 325 nm as an above band-gap excitation source. The samples were kept in a closed cycle optical cryostat (Oxford Instruments) for temperature dependent (~ 5-300 K) PL measurements. PL signal was dispersed in a spectrometer (Triax 550, Jobin Yvon, France) and detected by a UV sensitive CCD detector (Andor, UK). RESULTS AND DISCUSSION Figure 1 shows room temperature PL spectra of all the ZnO/TiOx multi-stacked layers with Ti concentration varying from 0-10.24 at.%. The films in zone-I (please refer to the experimental section above) having Ti concentration in the range of 0 (ZnO)-0.78 at.% (S5) exhibited detectable UV PL emission at room temperature. With further insertion of Ti i.e., in zone-II, UV PL intensity was found to be completely diminished for the films S6-S10. The inset of Figure 1 shows integrated UV PL intensity as a function of electron concentration. As can be seen, integrated PL intensity remained almost unaffected with increasing Ti incorporation up to ~ 0.4 at.% (S3). However, in this regime carrier density was found to be significantly increased from ~ 4.2x1019 cm-3 (intrinsic ZnO) to ~ 2.5x1020 cm-3 (S3). Thus n-type conductivity of the films were increased without deteriorating room temperature UV PL emission. Such multi-stack incorporation of sub-monolayer dopants in ALD is indeed useful to reduce dopant clustering which acts as luminescence killers in heavily doped ZnO films. However, beyond S3 i.e., for the films S4 and S5, integrated PL intensity was found to be decreased which could be attributed to the increase of non-radiative recombination processes in the films. With further incorporation of Ti i.e., in zone-II (S6-S10), UV PL intensity was found to be completely diminished which is


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plausibly due to the drastic decrement of doping-efficiency (i.e., formation of compensative native defects and electronically inactive dopant complexes) and enhanced degree of staticdisorder in the films.5 Increased strength of disorder in zone-II (S6-S10) could also be realized from the underlying electron transport properties of these films which exhibited a crossover from good metallic to incipient non-metallic resistivity behaviour in transition from zone-I (S1-S5) to zone-II (S6-S10).5 Therefore, in the present paper we have restricted our discussions on the optical properties of the films in zone-I. Figure 2 shows room temperature PL and optical absorption spectra of the films in zone-I which showed detectable UV PL emission at room temperature. All the films were found to be optically transparent (~ 80 %) in the visible spectral range (Supporting information, Figure S1). Optical absorption edge of the films showed a clear blue shift with increasing Ti concentration due to the well-known Burstein-Moss (BM) band-filling effect.7,22,23 As carrier concentrations in these films are well above the Mott critical density, Fermi energy EF lies inside the conduction band i.e., free electrons occupy states at the bottom of the conduction band. Hence, Pauli blocking prevents optical transitions into these states and optical-absorption onset occurs at higher energies. On the contrary, in case of PL emission, when the excited holes percolate from

k = kF to k = 0 , they can recombine at any point along the way, since the states directly above them (in k space) are always filled with electrons. Thus UV PL peak of each of the degenerately doped film was found to be red shifted (Stokes shift) compared to its optical band gap. The optical band gap as a function of carrier concentration is shown in Figure 3. According to BM effect, the increment in optical band gap compared to that of undoped ZnO ( E g0 ) is given by ∆EBM

h 2k F2 h 2k F2 h 2k F2 = + = * 2mc* 2mv* 2mvc

(1)


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where mvc* = (1 / mc* + 1 / mv* ) is the reduced effective mass with mc* and mv* are the effective −1

masses of carriers in the conduction and valance bands, respectively.7,22,23 The expression for the Fermi wave vector is given by k F = ( 3π 2 ne ) where ne is the electron density. Therefore, Eq. 1 1/3

becomes ∆EBM

h 2 (3π 2ne )2/3 = 2mvc*

(2)

As can be seen from Figure 3, measured band-gap values were significantly lower compared to that predicted by Eq. 2. This inconsistency could be attributed to the band-gap narrowing (BGN) effect.22,24 At electron concentration higher than Mott’s critical value of insulator-to-metal transition, the modification of electronic states begins to appear in the crystal because of the correlated motion (many body interactions) of conduction electrons and their scattering against ionized impurities. In the present study, all the ZnO/TiOx films are on the metallic side (i.e., σ≠0 at T→0 K) of the insulator-to-metal transition as confirmed by the temperature dependent electrical resistivity measurements of the films in our earlier report.5 Therefore BGN, which is dominated by two contributions, namely electron-electron interaction ( ∆Eee ) and electron-ion interaction ( ∆Eei ), is expected to have a crucial role on the optical properties of these films. The total contribution of BGN on the optical band gap is given by

∆EBGN = ∆Eee + ∆Eei

(3)

The electron-electron interaction is given by22,24

∆Eee = −

e2k F 2π ε sε 0 2

−

e 2kTF 8πε sε 0

 4  kF   1 − arctan    kTF    π

(4)

and electron-ion interaction is given by22,24


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∆Eei = −

e 2 ne ε 0 ε s aB* ( kTF )3

(5)

In the above equations k F is the Fermi wave vector and kTF = 2 k F / π a B* is the Thomas-Fermi wave vector. This negative effect of BGN is competitive with the positive BM shift. Therefore, considering the combined effects of BM shift (Eq. 2) and BGN (Eq. 3), the optical band gap values of these TCO films is given by22,24 E g = E g0 + ∆E BM + ∆E BGN

(6)

Eq. 6 provides fairly good theoretical predictions of the measured band gap values (see Figure 3) which indicates that at n > N c , free carrier density and many body interactions have significant role on the optical absorption spectra of the ZnO/TiOx films. In zone-II, electron concentration decreases with increasing Ti (at.%) which has already been reported in our earlier paper.5 However, band gap values for these films are found to be monotonically increased with increasing Ti incorporation (Supporting information, Figure S1). This is in sharp contrast to zone-I and therefore, variation of optical band-gap for these films could not be explained by considering only equation 6. In these films, conduction band minimum (CBM) energy position is possibly changed by the incorporation of Ti dopants i.e., it modulates the fundamental band gap of ZnO.25,26 The significant deterioration of the crystalline structure of ZnO is indeed confirmed by the X-ray amorphous nature of the samples beyond S6. Moreover, sample S10 with the highest Ti concentration of ~ 10.24 at.(%) did not exhibit the characteristic fundamental optical absorption of ZnO, further validating structural deterioration of the film. Therefore, a detailed study on the optical properties in zone-II requires ultra violet photoelectron spectroscopy (UPS) measurement and first principle calculation of the band structures of the films which is beyond the current scope of the manuscript.


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Dopant atoms except donating free carriers also introduce static-disorder due to their random distribution in the host lattice. In our earlier reports, the effects of varying degree of disorder on the underlying electron transport properties of the ZnO/TiOx layers has been addressed.5 Disorder also exhibits profound effects on the optical absorption spectra of these films as shown in Figure 2. The positional and energetic disorder as induced by the presence of native defects and incorporated Ti impurities results in unavoidable potential fluctuations in a microscopic scale which gives rise to band tails within the forbidden gap.9,10 The density of localized states in the band tails exponentially decreases with increasing energy from the band edge which results in an optical absorption coefficient α = α 0 exp(hν − E1 / EU ) in the sub-band gap region (

hν < Eg ) as shown in Figure 4.27 The above equation is known as Urbach relation where α0 and

E1 are constants and EU is the Urbach energy. The value of EU has been obtained by linear fitting of the ln (α ) vs hν plot in the sub-band gap region.27 The monotonic increment of EU with increasing Ti concentration (see Table 1) could be attributed to the increase of degree of static-disorder in the films. We are now interested to examine the PL characteristics of these degenerately doped semiconducting films. Usually, room temperature PL of high quality intrinsic ZnO films exhibit strong excitonic emission due to its large free exciton binding energy of ~ 60 meV.22 However, in presence of free carriers, the Coulomb interaction between electron-hole pairs in an exciton is screened which reduces the binding energy.28 When free carrier density reaches Mott’s critical limit, binding energy approaches to zero and dissociation of excitons (i.e. formation of electronhole plasma) occur. This phenomenon is usually referred to as the excitonic Mott transition which has been investigated in variety of materials including degenerately n-type doped semiconductors (InN, GaN etc), optically excited semiconductors etc.28,29 In the present case, the


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room temperature optical absorption spectra of the films did not exhibit any feature of excitonic resonance below the fundamental band-gap which indicates plausible excitonic Mott transition in the ZnO/TiOx layers. However, in a very recent report, Schleife et al. have carried out ab initio study on the frequency dependent absorption coefficient of intrinsic and n-type doped ZnO films.28 Results of this study indicates that above Mott critical limit, excitonic bound states can persist which may also influence the optical absorption even when the binding energies are negligibly small. However, experimental observation of such bound states as distinct peaks in the optical absorption spectra is quite difficult due to lifetime, temperature, and instrumental broadening effects. Therefore, in order to further investigate whether an excitonic Mott transition truly occurs in these films, we have carried out temperature dependent (~ 5-300 K) PL measurements. Figure 5 shows temperature dependent PL spectra for all the films in zone-I which showed detectable UV PL emission at room temperature. PL line shape of all the films were found to be symmetric Gaussian type at room temperature (~ 300 K) due to excitation of free electrons near the Fermi level and additional thermal broadening effects. However, with decreasing temperature, spectral shape of the PL emission gradually becomes asymmetric. The 5 K PL spectra of all the films (see Figure 6) exhibited a steep slope at the high energy side and a smooth slope at the low-energy side as typically observed in case of degenerately doped semiconductors.29–32 At low enough temperature, the high-energy edge represents the sharp cutoff of the Fermi level inside the conduction band and the broad low-energy wing is determined by the combined density of states washed out by the band tailing effect which arises due to dopant induced potential fluctuations. The high energy cut-off of the PL spectra corresponds to the Fermi energy position EF inside the conduction band as calculated with respect to the valence


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band maximum i.e., considering only the conduction band dispersion in Eq. 1.29 Therefore, the steep slope in the low-temperature PL line shape is caused by the decrease in free electron population above EF. Moreover, the systematic blue shift of the high energy PL edge (see Figure 6) with increasing carrier concentration (i.e., increasing EF) resembles the Burstein-Moss band filling (or Pauli blocking of optical transitions) effect which has also been considered in explaining the optical absorption spectra of the films in Figure 3. Such PL characteristics indicates band-to-band (B-B) transitions as the mechanism of PL emission in the ZnO/TiOx films. It is important to mention here that blue-shift in the PL spectra can also be induced by the strain in the lattice.24 In case of hetero-epitaxial growth, thin films are always subject to (a) residual biaxial strain due to the lattice parameter mismatch between the substrate and the deposited layer and (b) thermal mismatch strain as induced by the difference in the thermal expansion coefficients of the substrate and the grown material. In order to investigate the effect of strain on the PL characteristic of the as grown ZnO/TiOx films, we have carried out XRD measurements which showed growth of polycrystalline films with (10.0), (00.2) and (10.1) peaks corresponding to hexagonal wurtzite structure of ZnO (Supporting information Figure S2).8 The (00.2) peak position of the undoped ZnO film appears at 2θ = 34.420 which is close to the value 34.400 for the bulk ZnO (JCPDS card no. 79-2205).33 With Ti incorporation, (00.2) peak position did not show any significant shift in its peak position. However, XRD peaks are broadened and peak intensity is also found to be significantly reduced for the film S6 with Ti concentration of ~ 1.7 at.(%). With further incorporation of Ti, films become amorphous in nature. Thus XRD measurement confirm that the blue-shift in the high energy PL edge of the ZnO/TiOx films is not originated due to strain in the lattice rather increase of the Fermi level position inside the conduction band with increasing electron density might have results in the observed PL edge


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shift. In the literature, B-B PL emission is reported in a variety of heavily n-type degenerate semiconductors including GaN, InN, InGaP, GaAs etc.24,29–32 Therefore, in addition to the optical absorption, low temperature PL measurements further confirms the dissociation of excitons above Mott critical density which eventually triggers B-B transitions in the PL spectra of the ALD grown ZnO/TiOx films. The evolution of FWHM of the room temperature PL spectra with electron density could not be fitted by considering only the effect of dopant induced potential fluctuations. This further implies significant role of band filling and band gap narrowing on the PL line broadening above Mott critical density.34 In contrast to PL, both conduction and valence band dispersions are taken into consideration in explaining the optical absorption spectra of the films (see Eq. 1). This implies that above Mott critical density, non-k preserving (i.e., ∆k≠0) indirect transitions between free electrons in the conduction band and localized tail states above the valence band prevails in the radiative recombination process.29 The increased valence band tailing plausibly mediates between k F in the conduction band and the Γ point in the valence band. Moreover, it has also been argued that strong ionized impurity scattering in doped semiconductors can results in the breakdown of the momentum conservation in the PL emission process which gives rise to non k-preserving radiative recombinations.30,32 In addition to the band-to-band emission, a well resolved strong PL peak at ~ 3.31 eV was also observed for the films S1-S4. In the literature, PL emission from ZnO at ~ 3.31 eV is highly debated.11,12,35–37 It has been controversially ascribed to donor-acceptor pair (DAP) transitions or due to basal plane stacking faults (BSFs).11,12,35–37 DAP transitions are originated due to the recombination between neutral donor and acceptor pairs which restricts its observation only at lower temperatures.36,38 However, above Mott critical density (nc), donors are ionized even at T→0 K (i.e., zero ionization energy or metallic state) which consequently eliminate DAP


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recombination processes. In the present case, all the samples have carrier density n > nc and temperature dependent resistivity measurements as reported in our earlier paper also showed metallic behaviour (i.e., σ≠0 at T →0 K).5 Therefore, DAP transitions can be completely ruled out as the origin of 3.31 eV PL emission. In order to further elucidate the nature of this emission peak, we have carried out intensity dependent PL measurements for all the films at 5 K. As can be seen from Figure 7, the intensity of the PL emission for the film S1 increased with increasing excitation power. Interestingly, PL peak position remained unaffected to the excitation intensity. This is in contrast to the DAP recombination which exhibits a blue shift of the PL peak position with the increase of incident laser beam intensity.11,12,35 These observations eliminate DAP transitions as the mechanism of the 3.31 eV PL emission. Recently, Yang et al. have also reported observation of 3.31 eV PL emission in ALD grown ZnO films.37 They have attributed this characteristic PL peak to the basal plane stacking faults (BSFs) which are found to be the dominant structural defects in ALD grown epitaxial ZnO films on (0001) sapphire substrates. BSF in wurtzite lattice can be considered as a thin insertion layer of cubic zinc blend structure which is sandwiched by the wurtzite barriers. BSF forms a quantum well type structure for electrons in the conduction band and creates a potential barrier for holes in the valence band. Therefore, BSF emission plausibly resulted from the indirect recombination of confined electronhole pairs in which electrons are localized in the potential well of the BSF and holes at the interface of the BSF and wurtzite structure. The films S4 and S5 did not exhibit distinct BSF emission peaks which could be due to the large spectral broadening and weak PL intensity at higher dopant concentrations.


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CONCLUSION In summary, the influence of degenerate electron density on the optical properties of ALD grown ZnO/TiOx multi-stacked thin films has been experimentally investigated using a combined study of optical absorption and temperature dependent (5-300 K) photoluminescence measurements. Free carrier induced exciton screening effect in these films plausibly results in Mott transition i.e., exciton dissociation which results in non k-preserving band-to-band PL emission as evident from the blue-shift of the high energy PL edge at 5 K. The intensity dependent PL measurements have also been carried out at 5 K which ruled out the possibility of DAP transition as the mechanism of 3.31 eV PL emission. The optical band-gap values of these films are quantitatively analyzed by considering the combinatorial effects of Burstein-Moss band-filling (or Pauli blocking) and many body interactions induced band-gap narrowing effects. Moreover, the blue-shift in the optical absorption edge is found to be accompanied with potential fluctuations induced band-tailing effect as manifested by the increase in Urbach energy from ~ 54 meV for un-doped ZnO to ~ 271 meV for Ti concentration of ~ 6.42 at. (%). ASSOCIATED CONTENT Supporting Information The following files are available free of charge. Optical transmittance spectra of all the ZnO/TiOx films (PDF), X-ray diffraction spectra of all the ALD grown ZnO/TiOx films with varying Ti concentration (PDF).

AUTHOR INFORMATION *Authors for correspondence


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Email: [email protected] (Debabrata Saha) & [email protected] (P. Misra); Ph. No.: +91731-2488374, Fax. No.: +91-731-2488330. Present Addresses †

Department of Energy Science and Engineering, Indian Institute of Technology Bombay, Powai

400 076 Mumbai India. Author Contributions The manuscript was written by Debabrata Saha through discussions with all the co-authors. All authors have given approval to the final version of the manuscript. ACKNOWLEDGEMENTS Authors would like to thank Mr. V. K. Sahu of Laser Materials Processing Division, Raja Ramanna Centre for Advanced Technology, Indore for his help in temperature dependent photoluminescence measurements of some of the samples. Dr. Debabrata Saha sincerely acknowledge Prof. Dr. Martin Feneberg of Institut für Experimentelle Physik, Abteilung Materialphysik, Otto-von-Guericke Universität Magdeburg for his constructive suggestions and fruitful discussions on the temperature dependent photoluminescence characterizations. Debabrata Saha acknowledges HBNI, RRCAT (Indore) for providing financial support during the tenure of the research work.


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Table 1. Ti (at. %) and Urbach energy (EU) values for all the ZnO/TiOx films S1-S10 grown by ALD.

Sample name

Ti (at. %)

EU (meV)

ZnO

0.00

53.93±2.55

S1 (250, 1:4)

0.20±0.02

56.36±2.80

S2 (150, 1: 7)

0.27±0.02

74.46±1.87

S3 (100, 1:10)

0.41±0.03

87.56±2.71

S4 (60, 1: 16)

0.65±0.02

147.05±2.34

S5 (50, 1: 20)

0.78±0.02

170.06±2.62

S6 (30, 1: 33)

1.70±0.02

175.43±2.40

S7 (20, 1: 48)

2.04±0.03

184.84±2.03

S8 (15, 1: 63)

4.10±0.03

232.55±1.61

S9 (10, 1: 100)

6.42±0.02

271.73±2.48

S10 (5, 1: 167)

10.24±0.02

--

S ( n,1: s)


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Figure Captions Figure 1. Room temperature PL spectra for all the films in the Ti concentration range of ~ 010.24 at.%. The inset shows variation of integrated PL intensity and PL peak position as a function of carrier density. Figure 2. Room temperature PL spectra along with optical absorption for all the films in zone-I (ZnO and S1-S5). Vertical arrows indicates the point of inflection in the absorption spectrum of the respective films. Figure 3. Variation of measured optical band-gap values with carrier density for the films in zone-I. Theoretical predictions of band-gap values by BM (eq. 2), BGN (eq. 3) and their combined effects (eq. 6) are also shown. Figure 4. Variation of ln(α) as a function of hν in the sub-band gap regime of the ZnO/TiOx films. The solid lines show linear fit of the experimental data. Figure 5. Temperature dependent PL spectra for all the films in zone-I (ZnO and S1-S5) in the temperature range of 5-300 K. B-B and BSF indicates band-to-band and basal plane stacking faults related PL emissions respectively. Figure 6. PL spectra for the ZnO/TiOx films in zone-I at 5 K. PL spectra are normalized in intensity to clearly illustrate any shift in the PL peak position with Ti incorporation. Figure 7. Variation of 5 K PL spectrum of the film S1 with increasing incident laser beam intensity.


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TOC Graphic


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Figure 1. Room temperature PL spectra for all the films in the Ti concentration range of ~ 0-10.24 at.%. The inset shows variation of integrated PL intensity and PL peak position as a function of carrier density. 203x143mm (300 x 300 DPI)


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Figure 2. Room temperature PL spectra along with optical absorption for all the films in zone-I (ZnO and S1S5). Vertical arrows indicates the point of inflection in the absorption spectrum of the respective films. 279x361mm (300 x 300 DPI)



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

Figure 3. Variation of measured optical band-gap values with carrier density for the films in zone-I. Theoretical predictions of band-gap values by BM (eq. 2), BGN (eq. 3) and their combined effects (eq. 6) are also shown. 215x166mm (300 x 300 DPI)


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Figure 4. Variation of ln(α) as a function of hν in the sub-band gap regime of the ZnO/TiOx films. The solid lines show linear fit of the experimental data. 215x166mm (300 x 300 DPI)



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

Figure 5: Temperature dependent PL spectra for all the films in zone-I (ZnO and S1-S5) in the temperature range of 5-300 K. B-B and BSF indicates band-to-band and basal plane stacking faults related PL emissions respectively. 338x190mm (96 x 96 DPI)


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Figure 6: PL spectra for the ZnO/TiOx films in zone-I at 5 K. PL spectra are normalized in intensity to clearly illustrate any shift in the PL peak position with Ti incorporation.

203x143mm (300 x 300 DPI)



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Figure 7: Variation of 5 K PL spectrum of the film S1 with increasing incident laser beam intensity.

202x141mm (300 x 300 DPI)


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TiOx

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