Binder-Free CuâIn Alloy Nanoparticles Precursor and Their Phase

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Binder-Free Cu-In Alloy Nanoparticles Precursor and Their Phase Transformation to Chalcogenides for Solar Cell Applications Ye Seul Lim, Jeunghyun Jeong, Jin Young Kim, Min Jae Ko, Honggon Kim, BongSoo Kim, Unyong Jeong, and Doh-Kwon Lee J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/jp401637b • Publication Date (Web): 16 May 2013 Downloaded from http://pubs.acs.org on May 24, 2013

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The Journal of Physical Chemistry

Binder-Free Cu-In Alloy Nanoparticles Precursor and Their Phase Transformation to Chalcogenides for Solar Cell Applications

Ye Seul Lim†,‡, Jeunghyun Jeong†, Jin Young Kim†, Min Jae Ko†, Honggon Kim†, BongSoo Kim†, Unyong Jeong‡, and Doh-Kwon Lee*,† †Photo-electronic

Hybrids Research Center, Korea Institute of Science and Technology (KIST), Seoul, Korea

‡Department

of Materials Science and Engineering, Yonsei University, Seoul, Korea

*

Corresponding author: Tel. +82-2-958-6710, Fax.: +82-2-958-6649, E-mail: [email protected]

ABSTRACT A low-cost, non-vacuum fabrication route for CuInSe2 and CuInS2 thin films is presented. To produce these films, binder-free colloidal precursors were prepared using Cu-In intermetallic nanoparticles that were synthesized via chemical reduction method. The Cu-In alloy precursor films were transformed to CuInSe2 and CuInS2 by reactive annealing in chalcogencontaining atmospheres at atmospheric pressure. The as-synthesized nanoparticles and the annealed films were characterized by XRD, TEM, SEM, EDS, EPMA, Raman spectroscopy, and AES depth profile measurements to elucidate the phase evolution pathway and the densification mechanism of the Cu-In-Se-S system. Solar cell devices made with CuInSe2 and CuInS2 absorbing layers exhibited power conversion efficiencies of 3.92% and 2.28%, respectively. A comparison of the devices suggested that the microstructure of the absorbing 1

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layer had a greater influence on the overall photovoltaic performance than the band gap energy. A diode analysis on the solar cell devices revealed that the high saturation current density and diode ideality factor caused lower open-circuit voltages than would be expected from the band gap energies. However, the diode analysis combined with the microstructural and compositional analysis offered guidance about how to improve the photovoltaic performance of these devices.

Keywords: CIS, CISe, thin film solar cell, non-vacuum process, printing, colloidal precursor

1. INTRODUCTION I-III-IV2 compounds with a chalcopyrite structure, such as CuInSe2 and CuInS2, have been regarded as the most promising light-absorbing materials for thin film photovoltaic (PV) cells. On the laboratory scale, power conversion efficiencies over 20% have been achieved by partial substitution of gallium for indium.1 This outstanding performance, which is partly due to advantageous properties for solar cell applications such as direct band gap transitions and high absorption coefficients, and their high material stability including exceptional radiation hardness,2,3 have evoked the widespread commercialization prospects of Cux(In1-yGay)(Se1zSz)2

(CIGSeS) based solar cells. Besides, the ability to tune the band gap energies Eg in the

range from 1.0 eV to 2.4 eV by adjusting their composition (from y = z = 0 to y = z = 1)4,5 makes a multi-junction (tandem) solar cell feasible as an approach to surpass the current limits of CIGSeS cell efficiency.6 In this regard, CuInSe2 (CISe) could be an ideal candidate for the bottom cell in double-junction tandem solar cells. Alternatively, CuInS2 (CIS) has attracted considerable attention due to its wide band gap (Eg = 1.53 eV) and lower toxicity in comparison to its widely investigated selenide counterpart .7

2


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The absorber layers in state-of-the-art CIGSeS solar cells are generally fabricated using high vacuum techniques, e.g., a multi-stage co-evaporation or a two-step process of sputtering and subsequent chalcogenization.8 Recently, non-vacuum (wet) processes have been spotlighted as alternatives to these vacuum-based methods due to their potential to realize low-cost PV devices. The advantages of non-vacuum techniques include higher efficiency in materials utilization, lower initial capital cost for deposition equipments, lower energy demand for deposition, and ease in scale-up due to the compatibility with highthroughput roll-to-roll processes.9,10 Non-vacuum coatings of CIGSeS are carried out by using a paste or ink, which can be generated by dissolving various types of metal salts and/or selenides in a solvent to prepare a solution precursor11-21 or by dispersing nanoparticles in a solvent to prepare a colloidal precursor.22-31 The process to synthesize a solution precursor is relatively simple to carry out. Solution precursor routes often employ organic binders and/or thick solvents to get suitable rheological properties for coating. However, incomplete decomposition of these organic substances typically leads to carbon residue in the absorber films, which may induce high series resistance in the resulting solar cells. In contrast, well-dispersed nanoparticles in colloidal precursors may act as a thickening agent to some degree, which enables organic binder-free coating. The colloidal precursor routes have additional advantages in that the nanoparticles can be synthesized with high purity and a well-defined composition, and their high reactivity is beneficial during phase transformation and/or densification.28 Accordingly, various types of nanoparticles have been used to prepare the colloidal precursors including elemental metals or their alloys, binary metal selenides/sulfides, metal oxides, and ternary or quaternary selenides/sulfides.8-10,21-30 By employing metallic-type nanoparticles, one can utilize the volume expansion associated 3



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with chalcogenization of metals in densification of absorber layers. The use of Cu-In alloy powders in non-vacuum fabrication of CISe solar cell was first reported by Norsworthy and coworkers who demonstrated 10.5% efficiency employing H2Se(g) for selenization.22 The metallic particles in their work were 0.53 µm in size, however, which seem too large to bring about a well-packed precursor film of 3-4 µm thickness, possibly requiring a highly reactive but highly toxic H2Se gas for densification of the films. Moreover, the organic additives in their colloidal solution may improve the morphology of the precursor films but can leave carbon contamination in the sintered films. On the other hand, despite growing interest in the wet-processed CIGSeS solar cells, systematic interpretations on their characteristic properties of materials and devices are quite limited. A better understanding on the phase transformation and densification behaviors and on the origin of the limited device efficiency of the non-vacuum processed CIGSeS absorbers (apparently due to low open-circuit voltage and fill factor in general) is believed to be essential to improve their PV performance. In this study, Cu-In intermetallic nanoparticles with less than 100 nm diameters were synthesized via chemical reduction method32-35 under ambient conditions. Binder-free colloidal precursors were prepared therewith by ball milling in ethanol without adding any other organic substance. After preparing the precursor films by doctor-blading, a low-cost, non-vacuum chalcogenization of the films were carried out using Se(g) and H2S(g) as chalcogen sources to produce CISe and CIS thin films. The phase evolution pathway from the Cu-In alloy to CISe or CIS was systematically investigated in conjunction with the compositional and microstructural changes upon chalcogenization as a function of annealing temperature. The densification mechanism of the Cu-In-Se system is discussed in association with the observed compositional unmixing of cations. The CISe and CIS layers prepared 4


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from Cu-In alloy precursor were integrated into solar cell devices, which exhibited power conversion efficiencies of 3.92% and 2.28%, respectively. A diode analysis is presented to elucidate the origin of the limited performance of the present devices and to suggest ways to improve it.

2. EXPERIMENTAL 2.1. Synthesis of metallic nanoparticles, slurries, and precursor films Cu-In metallic nanoparticles were synthesized via chemical reduction method under ambient conditions using metal chlorides as metal sources and sodium borohydride as a reducing agent. A precursor solution was prepared by dissolving 0.747 g of copper (II) chloride (CuCl2, Aldrich, 99.999%) and 1.536 g of indium (III) chloride (InCl3, Aldrich, 99.999%) to yield a molar ratio of 0.8:1 in 50 mL of tetraethylene glycol (TEG, Aldrich, 99%). To prepare a reducing solution, 2.838 g of sodium borohydride (NaBH4, Junsei, 98%) was dissolved in 50 mL of TEG. To control the reduction kinetics,34,35 the precursor solution was added at a rate of 1 mL/min into the reducing solution that was kept at 0 °C in an ice bath with magnetic stirring. The resulting precipitates were separated from the solution through vacuum filtration and were washed with deionized water, ethanol, and methanol before being dried at 60 °C for 6 h. The yield of nanoparticle product was over 95%. The first step to produce a colloidal precursor (slurry) for printing was to disperse 4 g of nanoparticles in 25 mL of anhydrous ethanol without adding any organic dispersant or binder. Then, the solution was ball-milled at 200 rpm using ZrO2 balls with 1 mm and 5 mm diameters in a ratio of 7:3. Cu-In metallic precursor films were coated on Mo-sputtered soda-lime glass (SLG) substrates by doctorblading the colloidal precursor (Sheen Instruments, model 1117); the films were then dried under ambient conditions. Because the precursor films are prepared by a binder-free process 5



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using a volatile solvent, one can skip the oxidation step at elevated temperature that is often otherwise needed to remove organic residues either from the dispersant/binder or non-volatile solvents before chalcogenization of the films.

2.2. Formation of chalcogenide films by reactive annealing The phase transformation and densification that occur upon reactive annealing of the asdeposited films were investigated as a function of temperature in the range of 100 < T / °C < 550 by performing heat treatments in Se(g)- or H2S(g)-containing gas atmospheres for 30 min. For selenization of the precursor films, Se vapor was provided by melting selenium pellets (99.99%, Aldrich) of ca. 0.2 g placed 5 cm away from the films at the center of a horizontal tube furnace. As a carrier gas, a 4% H2-96% Ar gas mixture was used at a flow rate of 150 sccm. In this way, the temperature of the selenium pellets were kept 40 ~ 140 °C lower than the temperature of the films during annealing as measured by thermocouples that monitored the Se source and the film. According to the Clausius-Clapeyron equation, the Se vapor pressures PSe at the temperatures of annealing, Ta, were estimated from the equilibrium values for each temperature of Se, TSe,36 which are listed in Table 1. For sulfurization, an 1% H2S-99% Ar gas mixture was employed at a flow rate of 150 sccm. The precursor films were heated up to annealing temperatures at a rate of 10 °C/min. All heat treatments were carried out under ambient atmospheric pressure.

6


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Table 1. Selenium vapor pressures, PSe, at temperatures of annealing, Ta, controlled by the corresponding temperatures of the selenium source, TSe. Ta / °C

TSe / °C

PSe/ atm

100 °C

66 °C

-

200 °C

110 °C

-

300 °C

180 °C

-

400 °C

260 °C

7.07×10-5

500 °C

355 °C

1.84×10-3

550 °C

410 °C

8.02×10-3

2.3. Characterization of nanoparticles and thin films The surface and cross-sectional morphologies of the as-synthesized nanoparticles, asdeposited precursor films, and annealed films under chalcogen-containing atmospheres at different temperatures were characterized using field emission scanning electron microscopy at an acceleration voltage of 15 kV, and their compositions were analyzed by energy dispersive X-ray spectrometry at an acceleration voltage of 20 kV and an acquisition time of 60 s (FE-SEM/EDS, Hitachi, S-4200). The EDS spectra were collected from at least 5 randomly selected areas over an entire sample surface measuring 25 mm × 25 mm. The compositions of thin films were confirmed with electron probe X-ray micro-analysis with a beam spot radius of 30 µm (EPMA, Jeol, JXA-8500F) by measuring 5 times. Both chemical analyses showed good agreement in film compositions within the error bounds. The multiphase nanoparticles were further examined by transmission electron microscopy (TEM, FEI, Tecnai F20G2) using an acceleration voltage of 200 kV. The energy-dispersive spectrometer equipped on the TEM was used to semi-quantitatively analyze the compositions of nanoparticles in different phases. The crystal structure of the nanoparticles and annealed films 7



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were analyzed using an X-ray diffractometer (XRD, Rigaku, D/max 2500) operating at 60 kV and 300 mA in a θ-2θ scan mode with Cu-Kα radiation (λ = 0.15418 nm) and with Ni as a Κβ filter. Raman spectra of the annealed films were recorded by a Jobin-Yvon T64000 system with a spectral resolution of 0.6 cm-1 that employs an 1800 grooves/mm grating. The excitation beam of a 5 mW Ar+ laser with 514.5 nm radiation (Lexel Laser, Inc., model 95) was focused on a spot of ca. 1.4 mm diameter through an objective microscope with a magnification of 50×. The scattered radiation was filtered by a notch filter and collected in a backscattering geometry with a 2D liquid-nitrogen-cooled charge coupled device (CCD) detector. The relative atomic concentration profiles across the thickness of the selenized/sulfurized films were examined by Auger electron spectroscopy (AES, Scanning Auger Nanoprobe PHI-700 & LC-TOFMS LECO) with ion beam sputtering at an etching rate of 50 nm/min (estimated with SiO2).

2.4. Solar cell fabrication and analysis Solar cells were fabricated in a conventional configuration, Mo/(CISe or CIS)/CdS/i-ZnO/nZnO/Al. A CdS buffer layer of ca. 70 nm thickness was deposited on the chalcogenide film via chemical bath deposition (CBD) with a solution consisting of 2 mM CdSO4, 1.02 M NH4OH, and 84 mM thiourea at 60 °C for 20 min.37 Intrinsic ZnO (i-ZnO, 50 nm) and Gadoped ZnO (GZO, 250 nm) were deposited onto the CdS layer using a radio-frequency magnetron sputtering method. A Ni/Al grid (50 nm/500 nm) was deposited by an electron beam evaporation technique onto the ZnO layer with the remainder of the surface being as an active area measuring 0.3 to 0.4 cm2. Photocurrent-voltage (j-V) characteristics of the solar cells were investigated using a Keithley 2400 source meter. A class-AAB solar simulator (Yamashita Denso, YSS-200A) equipped with a 1600 W Xenon lamp and an AM 1.5G filter 8


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(ASTM E927-05, IEC 60904-9) was used as a light source. The light intensity was adjusted with an NREL-calibrated Si solar cell with a BK7 filter to approximate one sun of light intensity (100 mW/cm2). External quantum efficiencies (EQE) were measured under short circuit conditions as a function of wavelength with a spectral resolution of 10 nm using an incident photon-to-current conversion efficiency measurement system equipped with a 75 W xenon lamp and a grating monochromator (PV Measurements, Inc.). A calibration was carried out using a silicon photodiode (G425) in the spectral range from 300 nm to 1000 nm and a germanium photodiode from 900 nm to 1400 nm; both photodiodes had standard NISTcalibration.

3. RESULTS AND DISCUSSION 3.1. Characterization of the prepared metallic nanoparticles Figure 1 shows the morphology of the as-synthesized nanoparticles produced with a chemical reduction method under ambient conditions using CuCl2 and InCl3 precursors in a ratio of [Cu]/[In] = 0.8 together with NaBH4 in TEG as a reducing solution. Under SEM, the nanoparticles appear to be mono-dispersed with particle sizes of tens of nanometers. The corresponding EDS analysis revealed that the as-prepared nanoparticles have an overall composition of [Cu]/[In] = 0.84 ± 0.05, which was also confirmed by EPMA measurements on the as-coated precursor film on Mo substrates. The agreement between the batch composition of the starting materials and the composition of the as-synthesized nanoparticle product implies that copper and indium ions in the TEG solution were completely reduced in the presence of NaBH4. A copper-poor composition of the Cux(In1-yGay)(Se1-zSz)2 compound semiconductor (in this study, x = 0.84) is generally favored in solar cell applications because otherwise the segregation of conductive Cu-rich binary phases such as Cu2Se and CuSe 9



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would cause the formation of shunt paths in the absorber layer and would play a detrimental role in photovoltaic devices.22,38

Figure 1. SEM image of Cu-In intermetallic nanoparticles synthesized via chemical reduction method employing NaBH4 and TEG as a reducing agent and solvent, respectively.

The X-ray diffraction pattern of the as-prepared nanoparticles is shown in Figure 2. An analysis of the XRD peaks revealed that the obtained nanoparticles were composed of three crystalline phases, i.e., monoclinic CuIn (P21/m, JCPDS No. 35-1150), hexagonal Cu2In (P63/mmc, JCPDS No. 65-704), and tetragonal In (I4/mmm, JCPDS No. 65-7421). According to the Cu-In binary phase diagram,39,40 the equilibrium phases for a composition of [Cu]/[In] = 0.84 are as follows: Elemental indium coexists with the η-phase (Cu2In) from room temperature to 157 °C while the indium-rich liquid phase coexists with Cu11In9 from 157 °C to 310 °C and with the η’-phase (Cu2In) above 310 °C. Under the preparation conditions in the present study, i.e., at 0 °C, however, the prepared nanoparticles are likely in nonequilibrium state in part due to a lack of kinetic energy, which may be how CuIn is often identified in the literature.41-45 The coexistence of CuIn/Cu2In/In has also been reported by Chang et al.35 and Liu et al.46 who employed similar synthetic methods to this work. It is 10


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worth noting from the XRD pattern that, contrary to the seemingly homogeneous size distribution of the nanoparticles of the intermetallic phase mixture (Figure 1), the CuIn and Cu2In phases exhibited quite different peak broadening features. This effect is clearly seen in Figure 2b where the most intense reflections from both phases, i.e., the (200) plane for CuIn and the (102)/(110) planes for Cu2In are magnified along the 2θ-axis. It is recognized that the (200) peak of CuIn is approximately 3 times narrower than the (102)/(110) peak of Cu2In. The lower bound of the average particle sizes were estimated with the peak’s full width at half maxima (FWHM) according to Scherrer’s equation, and the CuIn and Cu2In phases were found to have particle sizes of DCuIn = (56 ± 6) nm and DCu 2 In = (17 ± 3) nm, respectively (The diffraction intensity for indium phase is very low suggesting that indium exists in a minor amount. Therefore, its particle size is not reported here due to the large uncertainty

0 20

30

40

50

2θ /

ο

60

70

C u 2In (102)/(110)

C uIn (200)

C u 2 In (212)/(300)

In (65-7421)

C uIn (213)

C uIn (365)/(192) C u 2 In (202)

C uIn (200)

2000

C uIn (102)

C u 2In (101) In (101)

Intensity / a.u.

4000

CuIn (35-1150) Cu2In (65-704)

C u 2In (112)

(a) 6000

C u 2 In (102)/(110) C uIn (134)

arising from the low diffraction intensity.).

(b) Intensity / a.u.

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35

41 ο 42

2θ /

43

Figure 2. X-ray diffraction patterns of the as-synthesized Cu-In alloy nanoparticles; (a) overall XRD pattern and (b) the most intense peaks from the CuIn and Cu2In phases, respectively, enlarged along the 2θ-axis. Note that the numerals in parentheses in the legend will henceforth denote the JCPDS card numbers. 11



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To investigate the detailed morphology and phase distribution of the bimetallic alloy mixture, as-synthesized nanoparticles were further investigated by using TEM after sonication in ethanol to break up agglomerates into primary particles. The results are shown in Figure 3. In Figure 3a, the Cu-In phase mixture visibly consists of two distinctive nanoparticles having different sizes; the larger particles have diameters in the range from 40 nm to 100 nm, and the smaller particles have diameters of ca. 15 nm. A semi-quantitative compositional analysis by EDS on two distinctive spots for the large and small particles (designated by spot 1 and 2 in Figure 3a) revealed that the atomic ratios of both particles are [Cu]/[In] = 1.1 ± 0.1 and [Cu]/[In] = 1.9 ± 0.2, respectively. The former (large particles) and the latter (small particles) can be reasonably assigned to the CuIn and Cu2In phases, respectively. Figure 3b shows high resolution images for the area marked by the dotted circle in Figure 3a. It is noted that the area encompassing spot 1 that corresponds to the composition close to CuIn was found to be in an amorphous state, which is presumably a consequence of sonication (see XRD pattern for nanoparticles after sonication in Figure S1). On the other hand, the area designated by the circle exhibited a crystalline nature as shown in Figure 3b. From the real and the reciprocal (Fourier transform) images, the spacing of the lattice fringes was estimated to be (0.306 ± 0.003) nm, which corresponds to the (101) plane of hexagonal η-phase (Cu2In). The (102) plane with d-spacing of (0.216 ± 0.003) nm was also identified in the Fourier transform as depicted in the inset of Figure 3b. Both the (102) and (101) planes are consistent with the two most intense reflections for Cu2In in the X-ray diffraction pattern (Figure 2), which validates that the nanoparticles of ca. 15 nm diameter are Cu2In.

12


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Figure 3. TEM investigation on Cu-In intermetallic nanoparticles; (a) transmission electron micrographs for the Cu-In alloy phase mixture and (b) high resolution TEM image for a Cu2In nanoparticle.

° Considering the standard reduction potentials of Cu ( ECu = 0.342 V), In ( EIn° 3+ /In = 2+ /Cu

° 0.338 V), and BH−4 ( EB(OH)

3 /BH 4

= -0.481 V),47 copper has a higher reduction potential with

respect to indium in BH−4 -containing solution. Thus, the formation of copper nuclei is more probable during the initial stage of the reaction and a single primary nucleation events are predominant whereas the much lower reduction potential of indium may lead to a slower multi-nucleation events.34,48 As a consequence, indium nuclei tend to have a wider size distribution.34 Moreover, the Gibbs free energy for the formation of Cu2In is more negative than that of CuIn in the current binary system because Cu2In is the thermodynamically more 13



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stable phase.39,40 Consequently, the critical nucleus size, r * , of Cu2In should be smaller than that for CuIn if the surface energies of both particles are comparable due to the equation,

r * = −2γ / ∆GV , where γ and ∆GV denote the surface free energy per unit area and the bulk free energy change associated with nucleation per unit volume,49 respectively. Thus, the larger initial size of CuIn nuclei together with the higher diffusion coefficient of indium ions in the solution34 facilitates the growth of CuIn particles. Even if the nucleation and growth mechanisms of Cu-In alloy nanoparticles seem to vary in different solvents,35 the thermodynamic argument above may provide a qualitative explanation for the different morphologies of CuIn and Cu2In nanoparticles prepared in this work.

3.2. Characterization of the selenized and sulfurized films The Cu-In alloy precursor films were prepared on Mo-sputtered soda-lime glass by doctorblading the colloidal ink consisting of CuIn, Cu2In, and In nanoparticles, and their phase, composition and microstructural evolution upon chalcogenization at various temperatures were investigated with XRD, Raman spectroscopy, EDS, SEM and AES depth profile measurements. Figure 4 depicts the X-ray diffraction patterns of the thin films prepared at various temperatures in the range of 100 < T / °C < 550 by selenization of the Cu-In alloy precursor films for 30 min. It is first recognized that the major reflections from the (200) and (134) planes of metastable CuIn at 34.7 ° and 43.4° diminished at 100 °C when compared with those for the as-synthesized nanoparticles (Figure 1) and disappeared completely at 200 °C leaving only the Cu2In phase (at 29.7 ° and 42.4°). At 300 °C, partial oxidation of the Cu-In alloy film was observed, which is attributed to negligible selenium supply as the temperature of the Se pellets is still below the melting point (221 °C) in the experimental setup of the 14


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present study (see Section 2.2). As the oxygen partial pressure inside the reaction chamber, log(Po2/atm) ≈ -30 measured by YSZ oxygen sensor, is higher than the equilibrium oxygen partial pressure for In/In2O3, log(Po2/atm)= - 45.1 e.g., at 300 °C while being lower than that of Cu/Cu2O, log(Po2/atm) = -23.3,50 indium in the Cu-In alloy was oxidized to In2O3 accompanying concomitant segregation of excess copper. The CuInSe2 phase in the chalcopyrite structure was first identified at 400 °C where the selenium vapor pressure was 7 × 10-5 atm. Impurity phases such as In2O3 and Cu were finally eliminated at 500 °C resulting in phase-pure CISe aside from elemental selenium, which was detected only at the surface and thus was presumed to form by condensation of selenium vapor during cooling. The Se condensation can be possibly avoided or minimized by reducing cooling rate or removing Se(g) supply during cooling. The intensity of the XRD peak at 40.8 ° in Figure 4 corresponds to the Mo (110) plane, which originates from the substrate, and its intensity decreases with increasing temperature up to 550 °C, indicating that the dense CISe layer is formed and attenuates the incident X-ray. In the Raman spectra (Figure S2), only an intense peak for the A1 vibrational mode and weak peaks for the B2/E modes of chalcopyrite CISe51,52 were found. No traces of impurity phases such as ordered defect compounds53,54 or CuxSe55 were detected within the limit of experimental precision.

15



C uIn (35-1150) ● C u 2 In (65-704) ▲ In 2O 3 (6-416) ▽ C u (4-836) ◆ M o (42-1120) ★ Se (51-1389) ▼ C uInSe 2 (87-2265)

Intensity / a.u.

10000

550 ℃

5000

▼

▼ ★▼ ▼

500 ℃

◆

▼ ▼

▲ ▲

▲

▽

▲ 300 ℃

200 ℃ 100 ℃

0 30

▲

40

▼

400 ℃◆

▼

●

●

20

▼

★

▼

50

o

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60

70

80

2θ / Figure 4. XRD patterns of the thin films printed using the Cu-In alloy colloidal precursor after selenization at different temperatures for 30 min.

The compositional change investigated by EDS is quite consistent with the phase evolution elucidated by XRD as follows: As seen in Figure 5, the mole fraction of selenium with respect to cation concentrations, [Se]/([Cu]+[In]), rises notably to 0.72 at 400 °C, which is the point where the CISe phase first appeared in the XRD pattern and indicates that the selenization of the Cu-In alloy proceeded to such an extent. The selenium content increases with increasing temperature and seems to be saturated at around 500 °C (The composition of the CISe film selenized at 500 °C was confirmed by EPMA as shown in Figure S3.). At the same time, the oxygen content decreases due to the exchange reaction between oxygen and selenium. On the other hand, the increase in oxygen concentration with increasing temperature in the range below 300 °C is attributed to low selenium vapor pressure as 16


Page 17 of 40

previously mentioned regarding In2O3 formation.

0.95

1.2

Mole fraction

selenization 0.90

[Cu]/[In]

1.0

[Se]/([Cu]+[In])

0.8

[O] 0.6 0.4

0.85

Mole fraction

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


0.2 0.0 0.80

0

100

200

300

400

500

600

O

Ta / C

Figure 5. Chemical compositions of the thin films based on Cu-In alloy nanoparticles on Mo/SLG substrate measured by EDS as a function of annealing temperature under selenium-containing atmosphere. Note that [k] stands for the atomic percent of the elements k with the total of all k normalized to 1. The solid curves are for visual guidance.

It is emphasized here that there was little amount of carbon residue detected in the CISe films by virtue of employing the binder-free colloidal precursor in a volatile solvent (ethanol). A thick carbon layer, which may remain in the CIGSe absorber film due to incomplete decomposition of organic substances, leads to high series resistance in solar cells and/or poor adhesion at the CIGSe/Mo interface.11,30,56 To solve this issue, an additional oxidation step is usually performed to burn out the organic substances,15,18,57 or a hydrazine-based solvent in which various binary metal selenides/sulfides and elemental chalcogens are soluble is employed to exclude the use of carbon-based chemicals.13 However, the former may increase the complexity of fabrication procedure while the wide application of the latter is limited due to the toxic and explosive nature of hydrazine. 17



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Page 18 of 40

Figure 6 shows the surface morphologies (a-d) and cross-sections (e-f) of the thin films corresponding to those samples for XRD and EDS measurements. There was little recognizable change in the microstructure of the films annealed at less than 300 °C compared to the as-coated film. Thus, the typical surface and cross-sectional morphologies of the film annealed at 200 °C are represented in Figures 6a and 6e. At 400 °C, a liquid-like microstructure developed at the film’s surface (Figure 6b). The thin film underwent noticeable re-crystallization and grain growth at 500 °C leaving a small number of pores at the surface (Figure 6c). The grain growth proceeded further at 550 °C, and almost all surface pores were eliminated (Figure 6d).

Figure 6. Surface morphologies and cross-sections of the thin films based on Cu-In alloy nanoparticles after selenization at (a, e) 200 °C, (b, f) 400 °C, (c, g) 500 °C, and (d, h) 550 °C for 30 min.

It is worth mentioning here that the CISe films have a layered structure at above 400 °C consisting of a porous bottom layer with particle sizes less than 200 nm and a dense top layer with much larger grains (Figures 6f-6h); this phenomenon has often been observed in the 18


Page 19 of 40

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literature.22,28,34,58-60 Comparing the morphological evolution with the compositional change shown in Figures 5 and 6, one can figure out that the densification of the top layer is concurrent with an increase in [Cu]/[In] ratio with increasing temperature above 300 °C. Due to attenuation of the incident beam with increasing depth, the X-ray signal from atoms near the CISe/Mo interface contributes less to the total EDS intensity than that from those atoms near the surface. Thus, the increase of the [Cu]/[In] ratio over that of the precursor film near the CISe surface (top layer) is likely accompanied by a lower [Cu]/[In] ratio near the CISe/Mo interface region (bottom layer). This implies that the formation of a bi-layer structure as shown in Figure 6 is related to cationic unmixing61 induced by the different diffusivities of cations62,63 under a chemical potential gradient, combined with the compositional tolerance toward a solely Cu-poor regime64 and the role of the Cu-containing binary phase in grain growth.65 As far as the CISe layer is concerned, the sample selenized at 550 °C looks more promising for PV applications than those prepared at lower temperatures in that it has a highly dense and crystalline microstructure (with a grain size of ca. 1.5 µm), which is an important feature of device-quality absorbers. It should be pointed out, however, that a selenium partial pressure as high as 0.008 atm resulted in the formation of an 1.5 µm thick MoSe2 layer at 550 °C at the cost of a 0.5 µm thick Mo layer (Figure 6h). An appropriate thickness of the MoSe2 layer is known to bring about better adhesion and semi-ohmic contact at the Mo/CIGSe interface,66,67 whereas a MoSe2 layer that is too thick causes a high series resistance and hence deteriorates the photovoltaic performance.68 Therefore, we further investigated the photovoltaic properties as well as the elemental distributions (by AES) across the thickness of the films that were chalcogenized at 500 °C; these measurements are discussed below. 19



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Page 20 of 40

In the case of sulfurization, the phase, composition, and morphology change of the Cu-In alloy films are rather simple in comparison with selenization. As shown in Figure 7, the precursor film consisting of CuIn and Cu2In nanoparticles was converted into CuInS2 (CIS) at 200 °C by employing an 1% H2S-99% Ar gas mixture. The compositions of CIS films measured by EDS as a function of annealing temperature (Figure 8) revealed that the replacement of oxygen with sulfur took place even at 100 °C and the composition of the CIS film was almost saturated at 300 °C with no significant changes upon further increase of the temperature. As the sulfurization temperature increased, the (112) reflection for CIS at 28.0 ° became stronger in general and the peak simultaneously narrowed to a little extent, which suggests slight grain growth. Indeed, Figure 9 shows that the grain growth of the CIS films was not as pronounced as that of the CISe films (see Figure S4 for direct comparison.). Although small grains ranging from 100 nm to 300 nm in size may have formed in the CIS phase, most of the surface pores were eliminated by sulfurization at temperatures over 500 °C. It is known that excess Cu acts as a fluxing agent in CuInS2 by forming CuxS.7 In this regard, the Cu-rich composition is deliberately maintained during the grain growth period in the socalled 3-stage co-evaporation method.8 Thus, a slight difference inadvertently induced in the composition of the precursor films, i.e., [Cu]/[In] = 0.84 for selenization and 0.75 for sulfurization, may have had a degree of influence on the grain growth mechanisms of both the CISe and CIS films. However, the remarkable difference in microstructure between CuInSe2 and CuInS2 as shown in Figures 6, 9, and S3, has also been found in samples prepared from precursor films with the same excess-Cu composition, [Cu]/[In] = 1.05.58 Accordingly, it can be inferred that the selenium plays a critical role in formation of top dense layer with large grains of CISe films. It is finally noted that unlike selenization, sulfurization did not induce the formation of MoS2 or a layered structure, and instead it maintains the 20


Page 21 of 40

homogeneous microstructure of CIS films. The change in [Cu]/[In] ratio upon sulfurization was less distinct than that from selenization.

C uIn (35-1150) ● C u 2 In (65-704) ◆ M o (42-1120) ▼ C uInS 2 (42-1432)

Intensity /a.u.

10000

550 ℃

500 ℃

400 ℃

5000 ▼

300 ℃ ▼

200 ℃

●

●

▼

◆

▼

◆

100 ℃◆

0 20

30

40

50

o

60

70

80

2θ / Figure 7. XRD patterns of the thin films printed using a Cu-In alloy colloidal precursor after sulfurization at different temperatures for 30 min.

0.85

1.2 [Cu]/[In] 1.0

Mole fraction

[S]/([Cu]+[In]) [O] 0.80

Mole fraction

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


0.8

sulfurization

0.6 0.4

0.75

0.2 0.0 0.70

0

100

200

300

400

500

600

O

Ta / C

21



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Page 22 of 40

Figure 8. Chemical compositions of the thin films based on Cu-In alloy nanoparticles on Mo/SLG substrate measured by EDS as a function of annealing temperature under a sulfur-containing atmosphere. The solid curves are for visual guidance.

Figure 9. Surface morphologies and cross-sections of the thin films based on Cu-In alloy nanoparticles after sulfurization at (a, e) 200 °C, (b, f) 400 °C, (c, g) 500 °C, and (d, h) 550 °C for 30 min.

The relative atomic concentration profiles were measured by AES on the CISe and CIS samples that were pre-annealed at 300 °C for 30-60 min and subsequently heated up to 500 °C. The pre-annealing step was found to be effective in improving the adhesion of CIS(e) films to the Mo surface for fabrication of solar cell devices. The results are shown in Figure 10. As for the CISe film (Figure 10a), one can recognize that the concentration of copper decreases stepwise at the position corresponding to an etching time of 20 min while that of indium increases similarly. This stepwise change in atomic concentration is believed to indicate the boundary in the bi-layer structure of the CISe film shown in Figure 6g. Namely, the dense top layer with large grains of CISe was found to have a higher [Cu]/[In] ratio than the porous bottom layer, thus supporting the hypothesis that the bi-layer structure is induced 22


Page 23 of 40

by cationic unmixing. On the other hand, the CIS film, with its homogeneous microstructure, exhibited a uniform distribution of constituent elements throughout the thickness of the film (Figure 10b).

Atomic concentration /%

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


100

100

(a)CuInSe2 300oC/30min, 500oC/30min

80

(b )CuInS2 300oC/1 h, 500oC/30min

80

Mo Mo

60

60

In

S

40

40

In

Se 20 0

Cu 0

20

O 20

40

60

Sputter Time /min

80

0

Cu 0

O 5

10

15

20

25

Sputter Time /min

Figure 10. Relative atomic composition depth profiles measured by AES for the thin films fabricated by printing a Cu-In alloy colloid precursor followed by annealing at 300 °C for 30-60 min and at 500 °C for 30 min; (a) CuInSe2 and (b) CuInS2 films.

3.3. Photovoltaic properties of CISe and CIS thin film solar cells CISe and CIS films that were prepared in the same manner as those for the AES depth profile measurements were integrated into solar cell devices with a chemically deposited CdS buffer layer and sputtered i-ZnO and GZO window layers. Figure 11 shows the current densityvoltage (j-V) characteristics under simulated solar illumination (100 mW/cm2, AM 1.5G) and darkness and the external quantum efficiency (EQE; ηQ) without a bias light for the resulting solar cells. The CISe and CIS devices exhibited active area efficiencies of 3.92% and 2.28%, respectively, with photovoltaic (PV) parameters as summarized in Table 2.

23



(a) 20

(c) 0.5

(b )1.0

ηQ

jSC = 31.0 m A /cm

0.4

O

-20

jill,CISe, 500 C

2

accordin g to Eq .(1)

2

O

jdark,CISe, 500 C

0.2

O

jill,CIS, 500 C

-40

jdark,CIS, 500 C

0.0

0.2

0.4

jSC = 7.2 m A /cm

O

0.6

0.0

2 2

0.6

[ln(1-ηQ)Eph ] / eV

CISe CIS

0.8 0

400

600

0.10

0.4

0.08

0.3

0.06

0.2

0.04

0.1

0.02

2

800 1000 1200

0.0 0.9 1.0 1.1

λ / nm

V/V

CIS

CISe

Mo/CISe or CIS/CdS/i-ZnO/GZO/Al

Mo/CISe or CIS/CdS/i-ZnO/GZO/Al

j / mA/cm

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

Page 24 of 40

0.00 1.4 1.5 1.6

Eph / eV

Figure 11. Photovoltaic properties of the solar cells employing CuInSe2 and CuInS2 thin films corresponding to those in Figure 10; (a) j-V characteristics under AM 1.5G illumination and in the dark, (b) external quantum efficiencies ηQ without a bias light under short-circuit conditions, and (c) [ln(1- ηQ)⋅Eph]2 vs. Eph curves near the band edge regime for estimating band gap energies of the thin films.

Table 2. Photovoltaic parameters evaluated from the j-V curves for the solar cells fabricated with CISe and CIS thin films annealed at 500 °C for 30 min. VOC

jSC

/V

/ mA/cm2

CISe, 500 °C

0.304

30.7

CIS, 500 °C

0.443

9.4

sample

A

η

/ cm2

/%

0.42

0.347

3.92

0.55

0.312

2.28

FF

As seen in Figure 11a and Table 2, the short-circuit current density of the CISe device (jSC = 30.7 mA) is comparable to that of a CISe solar cell based on colloidal precursors exhibiting an 8.2% conversion efficiency (jSC = 33.7 mA/cm2, VOC = 0.44 V, FF = 0.55).29 The jSC can also be estimated by the EQE when measured under short-circuit conditions (Figure 11b) combined with the solar spectral irradiance, Iλ (in Wm-2nm-1, referred to ASTM G-173) according to 24


Page 25 of 40

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jSC = ∫ e ⋅ I λ ⋅

λ hc

⋅η Q ⋅ dλ ,

(1)

where e, λ, h, and c denote the elementary electron charge, the wavelength of the photon, Plank’s constant, and the velocity of light, respectively. The estimated jSC for the CISe device from eq 1, 31.0 mA/cm2, was in good agreement with that measured from the j-V characteristics. Despite the relatively high overall photocurrent, the ηQ for the CISe device drops significantly at wavelengths longer than 600 nm, which suggests that the collection efficiency of the minority charge carriers still needs to be improved. This poor collection efficiency at longer wavelengths due to the short minority carrier diffusion length, LD ( ∝ D ⋅ τ where D and τ are the diffusion coefficient and the life time of the electron, respectively), could be attributed to high series resistance and/or a high recombination rate in the porous bottom layer where a lot of intergranular interfaces and tiny particle surfaces may act as trap and/or recombination sites. Therefore, achieving a homogeneously well-developed microstructure in the CISe film throughout its thickness would lead one to expect a higher jSC. On the other hand, the open-circuit voltage (VOC = 0.304 V) and fill factor (FF = 0.42) for the CISe device are much lower than those referred to previously,29 which results in inferior overall performance of the device in the present state. In comparison with the highly efficient CISe solar cell fabricated by co-evaporation techniques exhibiting 14.5% total area efficiency (jSC = 41.1 mA/cm2, VOC = 0.491 V, FF = 0.719),4 however, it is obvious that the wetprocessed CISe devices typically suffer from lower VOC and FF. The band gap energy Eg can be evaluated from the ηQ data near the band edge through the relationship between the absorption coefficient α and the photon energy Eph for a direct transition as αEph ∝ (Eph- Eg)1/2 along with the band edge approximation for α versus ηQ, i.e., α ∝ ln(1- ηQ).69 The linear regression on the plot of [ln(1- ηQ)⋅ Eph]2 versus Eph in Figure 11c thus yielded band gap 25



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Page 26 of 40

energies of (1.00 ± 0.03) eV for CISe and (1.46 ± 0.07) eV for CIS absorber films, which are fairly close to the reported values for single crystals, 1.04 eV70 and 1.53 eV,71 respectively. The small difference of only 40 to 70 mV between the reported band gap energies and those measured in this work cannot explain the low VOC values of the present devices. It can be recognized in Figure 11a that, by shifting the j-V curve upward, one cannot obtain an identical plot to the jdark-V characteristics over the entire range of applied voltage. The deviation could be interpreted as the consequence of illumination-driven changes in the PV parameters (such as series resistance Rs, shunt conductance Gsh, diode ideality factor n, and saturation current density jo) and/or of the applied bias-dependent photocurrent collection. It has been argued that the latter is the most likely explanation for the discrepancy in various types of thin film solar cells.69 Provided that the photocurrent density is a function of the applied voltage, i.e., jph= jph(V),69,72 the j-V characteristics under illumination can be described using jph(V) and the diode equation for the dark current density jdark, while taking the parasitic resistances into account as69,73   e(V − jRs A)   Gsh jill = jdark (V ) − jph (V ) = jo exp V − jph (V ) .  − 1 + nkT   A  

(2)

Here, k is Boltzmann’s constant and A is the surface area of the solar cell. By introducing considerable simplifications such as an assumption that jph is much higher than the shunt current density GshVOC/A, one can reach an expression for the open circuit voltage as74

VOC ≈

nkT  jPh  ln + 1 . e  jo 

(3)

Equation 3 indicates that the open-circuit voltage is mainly determined by the saturation current density because jo varies by orders of magnitude depending on the device quality and the band gap of the absorber.74,75 In reality, however, it is far from simple to deconvolute n and jo from non-linear least square (NLLS) fitting on the j-V curve due to their strong 26


Page 27 of 40

interdependence. Instead, the PV parameters Gsh, Rs, n, and jo can be determined from the j-V curves obtained in the dark using a standard diode analysis for thin film solar cells.69 In the plot of dj/dV vs. V (Figure 12a), the shunt conductance, Gsh, was evaluated from the constant dj/dV values in the range of V ≤ 0 based on the assumption that a linear shunt current

predominates over the diode current where V ≤ 0. The series resistance Rs and the diode ideality factor n were estimated from the y-intercept and the slope of a plot of dV/dj versus 1/(j-GshV/A) in the high bias regime (Figure 12b) according to −1

G dV nkT   = Rs A +  j − sh V  , dj e  A 

(4)

as can be readily shown by differentiating the diode current in eq 2 with the approximation that (Gsh/A)⋅(dV/dj) is negligibly small. A semi-logarithmic plot of (j-GshV/A) vs. V-jRsA (Figure 12c) enables us to extract the saturation current density jo and n from the y-intercept and the slope, respectively. The resulting PV parameters are listed in Table 3 where the ideality factor n is represented by the average values obtained from the two procedures

(b ) 40

2

30

1

2

0.70 m/Scm

0 Gsh/A =

-1 -0.2

2

0.07 mS/cm -0.1

0.0

V/V

CISe CIS 0.1

20

slope/mV = nkT/e

10 2

0.2

0 0.0

y-intercept/Ωcm = RsA 0.1

(c)

4 2

2

CISe CIS

2

Gsh/A =

dV/dj / Ωcm

2

(a)3

ln(j-GshV/A / mA/cm )

described above.

dj/dV / mS/cm

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


0.2

0.3 2

0.4

1/(j-GshV/A ) / cm /mA

CISe CIS

0 -2

-1

slope/V = e/nkT

-4 -6 -8

-10 0.0

y-intercept 2 = ln(jo/mA/cm ) 0.1

0.2

0.3

0.4

0.5

V-jRsA / V

Figure 12. A diode analysis on CISe and CIS solar cells; (a) dj/dV vs. V curves for Gsh evaluation, (b) dV/dj vs. 1/(j-GshV/A) plot for the determination of Rs and n, and (c) 27



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Page 28 of 40

semi-logarithmic plot of (j-GshV/A) vs. V-jRsA for determination of n and jo.

Table 3. Shunt conductance Gsh, series resistance Rs, ideality factor n, and saturation current density jo of CISe and CIS devices extracted from Figure 12. sample

Gsh/A / mS/cm2

RsA / Ωcm2

n

jo / mA/cm2

CISe

0.70 ± 0.03

1.2 ± 0.3

3.4 ± 0.3

(4.9 ± 0.3) × 10-2

CIS

0.07 ± 0.01

9.6 ± 0.1

1.9 ± 0.2

(4.5 ± 0.2) × 10-4

Despite the ambiguity of the best fit in Figure 12c, the estimated parameters provide insight into the inferior performance of the present devices, in particular, the low VOC and FF. According to Contreras et al.,74 the saturation current density of CIGSe solar cells increases with decreasing band gap below 1.16 eV and reaches ca. 10-4 mA/cm2 at Eg = 1.0 eV. Comparing this value with the measured jo for our CISe device (4.9 × 10-2 mA/cm2), one can infer that the poor VOC of the present CISe device is associated with a much higher jo. The large values of jo as well as n (3.4 for CISe) may be ascribed to the high recombination rate at the defective p-n junction i.e., CISe/CdS interfaces. In practice, the rough surface morphology together with the large number of surface states in the absorber films could result in such an inferior p-n junction.74 The jo and n values are known to significantly affect the fill factor as well72 in such a way that the reduced fill factor is associated with increased values of jo and n. Obviously, the higher values for Gsh/A (0.70 mS/cm2) and RsA (1.2 Ωcm2) for the CISe device compared to those for a 15.5% efficient CIGSe solar cell (Gsh/A = 0.05 mS/cm2 and RsA = 0.2 Ωcm2)69 are also responsible for the low fill factor. The series resistance of the CIS device (RsA = 9.6 Ωcm2) was found to be even higher than that of the CISe cell. We attribute this to lower crystallinity with a much smaller grain size in the CIS absorber as was discussed in Section 3.2. In this regard, it is surmised that the 28


Page 29 of 40

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relatively large discrepancy between the jSC values for the CIS device determined from the jV curve and the ηQ data is thus related to a lower collection probability for the photo-

generated electron under the measurement conditions of EQE than those for the j-V curve. By using the PV parameters described above and the microstructural and compositional characteristics of the absorber as measures by which the absorber quality is evaluated, further improvements in the PV performance of CISe and CIS solar cells prepared from a Cu-In alloy colloidal precursor are underway, particularly for the poor open-circuit voltage and fill factor.

4. SUMMARY AND CONCLUSIONS The Cu-In intermetallic nanoparticles synthesized via NaBH4–assisted reduction method under ambient conditions were found to be in a bimetallic compound mixture consisting of CuIn, Cu2In, and elemental In. TEM analysis in conjunction with EDS and XRD measurements revealed that as-synthesized nanoparticles were composed of CuIn particles of 40-100 nm diameter and Cu2In particles of ca. 15 nm diameter. The Cu-In alloy colloidal precursor enabled vacuum-free formation of CuInSe2 and CuInS2 thin films by employing elemental selenium and H2S(g) as chalcogen sources, respectively. The resulting chalcogenide films were found to be pure in phase after selenization or sulfurization at and above 500 °C and, notably, were shown to be carbon-free by virtue of the use of binder-free colloidal precursor in a volatile solvent. The CuInSe2 films exhibited a typical layered microstructure consisting of a porous bottom layer with particle sizes less than 200 nm and a dense top layer with much larger grains approximately 1.5 µm in size. It was further revealed that the densification of the top layer in CuInSe2 films is accompanied by an increase in the [Cu]/[In] ratio with increasing temperature, which suggests that the formation of a bi-layer 29



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Page 30 of 40

structure is associated with cationic unmixing. The re-crystallization of CuInS2 films was less pronounced than for CuInSe2 films. As a result, a homogeneous microstructure developed in the CIS films with small grains ranging from 100 nm to 300 nm in size. Accordingly, it is suggested that the selenium plays a critical role in grain growth and densification of chalcogenide films. The CISe and CIS solar cell devices employing binder-free Cu-In alloy colloidal precursors exhibited active area efficiencies of 3.92% and 2.28%, respectively. The higher efficiency of the former was due to the relatively high short-circuit current density (30.7 mA/cm2 for CISe vs. 9.4 mA/ cm2 for CIS), which most likely originates from the top dense layer of the CISe film. A diode analysis revealed that the present devices had very high saturation currents, jo, and diode ideality factors, n, in addition to high series resistance, Rs, and shunt conductance, Gsh, when compared to values for highly efficient devices fabricated by vacuum techniques. Such high values of jo and n, which may possibly be attributed to a high recombination rate at the absorber film/CdS junction, resulted in low open-circuit voltages and fill factors and, hence, tantalizing overall PV performances of the devices in the present state. An analysis on the PV parameters (jo, n, Rs and Gsh) and the microstructural and compositional characteristics shed light on how to improve device performance. Further studies are accordingly in progress to realize a well-developed microstructure with large grain size and low porosity throughout the absorber film thickness76 while simultaneously circumventing the formation of a resistive interlayer at the Mo/absorber interface and to improve the junction properties by surface morphology control and/or defect-chemical modification at the absorber/CdS interface.

ASSOCIATED CONTENTS 30


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Supporting Information Comparison of XRD patterns of Cu-In alloy nanoparticles before and after sonication, Raman spectra of the selenized films at different temperatures, EPMA mapping for the film selenized at 500 °C, and comparison of microstructure between the films selenized and sulfurized at 500 °C. This information is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *

Tel. +82-2-958-6710, Fax.: +82-2-958-6649, E-mail: [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work was supported by the program of Korea Institute of Science and Technology (KIST) and the “National Agenda Project” program of Korea Research Council of Fundamental Science & Technology (KRCF).

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Solar Cell with 6% Efficiency. ACS Appl. Mater. Interfaces 2013, 5, 1533–1537.

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