Highly crystalline p-PbS thin films with tunable optical and hole transport parameters by chemical bath deposition

Accepted Manuscript Highly crystalline p-PbS thin films with tunable optical and hole transport parameters by chemical bath deposition

Rekha Bai, Dinesh Kumar, Sujeet Chaudhary, Dinesh K. Pandya PII:

S1359-6454(17)30255-0

DOI:

10.1016/j.actamat.2017.03.062

Reference:

AM 13668

To appear in:

Acta Materialia

Received Date:

05 January 2017

Revised Date:

21 March 2017

Accepted Date:

25 March 2017

Please cite this article as: Rekha Bai, Dinesh Kumar, Sujeet Chaudhary, Dinesh K. Pandya, Highly crystalline p-PbS thin films with tunable optical and hole transport parameters by chemical bath deposition, Acta Materialia (2017), doi: 10.1016/j.actamat.2017.03.062

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Highly crystalline p-PbS thin films with tunable optical and hole transport parameters by chemical bath deposition Rekha Bai1, Dinesh Kumar1, Sujeet Chaudhary1, and Dinesh K. Pandya*,1 Thin Film Laboratory, Physics Department, Indian Institute of Technology Delhi, New Delhi 110016, India *E-mail address: [email protected]

Abstract Lead sulfide (PbS) thin films, consisting of well faceted (up to 400 nm) cubic-nanocrystals and possessing significantly improved opto-electronic parameters essential for photovoltaic applications, are grown by utilizing chemical bath deposition (CBD) technique with bath concentrations of 10–200 mM. X-ray diffraction (XRD) and Raman studies confirm the highly crystalline and pure phase of PbS. FESEM and HRTEM studies show that all the films possess uniform and compact (111) oriented nanocubic morphology. Bath concentration change provides tunability of nanocube size from 100–400 nm and the direct optical band gap from 1.50–0.94 eV. The PbS films exhibit p-type semiconducting behavior with hitherto unreported concurrent highest mobility of 29.3 cm2V–1s–1 and high carrier concentration of 1018 cm–3 with the lowest room temperature resistivity of 0.26 Ω−cm. The 25 mM and 10 mM films show significant surface plasmon absorption in 1200–2400 nm range making them suitable as efficient infrared absorbers in excitonic and multi-junction solar cells. Keywords: Lead sulfide thin films, Chemical Bath Deposition, IR absorption flux, Electrical conductivity, Excitonic and plasmonic solar cells

1Thin

Film Laboratory, Physics Department, Indian Institute of Technology Delhi, New Delhi 110016, India.

*E-mail: [email protected]; Tel: +91-11-2659 1347

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1.

Introduction Metal sulfide based binary nanostructured semiconductors have been attracting great

attention as an efficient material for solar energy conversion owing to their versatile optical and electrical properties [1–4]. These metal sulfide nanostructures comprise of a class of architecture that possess high absorption coefficient, good photo-stability, broad excitation spectra, capability to deal with hot electrons and generate multiple charge carriers per photon, as well as provide possibility of size dependent tunable band gap, which are favorable for thin film solar cells [1–2,5–8]. Thin films of n-type semiconductors such as ZnO, TiO2, CdS, CdSe etc. are reported to be the alternative environment friendly inorganic materials to achieve promising junctions with the p-type materials such as PbS, PbSe, CIGS etc. for application in excitonic and multi-junction thin film solar cells [2,4,9–13]. Among all of them, PbS (composed of earth abundant elements) has emerged as a promising candidate in photovoltaic applications as a relevant p-type absorber and hole carrier material due to its fundamental narrow band gap of 0.41 eV (that can be enhanced on quantum confinement allowing tunability over a much wider range from visible to infrared), high absorption coefficient of ~105 cm–1 in visible and near-infrared (NIR), high dielectric constant (ɛPbS = 15–20), small effective mass for the electrons and holes, and a large exciton Bohr radius ~18 nm, which allows the strong quantum confinement effects even for larger particles [1–2,14–16]. Among various wide band gap metal oxides, ZnO and TiO2 are considered as a suitable n-type electron transport and electron selective contact materials in next-generation photovoltaics because of their high carrier mobility and proper alignment of energy band levels. The appropriate band edge alignment of these n-type metal oxides with those of p-PbS (type–II heterojunction) is favorable for better carrier transport to improve the performance of

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solar cells. This is due to the efficient spatial separation of the photo-generated electrons and suppressed electron–hole recombination at the p-n interface [4,9–11,17–19]. Lead sulfide has rock salt crystal structure and its band structure [20] is complicated by large relativistic splitting, making it suitable for band gap engineering. Bulk PbS has an infrared band gap (0.41 eV) that shifts to the visible region for the nanoclusters [21]. Because of its small effective mass, it shows a large blue shift in its absorption edge with a small change in cluster size [22]. In addition, PbS nanoclusters are expected to have exceptional third-order non-linear optical properties [23] and thus are potentially useful in optical and optoelectronic devices. There have been few reports focusing on ZnO/PbS [4,8–9,16–17,19], TiO2/PbS [10– 11,18], CdS/PbS [13,24–25] bi-layer heterojunction solar cells for the enhancement of photovoltaic characteristics in quantum dot and planar thin film architectures. According to theoretical analysis on nanostructured thin film photovoltaics, single-junction quantum dot photovoltaic efficiency of up to 15% may be practically achievable [1]. However, solar cells possessing various architectures, NW and QD based PbS/ZnO and PbS/CdS heterojunctions, with efficiencies in the range of 3–8% have been reported [4,9,19,24]. In all these cells PbS layer absorbs the solar flux and generates photo-charge carriers. Therefore, any enhancement in the absorption of solar flux, in particular in near IR, will generate larger photo-current. In addition, an efficient charge transport requires high carrier mobility concurrently with high carrier concentration, achievable in materials that are single crystal like. But by and large the PbS that has been used in these cells lacks such optimal structural and opto-electronic properties.

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Though there are reports on the synthesis of PbS by various techniques, but the study of their structural, electrical and optical properties has not been comprehensive. The reports deal only with one or two of the essential aspects [2,14,26,28-32]. Further, there are wide variations among the previous reports that could be due to the use of different techniques and different growth regimes. Thus, we find that a complete picture on the growth and associated optoelectronic properties is lacking. There has been hardly any study on the electrical transport behavior of these metal sulfide thin films [2,26–28]. Also, the synthesized PbS absorber material lacks the good oriented crystallinity and morphology needed for an efficient photon absorber and good carrier transporter. For improved cell performance, there is need for inorganic nanostructures with well-defined crystal shapes, larger surface area, high carrier mobility, good conductivity and excellent charge transfer property. Such considerations have started gaining attention of the researchers [1,3]. In particular, poor charge transfer is expected to be a major hindrance in attaining high efficiency in solar cells based on PbS films due to higher charge recombination at grain boundaries which necessitates the growth of PbS with suitable morphology and having high conductivity and mobility. Whereas the presence of flat surfaces is shown to lead to lower charge transfer resistance due to high surface area of cubes which significantly affect the photovoltaic performance [29], yet there is no study of the electron transport and electrical parameters in such film architectures. Other than the carrier concentration, conductivity and mobility are also determined by the structural and morphological quality of the films, and therefore it is quite essential to understand the correlation among all these aspects. Further, in many of the reports an additional inhibitor has been used as the complexing agent to grow the PbS thin films, which slows down the growth process to yield larger grains. 4

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But it is also likely to get incorporated in the film material with possible consequences on its crystallinity and carrier transport. Also with the addition of inhibitor, the resulting films are not (111) oriented, as required for better absorption of light. Because of the random orientation, such inhibitor grown films are likely to suffer from low absorption, low mobility, low carrier concentration, and higher resistivity, contrary to the requirements for better carrier transport in solar cells [2,29–33,36]. In the present work, chemical bath deposition (CBD) technique [34,35], a simple, fast, inexpensive, large area scalable technique that is capable of yielding good quality material with widely varying nano-architecture, is employed to grow cubic-nanocrystal PbS thin films. The intent of this work is to scrutinize a proper correlation among structural, morphological, optical, and electrical properties of the nanocubic p-PbS thin films for the improved performance of the thin film solar cells. The growth of well adherent and uniformly packed large sized (up to 400 nm) cubic-nanocrystals is demonstrated without employing any inhibitor in the reaction bath. Highly (111) oriented nanocubes of p-PbS with band gap controllable from visible to near-IR with high absorption in visible region, lowest room temperature resistivity of ~0.26 Ω-cm with exceedingly high mobility ~ 29.3 cm2V–1s–1 and high carrier concentration of 1018 cm–3 have been demonstrated at process temperatures close to room temperature. An excellent correlation among structural, morphological, optical, and electrical properties of these nanocubic p-PbS thin films has been established. 2.

Experimental

2.1.

Growth and characterization PbS thin films were fabricated by the CBD technique on glass substrates in alkaline

medium with different bath concentrations. All analytical grade reagents (Merck) were used

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without further purification for deposition. The films were grown by using lead acetate [Pb(CH3COO)2.3H2O] as a source of lead ions (Pb2+) by varying the concentration from 10 mM to 200 mM, and thiourea [SC(NH2)2] as a source of sulfide ions (S2–) by keeping Pb:S molar ratio of solution constant as 1:2. Ammonia solution that also acts as a complexing agent was added slowly to adjust the bath pH to 11. The solution was stirred for 30 minutes to ensure homogeneous mixing. The cleaned glass substrates were immersed vertically into the reaction bath maintained at 75°C with continuous stirring during deposition. After deposition, the substrate coated with PbS was taken out and thoroughly washed with distilled water and dried in ambient air. As-deposited PbS films were found to be dark brownish black in color, homogeneous, specularly reflecting and well adherent to substrate. As-prepared PbS thin films were characterized structurally using X-ray diffractometer (PANalytical X’pert Pro model) having CuKα radiation (λ = 1.54 Å) in 2θ range from 20°–80° using Bragg-Brentano geometry. The surface morphology of the prepared samples was investigated by using Field Emission Scanning Electron Microscopy (FESEM, FEI Quanta 200F SEM Model). The morphology of material was observed by using high resolution transmission electron microscopy (HRTEM, Tecnai G2 S-Twin, FEI, operated at 200 keV). The atomic percent composition of Pb and S elements in the synthesized PbS thin films was determined by the energy dispersive X-ray spectroscopy (EDX). The topography and roughness were imaged under ambient conditions by a Nanoscope-V scanning probe microscope from Bruker using tapping mode. The transmittance (T), reflectance (R) and absorbance (A) spectra of the samples in the wavelength range 350–2500 nm was recorded using Perkin Elmer Lambda 1050 UV–Vis–NIR spectrophotometer. Raman spectra were recorded by using Renishaw InVia Micro Raman Spectrometer in 120–800 cm–1 range,

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utilizing the 514-nm line of Argon ion (Ar+) laser and a 2400 lines/inch grating. The electrical properties were investigated by Hall Effect measurements at room temperature using Van der Pauw configuration. 2.2.

Reaction Mechanism Growth of PbS thin films take place when the ionic product of Pb2+ and S2– ions

exceeds the solubility product of the PbS (Ksp = 10–29) [34,35] and therefore Pb2+ and S2– ions concentration has to be controlled very carefully during the growth. The deposition process is based on the slow release of the Pb2+ and S2– ions in the reaction solution, which arrive on the substrate surface by an ion-by-ion mechanism. The ions mobility is larger in ion-by-ion deposition process than the colloid-in-complex decomposition mechanism [33] and is controllable over a wide range through ion-flux rate. The formation of Pb(NH 3 ) 24 complex with the Pb2+ ions by addition of liquid ammonia reduces the rate of PbS precipitation in the bulk solution and thereby lowering the concentration of free Pb2+ ions into the solution. The nucleation centers on glass substrate for the growth of PbS thin films are provided through judicious formation of the traces of Pb(OH)2 in the reaction bath. The white turbid appearance in the solution is an indication of hydroxide formation and is the pre-requisite for the film deposition by homogeneous nucleation process. The color of the solution slowly changes brownish to dark brown after heating, which indicates initiation of chemical reaction and PbS film growth on the substrate. The chemical reaction involved for the formation of PbS films from Pb(CH3COO)2.3H2O is as follows (Eq. 1) [34]: Pb(CH3COO)2 (aq) + SC(NH2)2 (aq) + 2OH– → PbS + H2CN2 + 2H2O + 2CH3COO–…… (1)

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3.

Results and discussion

3.1.

Morphology and Composition Fig. 1 shows the top-surface FESEM images of chemically deposited PbS thin films.

The micrographs reveal that the PbS films exhibit well oriented nanocubic morphology and the whole surface of the substrate is uniformly covered with these nanocubes that are well faceted and are oriented such that the (111) axis of the cubes is normal to substrate. The surface morphology of the films is greatly influenced by the bath concentration and all the films are quite uniform. Lowering of bath concentration results in better (111) aligned and well-formed nanocubes with larger facets. The compact films formed at higher bath concentrations become just separated at 25 mM concentration and at 10 mM concentration the cubes are separated. It can be further seen that with the decrease in bath concentration the facets of PbS nanocubes get well defined and there is a corresponding reduction in density of nanocubes and an increase in cube-size from ~100 nm to ~400 nm. The composition of the PbS thin films deposited on glass substrates is determined by EDX analysis. The data obtained from the quantitative analysis indicate that the atomic ratio of Pb to S is approximately 1:1 for all the films and confirms the stoichiometric growth of PbS. 3.2.

Structural properties The XRD spectra of PbS films deposited at various bath concentrations 200 mM, 100

mM, 50 mM, 25 mM, and 10 mM are shown in Fig. 2. All the observed peaks correspond to the face centered cubic rock-salt structure of PbS. The absence of any additional peak suggests that no other impurity phase is formed. However, the peak intensities do not correspond to the powder standard, and the (111) is the most intense reflection rather than the (200). The lattice parameter ‘a’ of PbS is calculated by using the Eq. 2 [14, 31]:

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1 h2  k 2  l 2 ………… (2)  d2 a2

The mean value of ‘a’ is estimated by averaging lattice parameters calculated for (111), (200), (220), and (311) reflections and is shown in Table I. It found to be a = 5.940 Å and is in good agreement with the reported value of 5.936 Å [JCPDS 05-0592]. It can be further seen that the intensity of four main peaks (111), (200), (220), and (311) is sharply increased with the decrease in bath concentration to 25 mM, which indicates a significant increase in crystalline quality of the films. The average crystallite size (D) is calculated by using Scherrer formula (Eq. 3) [26]:

D

0.9  ………… (3)  cos 

where, λ is wavelength of CuK X-ray radiation (1.54 Å), β is the full width at half maximum (FWHM) in radian and θ is Bragg angle. The calculated crystallite sizes are found to increase from 20 nm to 26 nm with the decrease in bath concentration up to 25 mM, but it decreases for 10 mM. The observed trend in crystallite size can be correlated with reaction rate which is controlled by varying the bath concentration. For 200 mM film, large numbers of Pb2+ and S2– ions are available for the reaction and the decreasing bath concentration to 25 mM leads to lesser ion flux arriving at the substrate resulting in larger grain size as well as nanocube size and higher texture. This confirms that the growth process follows nucleation controlled mechanism. But the slight decrease in crystallite size for 10 mM indicates that 25 mM is a threshold concentration below which, though the reaction rate is reduced, the number of ions are also decreased to the extent that crystallite size could not grow beyond certain size as well as substrate coverage is not complete.

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Texture coefficient (TC) corresponding to (111), (200), (220), and (311) planes is calculated in all the films by Hall method to find the quantitative information of the preferred growth direction using the Eq. 4 [26,32,33]:

TC (hkl ) 

I m hkl / IC hkl n

(1/ n) I i 1

m

………… (4) C

hkl / I hkl

where I m hkl and IC hkl are, respectively, the measured intensity and the standard intensity of (hkl) plane based on standard data file (JCPDS 05-0592), and n is number of diffraction peaks considered for analysis. TC(hkl) values represent the abundance of crystallites oriented in a given (hkl) plane. TC is expected to be maximum and greater than unity for the preferentially grown facet, and if it is ≤ 1 the crystallites are considered to be randomly oriented [26]. From Table I it can be seen that for all the samples TC along (111) plane is maximum and more than 1, whereas for other planes it is either less than or equal to 1. So, we can conclude that (111) is the preferred orientation of all the samples. This deduction is further confirmed by our electron microscopy studies presented next. 3.3.

High resolution-transmission electron microscopy (HRTEM) HRTEM is employed to further explore the morphology and crystal structure of as-

deposited PbS thin films. The bright field TEM images of three selected samples 100 mM, 50 mM, and 10 mM are shown in Fig. 3(a–c). It reveals very well formed nanocubes in consonance with the observed morphology in FESEM. The size of nanocubes determined from TEM images show good agreement with that determined from FESEM (~100–400 nm). No substructure is observed, indicating that nanocubes are quite compact and not allowing the

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electron beam to pass through them. The HRTEM images taken for different PbS films are shown in Fig. 3(d–f). The observed lattice fringes in HRTEM images reveal the good crystallinity of PbS films. All the nanocubes exhibit an inter-planar spacing of ~0.357 nm belonging to the (111) lattice plane of the face centered cubic phase of PbS, suggesting that (111) plane in the all the films grow normal to substrate, which is in good agreement with the XRD analysis. 3.4

Atomic force microscopy (AFM) The surface topography and roughness of as-deposited PbS thin films were investigated

by employing AFM in tapping mode. Figs. 4(a) & 4(b) show AFM images of two representative films, grown at 100 mM and 25 mM bath concentrations, respectively. It can be observed that the topography is consistent with the FESEM study according to which size of cubic nanocrystal increases with decrease in bath concentration. The rms-surface roughness of the PbS thin films is found to lie in the range of 9−12 nm increasing with decrease in bath concentration. Also, somewhat greater surface roughness at lower bath concentration of PbS thin films is beneficial for better absorption of light, but its magnitude is quite small to be of concern for device fabrication. 3.5.

Raman spectroscopy The Raman spectra recorded at room temperature on all the PbS thin films in the range

of 120–800 cm–1 using the 514-nm excitation line are shown in Fig. 5. The Raman bands are observed at following positions: ~178 cm–1, ~200 cm–1 and ~441 cm–1. The band at ~200 cm–1 can be ascribed to the LO phonon at Brillouin zone center and the weak band at ~441 cm–1 is assigned to the first overtones of the LO phonon (2LO) mode [28,37]. So, we can say that Raman bands correspond to the PbS, and absence of any additional band confirms the pure 11

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phase of PbS. It should also be noted that a strong peak is observed at ~178 cm–1 and it can be correlated with the surface phonon (SP) mode. This mode is generally observed if the surface to volume ratio is large [38]. In our case where the nanostructure in form of nanocubes is observed, it accounts for the large surface to volume ratio and hence the SP band is observed. It is noticeable that as we go from 200 mM sample to 25 mM sample, not only the intensity of the overall Raman signal significantly increases but the intensity of the SP mode peak becomes larger than the LO phonon peak. This gradual increase in intensity with lowering of bath concentration provides support to the increasing area associated with the (100) facets, consistent with the FESEM results. The slight red shift is found in the peak positions and is reported to be observed in films showing quantum confinement [14]. It is to be noted that the crystallite size observed in our case is < 30 nm, as required for the quantum confinement in PbS thin films. Thus, the red shift in the peaks can be ascribed to quantum confinement. 3.6.

Optical properties Fig. 6(a) shows the transmittance spectra of as-deposited PbS thin films recorded in

wavelength range of 350–2500 nm. All the spectra show on-set of transmission in the wavelength range 600–800 nm, suggesting the occurrence of direct energy bandgap in this spectral range, higher than that of bulk PbS at 3000 nm. Whereas the on-set of transmission is sharp for samples with bath concentration ≥ 50 mM, the lower bath concentration samples show a gradual on-set suggesting enhanced absorption in the NIR region, a requirement for absorbers capable of capturing lower wavelength photons as well that are abundant in solar spectrum. The average transmittance in NIR region (1000–2000 nm) is 20% for < 50 mM samples in comparison to 40% for ≥ 50 mM samples. The reflectance spectra recorded at normal incidence are shown in Fig. 6(b) and these are complimentary to transmission spectra, 12

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signifying specular nature of the samples. More important to note is the lower average reflectance of 25% for < 50 mM samples in comparison to 40% for ≥ 50 mM samples. This clearly establishes significant enhancement in absorption of the 25 mM and 10 mM films. The absorbance spectra are shown in Fig. 6(c) and it shows significantly higher absorption in the sub-bandgap regime of 1200–2400 nm (NIR range) for the 25 mM and 10 mM samples. There is also a red shift of this NIR absorption band. This additional NIR absorption band is missing in  50 mM samples that are very compact and show absorption in the above bandgap regime only. A possible interpretation for this NIR band is that the PbS nanocubes are in just separated state in the 25 mM and separated in 10 mM samples, leading to generation of surface plasmons and their coupling with neighboring particles [39–40]. Also, the observed red-shift and slightly higher absorption for the 10 mM film can be linked to the change in the proximity and increase in facet-area of nanocubes due to increase in their size. In order to quantify the absorption of solar photon flux, we have calculated the photon flux within standard AM1.5G solar spectra [ASTM G-173-03 (International standard ISO 9845-1, 1992)] in full spectrum range (400–2400 nm) and NIR range (1200–2400 nm). The calculated value of flux in standard full spectrum range is 4.06  1021 photons.s–1m–2, and in NIR range is 1.24  1021 photons.s–1m–2 that is ~30% of full solar spectrum. The conventional solar absorbers like silicon are unable to absorb this sizable part of solar spectrum. The calculated absorbed solar photon flux in NIR range increases from 6.66  1019 to 5.71  1020 photons.s–1m–2 as we move from 200 mM to 10 mM samples and is shown in Table I. The absorbed fraction of NIR photon flux is ~5.37% for 200 mM film and increases to ~41% and ~46% for 25 mM and 10 mM bath films respectively. This data is shown in Fig. 6(d) for various samples. In comparison to silicon, the conventional solar absorber, the NIR photon 13

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absorption gain is very substantial [41,42]. Thus, our 25 mM PbS nanocubic films absorb ~61% fraction of total solar flux in the range 400–2400 nm. The optical absorption coefficient (α) at different wavelengths is calculated by combining the transmittance (T) and reflectance (R) of the film and described in Eq. 5 [26,43]:

1  (1  R) 2    ln   ………… (5) t  T  where, t is the film thickness, T and R are the transmittance and reflectance of the films respectively. The absorption coefficient (α) is related to the band gap (Eg) of the material by the Tauc’s relation (Eq. 6) [14,31,33]:

 h  Ao (h  Eg ) n ………… (6) where n = ½ for allowed direct transition and n = 2 for allowed indirect transition. Ao is the parameter which depends on the transition probability. The linear dependence of (αh)2 vs h indicates the presence of direct band gap in all the deposited films (Fig. 6(e)). The direct band gap energy (Eg) is obtained from the intercept on energy axis by extrapolating the linear portion of the graph. Eg values are found to be 1.50–0.94 eV and showing decreasing trend with the decrease in bath concentration from 200 mM to 25 mM. However, Eg is slightly increased to 1.06 eV for 10 mM (Fig. 6(f)). This can be correlated with the slight decrease in crystallite size calculated from XRD data. The Eg is found to be varying inversely with the crystallite size as is reported earlier [21,22]. The relatively large size of nanocubes may be due to the presence of nearly aligned crystallites in very close proximity inside the nanocubes. The quantum confinement is reported in PbS films having crystallite size smaller than 30 nm [14]. Thus, the observed red shift of the band edge with the decrease in bath concentration arises from 14

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quantum confinement effect in PbS thin films. Thus, we have been able to tailor the optical band gap from near infrared to visible region via a simple change in bath concentration. 3.7.

Electrical Properties

Electrical properties of CBD deposited PbS films at room temperature are obtained by using Van der Pauw Hall effect measurement at room temperature, from which we obtain different parameters viz. carrier concentration (p), carrier mobility (μ), and resistivity (ρ) of the deposited PbS nanocubic films and shown in Fig. 7. Hall measurement also confirmed the ptype semiconducting behavior of all the films. The values of the various electrical parameters are given in Table II. It is found that ρ decreases with the decrease in bath concentration from 200 mM to 25 mM. The lowest resistivity obtained is 0.26 Ω–cm at bath concentration of 25 mM. It may be pointed out that it is the lowest value reported so far [2,26,28,31] for PbS films. This observed trend in ρ can be understood in terms of mobility changes. It can be noted that the mobility increases significantly (by 5 times) with the decrease of bath concentration from 200 mM to 25 mM. The carrier concentration also increases with decrease in bath concentration, but not as sharply as carrier mobility. So, in our films the increase in mobility is mainly deciding the resistivity decrease and it can be interpreted in terms of the possible improvement in alignment of crystallites which allows the carriers to easily cross the intergrain boundaries due to reduced grain-boundary potential. This interpretation can be seen in light of observation of ~250 nm size nanocubes consisting of ~25 nm size grains, which seem to be almost aligned crystallographically, thus allowing inter-grain carrier transport with minimal scattering. Highest values of p and μ obtained are 8.1 x 1017 cm–3 and 29.3 cm2V–1s–1, respectively, for 25 mM nanocubic films. For the lowest bath concentration of 10 mM compactness decreases compared to higher concentration films (Fig. 1) and nanocubes are seen 15

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as separated. This causes an increase in inter-nanocube distance and hence the increase in resistivity to 0.72 Ω–cm is observed. So, it can be concluded that 25 mM is a threshold concentration below which contribution from inter-nanocube distance pre-dominates over the increase in alignment of grains. It may also be noted that at 10 mM concentration the nanocubes have (111) as well as (200) alignment. There are some previous reports [28,30,32,33,36] on the growth of PbS thin films using CBD technique, either employing very high bath concentrations [28,32,33] or high bath pH [30,36]. Moreover, many of these are additionally using inhibitors [28,30,32,33,36] such as sodium sulfite, trisodium citrate, and triethanolamine to grow the PbS films at lower rates. The inhibitor slows down the reaction rate in solution by complexing with Pb2+ ions. As a result, rapid reaction and precipitation of metal hydroxide is prevented, but at the same time the presence of the inhibitor seems to restrict the grain growth and grain sizes 10-50 nm are generally reported. Further, such grown films exhibit poor electrical transport parameters due to the inhibitor induced enhancement in the grain boundary scattering and also as a consequence of the creation of additional defects and trap states on the surface. Moreover, bath pH > 12 containing inhibitors seem to promote either (100) or (110) oriented planes, contrary to the required (111) plane for better light absorption. In accordance with the nucleation controlled growth mechanism we expect that lowering of growth rate will yield better crystal quality and larger grain size. However, it was not achieved by the use of inhibitors. Therefore, to replace the role of inhibitor for lowering the growth rate, we have used lower bath concentration (10–200 mM) with bath pH < 12 to keep the ion flux low. This significantly inhibits the precipitation by permitting a slow and controlled release of Pb+2 ions and thus allowing the growth to occur via ion-by-ion deposition

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mechanism. Thus, the resulting films in our case are well faceted cubic nanocrystals having good crystallinity and larger grain size (up to 400 nm). In addition, the inhibitor free growth promotes the formation of (111) oriented PbS films, having high optical absorption and very good electrical properties which are essential requirements for higher photocarrier generation and better carrier transport in solar cell materials. Now we compare our results with previously reported data on PbS films. Yeon et al. [2] have reported the efficiency of 3.1% for chemical bath deposited CdS/PbS heterojunction thin film solar cells with highest carrier concentration 5.6 × 1016 cm–3 of PbS which is one order lower than that of our PbS nanocubic films. Khot et al. [14] have prepared PbS films with random orientation of cubical grains (surprisingly the orientation keeps changing with thickness increase) and there is absence of the electrical transport behavior as well as any proper correlation among reported measurements. Hernandez-Borja [25] has reported the efficiency 1.63% for planar PbS/CdS solar cell, which was low as their films had poor photon absorption. Mathews et al. [26] and Preetha et al. [31,32] have reported PbS thin films with very high room temperature resistivity of the order of 104–107 Ω–cm, which is huge compared to our nanocubic films and not suitable at all for solar cell application. On CBD grown PbS films Slonopas et al. [28] reported electrical properties comparable with ours but the grains are with random orientation and there is no study of optoelectronic behavior of these PbS films. Rao et al. [29] have used urea as inhibitor to grow oriented nanocubes. Without inhibitor, the growth is of different sized random cubes. Sengupta et al. [30] have also used citrate as additional inhibitor to grow (200) oriented films with < 100 nm nanocubes at low temperature. There is absence of electrical and optical study of these films. But we have done a controlled growth of well oriented cubic nanocrystals of sizes as large as 400 nm without using any

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inhibitor and carried out a complete study of tunable optical and electrical behavior of PbS films with proper correlation among them. 4.

Conclusions Highly crystalline PbS films consisting of highly oriented cubic nanocrystals of up to

400 nm sized facets possessing exceedingly good structural, crystallographic, electrical, and optical quality and having good adhesion and specular nature are successfully grown by CBD technique. Their optoelectronic properties and cubic-nanocrystal morphology has been tailored with bath concentration variation from 10 mM to 200 mM. It is established that all the films consist of very well faceted and near parallel aligned nanocubes that are preferably oriented along (111) axis of the cube. The nanocube size as well as the crystallite size are controlled by the ion-supply rate and are found to increase from 100–400 nm and 20–26 nm, respectively, with the decrease in bath concentration. XRD and Raman studies confirmed the phase pure nature of the films. The optical direct band gap is continuously varied from NIR to visible region (0.94–1.50 eV) via bath concentration and is interpreted by the quantum confinement effects. All the films are p-type semiconducting and their electron transport properties are tunable by the bath concentration. The hitherto unreported lowest room temperature resistivity of 0.26 Ω–cm, and concurrently highest mobility of 29.3 cm2V–1s–1 and high carrier concentration of 1018 cm–3 are observed for PbS film deposited at 25 mM bath concentration and understood in terms of decrease in grain boundaries on decreasing the bath concentration. Occurrence of additional NIR photon absorption due to surface plasmons is a very special feature (unobserved so far) of the 25 mM and 10 mM films possessing well-resolved and presumably optimally connected nanocubes having large facet area.

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Figure Captions Fig. 1. FESEM images of PbS thin films deposited at different bath concentrations, (a) 200 mM (b) 100 mM (c) 50 mM (d) 25 mM (e) and 10 mM. Fig. 2. XRD pattern of PbS thin films deposited at 200 mM, 100 mM, 50 mM, 25 mM, and 10 mM bath concentrations. Bottom pattern shows standard JCPDS data (No.05-0592). Fig. 3. (a–c) TEM images and (d–f) HRTEM images of PbS thin films deposited at 100 mM, 50 mM, and 10 mM bath concentrations, respectively. Fig. 4. AFM images for PbS thin films deposited at (a) 100 mM and (b) 25 mM bath concentrations. Fig. 5. Room temperature Raman spectra of PbS thin films deposited at different bath concentrations. Fig. 6. (a) Transmittance spectra, (b) Reflectance spectra, (c) Absorbance Spectra, (d) Fraction of absorbed NIR Photon Flux, (e) Tauc’s plot for direct band gap (Eg) and (f) Variation in Eg values, of PbS thin films deposited at different bath concentrations. Fig. 7. Variation of room temperature resistivity, mobility and carrier concentration for PbS thin films deposited at 200 mM, 100 mM, 50 mM, 25 mM, and 10 mM bath concentration.

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ACCEPTED MANUSCRIPT TABLES Table I. Structural and optical parameters of the PbS thin films at various bath concentrations. Sample

(hkl)

200 mM

111 200 220 311 111 200 220 311 111 200 220 311 111 200 220 311 111 200 220 311

100 mM

50 mM

25 mM

10 mM

a=b=c (Å)

D(nm)

5.95

20.3

5.95

20.9

5.94

25.0

5.94

25.7

5.94

23.8

(TC)

1.58 1.02 0.81 0.58 1.50 1.16 0.76 0.57 1.40 1.09 0.84 0.65 1.38 1.03 0.89 0.68 1.41 1.08 0.84 0.67

Eg (eV)

Absorbed NIR Photon Flux (photons s–1m–2)

NIR Absorption Fraction (%)

1.50

6.66  1019

5.4

1.30

1.36  1020

11.0

1.12

2.13  1020

17.2

0.94

5.02  1020

40.5

1.06

5.71  1020

46.1

ACCEPTED MANUSCRIPT Table II. Electrical parameters of the PbS thin films deposited at various bath concentration

Sample

ρ (Ω–cm)

p (cm–3)

μ (cm2V–1s–1)

200 mM

2.05

5.1  1017

6.0

100 mM

1.46

5.6  1017

7.7

50 mM

0.93

6.5  1017

10.3

25 mM

0.26

8.1  1017

29.3

10 mM

0.72

5.4  1017

16.0