Carbon-Impurity Affected Depth Elemental Distribution in Solution-Processed Inorganic Thin Films for Solar Cell Application.

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Carbon-Impurity Affected Depth Elemental Distribution in SolutionProcessed Inorganic Thin Films for Solar Cell Application Shanza Rehan, Ka Young Kim, Jeonghyeob Han, Young-Joo Eo, Jihye Gwak, Seung Kyu Ahn, Jae Ho Yun, KyungHoon Yoon, Ara Cho, and SeJin Ahn ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.5b10789 • Publication Date (Web): 28 Jan 2016 Downloaded from http://pubs.acs.org on January 31, 2016

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ACS Applied Materials & Interfaces

Carbon-Impurity Affected Depth Elemental Distribution in Solution-Processed Inorganic Thin Films for Solar Cell Application Shanza Rehan,§, ¥ Ka Young Kim, § Jeonghyeob Han,§, ¥ Young-Joo Eo, § Jihye Gwak, § Seung Kyu Ahn, § Jae Ho Yun,§, ¥ KyungHoon Yoon, § Ara Cho* §, ¥ and SeJin Ahn* §, ¥ §

Photovoltaic Laboratory, Korea Institute of Energy Research, Daejeon 305-343, Korea.

¥

Department of Renewable Energy Engineering, Korea University of Science and Technology

(UST), DaeJeon 305-350, Korea Keywords: Non-vacuum, carbon-residue layer, ohmic, diffusion, chalcopyrite, resistance, solution deposition, Photovoltaics.

ABSTRACT: A common feature of the inorganic thin films including Cu(In,Ga)(S,Se)2 fabricated by non-vacuum solution-based approaches is the doubled-layered structure, with a top dense inorganic film and a bottom carbon-containing residual layer. Although the latter has been considered to be the main efficiency limiting factor, (as a source of high series resistance), the exact influence of this layer is still not clear, and contradictory views are present. In this study, using a CISe as a model system, we report experimental evidences indicating that the carbon residual layer itself is electrically benign to the device performance. Conversely, carbon was

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found to play a significant role in determining the depth elemental distribution of final film, in which carbon selectively hinders the diffusion of Cu during selenization, resulting in significantly Cu-deficient top CISe layer while improving the film morphology. This carbonaffected compositional and morphological impact on the top CISe films is a determining factor for the device efficiency, which was supported by the finding that CISe solar cells processed from the precursor film containing intermediate amount of carbon demonstrated high efficiencies of up to 9.15 % while the performances of the devices prepared from the precursor films with very high and very low carbon were notably poor.

1. Introduction

Cu(In,Ga)Se2 (CIGSe)-type solar cells are the most promising among all thin-film photovoltaic technologies because of their high efficiency and stable performance for both terrestrial and space applications. CIGSe has high light absorption coefficient (~ 105 cm-1) and a tunable direct band gap (1.0 - 2.4 eV), which can be controlled by varying the ratios of In/Ga and S/Se.1, 2 The highest efficiencies of CIGSe solar cells of up to 21.7 % have been achieved using vacuum-based 3-stage co-evaporation technique, in which precise in-situ compositional control is possible at every stage of film deposition.3 However, the high manufacturing cost of the vacuum-based deposition limits its large-scale commercial production. In contrast, non-vacuum solution-based deposition techniques offer a simple and cost-effective fabrication option, with efficient material utilization and high throughput.4 The highest efficiency of CIGSe-type solar cells through non-vacuum solution based routes is 15.2 %, achieved by so-called “hydrazine-based process”,5 in which metal selenide or sulfide


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powders were used as the precursor materials and hydrazine as the solvent. However, hydrazine is an extremely toxic and explosive solvent, and the entire process must be performed in a glove box6-8 which renders this process far from commercialization. As a result, hydrazine-free solution-based processes have been extensively studied for over a decade.7-13 In line with these efforts, the remarkable efficiencies of CI(G)Se solar cells have been reported using a variety of manufacturing techniques. Some of the representative results are summarized in Table 1.

Table 1. Summary of representative “hydrazine-free” solution processes for CI(G)Se-type solar cells

Carbon-

Absorber

Precursor

Solvent / Organic additives

Eff. (%)

Ref.

CIGSSe

Crystalline nanoparticle

Hexanethiol

15

McLeod et al.11

CISe

Multi-phase nanoparticle

Ethylene glycol, ethanol / 8.2 Polyvinylpyrrolidone (PVP)

Jeong et al.12

CISe

Hydroxide nanoparticle

Uhl et al.13

CISe

Amorphous nanoparticle

Butyl glycol acetate / 8.0 Polymeric dispersant (NUOSPERSE, type FX) Methanol / 7.94 Monoethanolamine (MEA)

CIGSe

Molecular ink

Ethanol, 1-2-propanediol

7.7

Uhl et al.15

CIGSSe

Molecular ink

Water

10.7

Septina et al.16

CIGSSe

Molecular ink

Water

10.5

Hossain et al.17

CIGSSe

Molecular ink

Methanol / Polyvinyl acetate 8.8 (PVA)

containing

Carbonfree

Ahn et al.14

Park et al.18


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A technological issue of long debate in the hydrazine-free solution processes is the existence of carbon impurity in the final film. CI(G)Se-type films with a double-layered structure consisting of the top dense crystalline CI(G)Se layer, which is believed to dominate the electro-optical properties of the entire film, along with the bottom carbon-containing, often porous, residue layer, has been frequently reported.11-15,

19, 20

This carbon-containing residue layer (CRL) is

commonly reported to be originated from incomplete removal or decomposition of the various organic additives, such as chemical modifiers, sacrificial ligands and surfactants for solubility, coatability and reaction kinetics control.21 To date, the formation of the CRL is generally considered as one of the main reasons for limiting device efficiency in hydrazine-free routes because it is reported to have the following characteristics: acts as a resistive element thereby increasing the series resistance (Rs);22-24 inhibits the merging of nanoparticles and the crystal growth,25 causes poor adhesion with Mo substrate,24,

25

and acts as recombination centers.26

Based on this belief, most of the researches in this field have been focusing on the removal of this CRL for several years. In fact, ‘carbon-free’ CI(G)Se-type films have been successfully demonstrated mainly by two routes; a spray-pyrolysis approach using aqueous ink consisting of inorganic metal-salts and thiourea16, 17 and a multi-step annealing approach.18 The aqueous ink used in the former does not include any organic additives ensuring the carbon-free final films while it typically requires rather long deposition time to obtain thick films due to the low viscosity of the ink.4 The latter approach was reported by Park et al.18 in which they used a molecular ink containing organic additives for viscosity adjustment but the carbon-containing precursor film was subjected to multi-step heating process of aggressive air-oxidation to completely remove carbon followed by high temperature sulfurization in H2S gas and selenization with Se vapor. However, in spite of the successful demonstration of these carbon-


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free devices, the efficiency level (10.7 % from spray pyrolysis and 8.8 % from multi-step heating process, Table 1) is still not as high as the maximum efficiency of the carbon-containing CIGSSe solar cells (15 %, Table 1). Therefore, this comparison raised doubts regarding the actual effects of the CRL on the device performance, i.e., is the CRL really a main efficiency limiting factor? Recently, an interesting paper was published claiming that the CRL is benign to device performance, as it is electrically conductive with low bulk resistance, ensures favorable hole transport at absorber/CRL interface and forms an ohmic contact with Mo.27 Although this conclusion was drawn for the case of nanoparticle-based CuZnSn(S,Se)4 (CZTSSe) solar cells, we believe that the nature of the CRL is universally equivalent in most of the solution-processed CIGSe and CZTSe systems. Thus, there is still a lack of consensus regarding the impact of the CRL in CIGSe-type solar cells prepared by non-vacuum routes, and little direct experimental observation exists on this issue. In addition to the device performance, if we consider that the top CIGSe layer is formed by the reaction between metal ions vigorously diffusing to the surface of the precursor film and Se vapor, then the neighboring carbon atoms must affect this diffusion behavior of metal ions and hence the properties of the top CIGSe film. However, no such report was found in the literature addressing this issue. Therefore, the purpose of this research is to provide insight regarding the impact of carbon existing in the precursor film on the properties of the final absorber layer and device performance of CISe solar cells prepared using a hydrazine-free solution-based route. Experimental results revealed that CRL itself is electrically conductive and forms an ohmic contact with both CISe and Mo films, indicating that CRL is not a determining source of RS. However, carbon significantly affected the diffusion behavior of metal ions during selenization


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by selectively hindering the diffusion of Cu, while both In and Na easily diffused through the precursor film matrix. The importance of this non-uniform diffusion of metal ions is that it determines the depth compositional uniformity. An additional impact of carbon was in enabling the top CISe films to become more compact and flatter with increasing carbon content in the precursor films. The carbon-affected compositional and morphological characteristics of the top CISe film significantly influenced device performance, as evidenced by experiments in which the effects of carbon content in the precursor film on the device performance were investigated. CISe solar cells processed from the precursor film with an intermediate amount of carbon demonstrated high efficiencies of up to 9.15 %, while both the carbon-free devices and the devices containing very high amount of carbon exhibited very poor performance. The lower performance of the carbon-free and the very carbon-rich devices were attributed to rough morphology and extremely undesirable depth elemental distribution of the absorber film, respectively, which are affected by the carbon content in the precursor film.

2. Results and Discussion

2.1. Characteristics of typical CISe thin film prepared via the molecular ink approach As a model system to address the impact of carbon, a direct solution coating technique based on a molecular precursor ink was used to form the precursor film. This technique was used in particular because this route provides the simplest precursor system in which every metal ion chelated with organic ligands is uniformly dispersed in the ink. Thus, we can avoid the additional complexity of a non-uniform distribution of carbon, as in the case of nanoparticle-based routes, in which only the surface metal ions are chelated and those in the inner part of nanoparticles are


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not. Among the various recipes suggested in the literatures, a metal salt/alcohol ink-based route, in which metal acetates chelated with ethanolamine are dissolved in alcohol solvent, was chosen. This recipe is almost the same as that reported by Ahn et al.,28 the only difference being the use of 2-methoxyethanol in this study instead of ethanol to ensure milder evaporation of the solvent during air drying to reduce crack formation in the precursor films. CISe films prepared via this route demonstrated a characteristic double-layered structure of a densely packed crystalline layer on a bottom fine-grained film (Figure 1(a) and (b)). A 400 nm thick MoSe2 layer was also formed by the direct contact of elemental Se with Mo during selenization (Figure 1(b)). Raman analysis and AES depth profile (Figure 1(c) and (d), respectively) revealed that the upper layer is CISe with trace amount of ordered vacancy compound (OVC) while the bottom part is CRL with enrichment in Cu. (Due to the standard-less measurement mode of AES and the porous nature of the bottom layer, the AES depth profile was considered only semi-quantitatively in this work.) The presence of OVC phase is believed to cause a rather positive effect on the device performance due to its hole-blocking nature, which intern reduces surface recombination.29, 30 The Cu enrichment in CRL will be addressed more in detail when the growth mechanism of the CISe film is discussed. The carbon in the bottom CRL was found to exist in the amorphous state as evidenced by XRD pattern of the same sample in which no carbide peaks were detected (Figure S1 in supporting information). Note that the Raman scattering data shown in Figure 1(c) provides information only for the top-most part of our film due to the penetration depth of approximately 200 nm of the Ar-ion laser light used. Thus, to obtain further structural information throughout the entire film depth, we performed additional Raman analysis for the CISe/CRL/(MoSe2)/Mo film after performing inclined polishing, as shown in Figure 2(a). This inclined polishing lengthens the lateral distances for the


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Raman measurement, thereby providing full structural information from the top-most part of the CISe layer to the MoSe2 layer. At first glance, each layer of CISe, CRL and MoSe2 is well distinguished by CISe-related Raman peaks in Figure 2(b) and MoSe2-related peaks in Figure 2(d). For the CISe part (Figure 2(b)), note that all the Raman peaks correspond to the CISe phase, except a trace of OVC that is found only near the CISe/CRL interface (point 3). No distinct Cu-Se or In-Se binary phase-related peak is visible in Figure 2(b), indicating that the top layer can be considered as a CISe without any binary selenide phase. Figure 2(c) demonstrates that Raman signal from CRL is in general broad and low in intensity, showing weak signals of Cu-Se and CISe related peaks. This result indicates that CRL consists of almost amorphous or low crystalline CISe and/or Cu-Se phases mixed with amorphous carbon.

Figure 1. Characteristics of a typical CISe film; (a) planar SEM image, (b) cross-sectional SEM image, (c) Raman scattering spectra and (d) AES depth profile.


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Figure 2. (a) Locations where the Raman spectra are taken in a CISe/CRL/Mo sample after inclined polishing; Raman spectra obtained from (b) top CISe layer, (c) CRL and (d) MoSe2 layer. 2.2. Optical and electrical properties of the CRL The characteristics of our CISe film described above are in general similar to those reported by other researchers.12-14, 28 The following sections are devoted to the understanding of the impact of the CRL on device performance and the carbon-affected diffusion behavior of metal ions during selenization.


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Figure 3. Cross-sectional SEM images of a selenized samples on SLG (a) before and (b) after etching of the top CISe layer. Optical transmittance measurement of the top-layer-removed-CRL was performed to investigate the optical properties of the CRL; the result is presented in Figure 4(a). For comparison, an experimentally measured transmittance curve of a CISe/CRL sample and a calculated transmittance curve of the CISe layer are also presented in Figure 4(a). The latter was required to analyze the optical properties of the top CISe layer, and it was calculated by the following equation; for clarity, the meaning of each transmittance curve is schematically shown in Figure 4(b). %T

%T %T


10

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Figure 4. (a) Transmittance curves and (b) conceptual schematic drawing of (i) CRL, (ii) CISe+CRL and (iii) CISe only cases. Note that (i) and (ii) are experimental data and (iii) is calculated using (i) and (ii). (c) Bandgap of CISe film estimated from (iii) of Figure 4(a). The results revealed that while the transmittance curve of the CISe film, (iii) of Figure 4(a), shows an characteristic absorption edge of a semiconductor with a bandgap of 0.95 eV (Figure 4(c)), which is consistent with the literature,12, 28 the CRL does not have such an absorption edge ((i) of Figure 4(a)). The CRL shows a gradual decrease in %T to the short wavelength direction,


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which is very similar to the transmission curve of the conventional amorphous carbon film synthesized by a pulsed filtered arc deposition using a graphite cathode.31 This result indicates that even though our CRL is mixture of amorphous carbon and low crystalline CISe and/or CuSe phases, its optical behavior is dominated mainly by the nature of amorphous carbon, and it is unlikely that CRL contributes to generating photocurrent when illuminated. The electrical characteristics of the CRL, in terms of the resistance of the CRL itself and its electrical contact with the underlying MoSe2/Mo film, was evaluated by comparing the I-V curves of a Ag/CRL/MoSe2/Mo device and a completed CISe device, with the schematics of each device shown in Figure 5(a) and (b), respectively. While the completed CISe device shows a typical diode-like I-V curve in Figure 5(c), the Ag/CRL/MoSe2/Mo device exhibits a perfect ohmic nature. Furthermore, comparison of the slopes of the linear portions in the curves in the high forward bias region revealed that RCRL is smaller than R(S,full device). These data suggest that the CRL itself is a simple conductor with low resistivity that forms an ohmic contact with the MoSe2/Mo back contact; as a result, the CRL is unlikely the main source of RS of the completed device.


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Figure 5. Schematics of (a) Ag/CRL/MoSe2/Mo device and (b) completed CISe device; (c) I-V curves of the two devices. Another issue to be considered is the contact property between the CRL and the top CISe layer which was addressed by investigating the impedance spectra of the completed CISe device. This experiment was motivated by the idea that if there is any electrical junction barrier impeding the hole transport from CISe to CRL at the CISe/CRL interface, it should be visually detected as an additional capacitive element to the main junction capacitance in the impedance spectra. Figures 6(a) and 6(b) are the Nyquist and Bode plots, respectively, of a completed CISe device measured at various forward biases. The red solid lines shown in Figure 6(a) are perfect semicircles generated according to the simple Randle circuit. In general, all the impedance spectra can be reasonably described by a single semi-circle (Figure 6(a)) with a single critical frequency (Figure 6(b)) within frequency range of 10-1 to 106 Hz, reflecting that the device has seemingly only one main electrical junction, which is believed to exist at the CdS/CISe interface. Note that in Figure


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6(a), the measured spectra exhibit a slight deviation from a perfect semicircle in the very low frequency region; this deviation becomes more prominent as the applied bias increases. One may attribute this inductive loop-like deviation to the CRL in our device. However, Bisquert previously demonstrated that the inductive loop in the impedance spectra at very low frequency was observed in many types of inorganic and organic solar cells, including CdTe, extremely thin absorber (ETA) and standard organic bulk hetero-junction solar cells, although the phenomenon still lacks a proper interpretation.32 As a result, it is not reasonable to simply relate this inductive loop-like behavior shown in Figure 6(a) to the existence of the CRL. To strengthen this point, we performed impedance measurements for vacuum-processed CISe solar cells with efficiencies in the range from 9 to 11 %, in which CISe films were fabricated using a 3-stage co-evaporation technique (Figure S2 in supporting information). The general shape of the Nyquist and Bode plots (Figure S2(a) and S2(b), respectively) of the vacuum-processed CISe devices is clearly very close to those of the solution-processed ones shown in Figure 6. Furthermore, the deviation from a perfect semi-circle at very low frequencies in the Nyquist plot (Figure S2(a)), which is more prominent at higher forward biases at 0.5 and 0.6 V, is also detected even in this case. This result confirms that the inductive loop-like behavior should not be attributed to the CRL itself or to the CRL/CISe interface because vacuum-based CISe devices have no carbon at all. Thus, we could not find any evidence of any additional electrical junction barrier to the main junction in the impedance spectra of our solution-processed device, and hence the CRL/CISe interface is believed to be not causing any hole-blocking behavior.


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Figure 6. (a) Nyquist and (b) Bode plots of a solution-processed CISe device measured at various forward biases. From all the electrical measurement data described above, we suggest that CRL itself is indeed electrically benign to the device performance, which is in stark contrast to the historical claims


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that CRL is a highly resistive layer and a primary efficiency limiting factor,20, 22, 23, 26 whereas our results support the recent report from Wu et al.27 2.3. Carbon affected depth elemental distribution in the selenized film In the previous section, the CRL itself was found to be electrically not detrimental to the device performance. However, as stated in the introduction, carbon existing in the precursor film, in addition to being a residual layer in the final film, must affect the diffusion behavior of metal ions through the carbon-containing matrix during selenization and hence the properties of the top CISe. This issue motivated us to investigate the detailed reaction pathway of the top CISe film formation during selenization with a focus on the influence of carbon in the precursor film. To this end, selenization of air-dried precursor films was interrupted at various temperatures, as shown in Figure 7, and the samples were removed for SEM, AES and SIMS analysis.

(e)

(d) 560 sub. temp. / oC

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(c)

400 (b)

300 200

(a)

30min

time / min

Figure 7. Thermal history of selenization, showing the points (a)-(e), where the process was interrupted and the samples were removed for analysis. The columns 1, 2, and 3 in Figure 8 show the SEM planar view, cross-sectional view, and AES depth profile of elemental diffusion in the precursor film, respectively. The five rows, (a), (b), (c),


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(d), and (e) in Figure 8 stand for the five points designated in Figure 7 at which selenization was interrupted. Point (a) and (b): For (a) and (b) in Figure 8, SEM images show that the film morphology is close to that of the air-dried precursor film as the vapor pressure of Se is not completely developed inside the graphite container at these temperatures, therefore very less or no reaction is occurring at these points. AES depth profiles of these films demonstrate that all the elements are uniformly distributed throughout the entire film. Co-existence of oxygen and carbon in the films indicates that air-drying induced partial decomposition and oxidation of metal-MEA complexes. Point (c): Distinct crystal formation with layer separation appeared to start at 400 °C ((c) of Figure 8) with AES profile clearly showing the newly forming top Cu-In-Se layer which is almost free of carbon. It is notable that at point (c) carbon content of the bottom CRL increased to approximately 20 at. % compared to those at point (a) and (b), which is approximately 15 at. %. This result indicates that the first selenide film starts to form at the film surface by reaction between metal ions (which are diffused from the underlying precursor film) and Se vapors, while carbon moves toward the back of the film, causing its accumulation in the bottom layer. Also note that in the newly forming film, the In content is strikingly higher than the Cu content. Considering that the initial elemental distribution was uniform (AES data of point (a) and (b) of Figure 8) and that a new film forms by diffusion of Cu and In to the top surface, the significantly high In content in the new film indicates that In moves towards the top much faster than Cu in our film. Point (d) and (e): When the temperature increases to 560 °C ((d) in Figure 8), the doublelayered feature becomes more prominent, in which the top layer is a 560-nm thick densely


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packed CISe film and the thickness of CRL is 650 nm. Dwelling at that temperature for 30 min does not significantly change the general morphology of the film ((e) in Figures 8), except for the slight increase in the thickness of top CISe layer and the decrease in that of the CRL. An interesting feature in the AES profiles in Figure 8 is that the more rapid diffusion of In than that of Cu, first observed at 400 °C (point (c)), resulted in significant non-uniformity in the composition between the top and bottom layers as the selenization proceeded (point (d) and (e)). Almost all the In was used to form the top CISe layer, while a large fraction of Cu still remained in the CRL (point (d)). Dwelling for 30 min at 560 °C (point (e)) was found to induce more Cu diffusion from CRL to the top layer. This finding of entrapment of Cu in CRL is consistent with literature.15


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Figure 8. Column 1, 2, and 3, showing planar SEM images, cross-sectional SEM images, and AES depth profiles of elemental diffusion in the precursor film at points designated as (a), (b), (c), (d), and (e), in Figure 7.


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From the results of the morphological and compositional study, we can draw following conclusions on the reaction pathway of our CISe/CRL films; i) The top CISe film starts to form at the surface of the precursor film, where Cu and In (diffused from the bottom layer) react with Se, while carbon moves and accumulates towards the backside. Double-layer formation of the carbon-free top CISe film on CRL can be reasonably explained by this mechanism. ii) Diffusion of In to the top surface to react with Se is much faster than the diffusion of Cu, resulting in the significant compositional non-uniformity in terms of the Cu/In atomic ratio between the top CISe and bottom CRL. In particular, the second conclusion of faster In diffusion is surprising because it is opposite to the general perception that Cu is the most mobile atom in most CIGSe material systems.33-36 The unusually high In mobility in our film indicates that there should be a critical difference in the film matrix (which determines the diffusion behavior of metal ions) of our solution-processed films compared to those reported in the literature. Because the most prominent difference is that our film contains a significant amount of carbon, while the reported vacuum-based CI(G)S films are completely carbon-free, we propose the hypothesis that carbon plays a role in controlling the diffusion behavior of metal ions during selenization. Carbon should selectively hinder the movement of Cu, thereby causing the significant non-uniformity in the depth compositional profile as shown in AES depth profiles of Figures 8(d) and (e). As a first step to prove our hypothesis, we prepared separate Cu- and In-precursor films from Cu-acetate/MEA/2-methoxyethnaol and In-acetate/MEA/2-methoxyethanol inks, respectively. The inks had the same solute concentration as that of the original Cu-In-mixed ink and the film deposition was also performed in a same manner as that for forming the original Cu-In-mixed precursor film, i.e., the same conditions for spin coating and air-drying. AES depth profiles of


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these two films reveal that the Cu-film still contains a significant amount carbon (Figure 9(a)), while the In-film is completely free of carbon (Figure 9(b)), signifying that the carbon in the original Cu-In-mixed precursor film (Figure 8(a)) is mostly from Cu-acetate/MEA complex and, more importantly, that Cu ions are more strongly bonded to this chelating agent than In ions in our ink system. This result is consistent with the report from Cho et al., who reported that chelate complex formed between Cu and MEA is more stable than that between In and MEA.37 Furthermore, this result can explain our finding of the hindered movement of Cu compared to the much faster diffusion of In to the top surface during selenization. Because a significant amount of Cu is still trapped or captured by carbon, even after the air-drying step, supply of this Cu from the (unreacted) bottom layer to the newly forming top film can occur only by breaking the Cucarbon bonds. In this way, the carbon in the precursor film is believed to play a critical role in determining the relative diffusion behavior of Cu and In as well as the depth compositional uniformity of these elements, and hence the quality and/or property of the final top CISe layer.

100

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C 6

sputter time / min

0

2

4

6

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Figure 9. AES depth profiles of (a) Cu-precursor film and (b) In-precursor film.


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The importance of carbon in the precursor film on the properties of the top CISe film was further verified by the following set of experiments, in which two extreme conditions of the carbon content in the CISe film were intentionally taken. If the carbon in the precursor film indeed affects the diffusion behavior of metal ions in the way described above, it is expected that in the precursor film with higher carbon content there will be greater number of initial Cu-MEA bonds, which should result in a more severe hindrance of Cu diffusion to the top surface during selenization. This may induce an overly Cu-deficient top film. Conversely, the final top CISe film selenized from a precursor layer containing less amount of carbon should demonstrate an increased Cu content. In this regard, a “high carbon” and “low carbon” films were intentionally prepared by adding twice the amount of MEA to our standard recipe and then air-drying at lower temperature of 200 °C and by air-drying the film deposited using the standard ink at higher temperature for a longer time, respectively. It should be noted that “high” and “low” carbon content means the relative amount of carbon in the precursor film, which can be understood as a relative numbers of Cu-MEA bonding. The selenization conditions were same for both samples. Cross-sectional SEM images and a depth profile of the Cu and In contents measured by AES (normalized to (Cu+In)=100 % for clarity) of the two films are presented in Figure 10. The excess amount of carbon in the “high carbon” film again resulted in a double-layered structure, as shown in Figure 10(a). As expected, most of Cu was captured by the carbon in the bottom CRL, while the top carbon-free film is extremely In rich, resulting in an almost complete separation of the two layers of a carbon-free In-selenide top film on an Cu-selenide embedded bottom CRL. In contrast, in the “low carbon” film, due to strong air-drying of the precursor film, all of the carbon was removed (Figure 10(c)). The resultant film shows a uniform distribution of Cu and In throughout the entire film (Figure 10(d)), which can be explained by the lack of


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carbon to hinder the diffusion of Cu, further supporting the validity of our postulation. Note that, apart from the compositional distribution, the morphology of the carbon-free films is greatly affected by carbon content in the film. In general, carbon provides flatness to our CISe/CRL films, as evident from SEM images of Figure 8(d) and (e), and Figure 10(a), while a CISe film grown under carbon-free condition exhibits a very rough and porous morphology (Figure 10(c)). Possible explanations for this phenomenon are as follows. i) Suppression of Cu-Se liquid phase formation at the initial stage of the top film growth in the case that the precursor film contains carbon. In the carbon-free precursor film, the mobile Cu not captured by carbon is expected to easily diffuse to the top surface to form Cu-Se liquid, which is the first selenide phase formed in this case. This Cu-Se liquid might induce a long-range and fast diffusion, leading to a facile but excessively rapid growth of the grains, according to Kim et al.38 and Ahn et al.39 They also claimed that the rapid grain growth was the main reason for the pit and pore formation or the abnormal growth of the particles.38, 39 According to this explanation, carbon in the precursor film can suppress the initial formation of Cu-Se and hence mitigate the rapid grain growth, ensuring a much flatter film morphology. ii) Surface energy homogenization by carbon. As explained above, the very initial phenomenon of the top film growth is the formation of metal-selenide nuclei on the precursor film. This nucleation and initial film formation is expected to occur at the site where the surface energy is high, which is expected to be influenced by the existence of carbon at that surface region. When amorphous or very low-crystalline metal-oxides, whose surface energy is expected to significantly vary from location to location, are exposed at the surface of carbon-free precursor film, the initial metal-selenide nuclei formation might occur only at local high energy site, thus causing the film growth to proceed in a quite random and non-uniform manner. However, if the entire surface is covered by carbon, then the bonding of carbon with


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metal ions may passivate the high surface energy sites and enable the whole system to have a uniform surface energy. In this case, nucleation of selenide crystals can occur more uniformly, resulting in a flat and dense film. To clearly understand this carbon-assisted flat and uniform film growth, more studies are needed.

Figure 10. (a) Cross-sectional SEM image and (b) Cu and In depth composition profile of “high carbon” CISe film. (c) Cross-sectional SEM image and (d) Cu and In depth composition profile of “low carbon” CISe film. The atomic concentration is normalized to (Cu+In)=100 %. In addition to the main elements of Cu and In, another important element to be considered is Na. To date, there is a lack of information regarding how the diffusion of Na is influenced by a relatively thick CRL, which is approximately 600 nm thick in our study. To address this issue,


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we performed a SIMS analysis for a CISe/CRL film selenized for 30 min at 560 °C (point (e) in Figures 7). The obtained SIMS depth profile of Na is presented in Figure 11. For comparison, the Na SIMS signal obtained from a CISe film fabricated using a 3-stage vacuum co-evaporation technique on the same Mo/SLG substrate is also shown in Figure 11. First, the Na intensity in the Mo part is similar in both the solution-processed and vacuum-processed films, confirming that the Na contents of the substrates are almost identical. Comparing the Na intensity in the carbonfree parts of the two samples (top CISe layer in the solution-processed film and whole layer in the vacuum-processed film), the Na level in the top CISe layer in the solution-processed film is even slightly higher than that in the vacuum-processed film. More interestingly, we found unexpectedly high levels of Na in the CRL of the solution-processed sample. This result indicates that the CRL in our film does not act as a barrier for Na diffusion to the top CISe at all. Rather, the CRL acts as another high Na reservoir just beneath the CISe layer and allows for facile penetration of Na towards the top. It is difficult to address the detailed electrical impact of this Na distribution in our film at this stage; however, it is clear at least that solution-processed CISe films have a Na concentration comparable to that in the co-evaporated films.


25


solution CISe/CRL vacuum CISe 10

intensity / counts

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7

CRL

CISe

10

6

10

5

(MoSe2)/Mo

CISe

10

4

0

500

1000

1500

2000

sputtering time / s

Figure 11. Na SIMS depth profiles for a solution-processed CISe/CRL film and a vacuumprocessed CISe film. A three-stage co-evaporation technique was used to form the vacuumprocessed film. 2.4. Influence of carbon on the device characteristics In the previous sections, carbon was found to play a significant role in determining the depth elemental distribution of Cu and In by affecting their diffusion behavior during selenization. Although CRL itself is electrically benign to the device performance, the non-uniform depth distribution of metals should significantly affect the characteristics of the final devices as the composition of the top CISe film will be deviated from that original value kept during the preparation of the molecular ink. This motivated us to investigate the effects of the carbon content in the precursor films on the characteristics of the final device. In this regard, three devices were fabricated with the CISe/(CRL) layers selenized from the precursor films with different amounts of carbon content,


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classified as; “intermediate carbon” (Figure 1(b)), “high carbon” (Figure 10(a)) and “low carbon” (Figure 10(c)), according to conventional Mo/absorber/CdS/i-ZnO/n-ZnO/Al structure. (Here, the “intermediate carbon” device was built with the film shown in Fig. 1 using our standard ink while “high carbon” and “low carbon” devices were built with the films shown in Fig. 10(a) and (b), respectively). To check the reproducibility and validity of our finding, more than 10 samples of each carbon content were prepared and the standard deviation was calculated. Almost all of the samples of same amount carbon showed similar behavior with their representative dark I-V curves of shown in Figure 12. The “intermediate carbon” device, prepared using our standard ink, exhibited reasonably good rectifying behavior, while the “low carbon” and “high carbon” devices exhibited large leakage currents and complete short-circuit characteristics, respectively. For the “high carbon” case, due to the significantly limited diffusion of Cu, the top layer is almost In-selenide (Figure 10(b)), which is typically considered as an ntype semiconductor, and thus a p-n junction is not formed. This lack of a p-n junction is the reason for the perfectly linear I-V characteristic, with no detectable photovoltage. This behavior was common for all the samples of the “high carbon” device. The “low carbon” device is expected to have a CISe/CdS junction for a uniform Cu and In distribution (Figure 10(d)). However, a very rough and porous film may suffer from imperfect coverage of CdS and/or iZnO on the CISe film, resulting in a significant amount of shunting, as can also be seen in Figure 12. The “low carbon” devices showed a mean efficiency of 1.5 % ± 0.4 when illuminated. In contrast to “high carbon” and “low carbon” devices, the “intermediate carbon” devices demonstrated high efficiencies of up to 9.15% with average efficiency of 8.4 % ± 0.3.


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100

high carbon intermediate carbon low carbon

50 0

-50

-100

-1

0

1

voltage / V Figure 12. Dark I-V curves of “high carbon”, “low carbon” and “intermediate carbon” CISe devices. The light illuminated I-V and EQE curves of the best “intermediate carbon” device are presented in Figures 13(a) and 13(b), respectively. Note that this high efficiency, which is comparable to the efficiencies reported from the carbon-free, solution-processed CIGSSe devices,16-18 is achieved with a relatively thick CRL (as shown in Figure 1(b)) even without Ga addition, re-confirming that the CRL itself is not a critically detrimental factor in our device. In addition, it should be noted that this “intermediate carbon” film still suffers from an unoptimized composition of the top CISe layer, which is presumably excessively Cu-deficient compared to the typical optimum Cu/In ratio of approximately 0.8~0.9 due to the captured Cu in the bottom CRL. Therefore, further optimization of the top CISe composition, even with the CRL, can improve the device performance to a level higher than 9.15 %. Finally, we suggest that an appropriate direction for future research in solution-processed CIGSe film fabrication should not focus only on the removal of the CRL itself, rather, such studies should pay more attention to better compositional optimization of the top CIGSe layer.


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Active area : 0.44 cm without AR coating

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60 40 20 0

400

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wavelength / nm

Figure 13. (a) Light illuminated I-V curve and (b) EQE curve of the best “intermediate carbon” CISe device

3. Conclusions

We performed a systematic investigation to elucidate the role of carbon in the solution-processed CISe thin film solar cells. The CRL itself was determined to be a good electrical conductor that forms an ohmic junction with the adjacent CISe and (MoSe2)/Mo layers. This result implies that CRL is electrically benign to device performance, in contrast to the generally accepted belief that the CRL is the cause of high series resistance. However, carbon was found to play a significant


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role in determining the depth elemental distribution of the final selenized film, in which a higher carbon content in the precursor films causes the top CISe films to become more Cu-deficient and morphologically more compact and flatter while the CRL more Cu-rich. When a new film formed on the surface of the underlying precursor layer by reaction between Se vapor and the metal ions diffused to the surface, carbon selectively hindered the diffusion of Cu due to the stronger bonding between Cu ions and organic ligands than the bonding between In ions and organic ligands. This non-uniform diffusion of metal ions resulted in the non-uniform depth composition of the final film, i.e., Cu-deficient top CISe layer, with the CRL enriched with Cu. As a consequence, this deviation of composition of the top CISe layer from the optimal composition, designed when the precursor ink was prepared, is found to significantly affect the device performance. As a proof of this claim, the effects of carbon content in the precursor film on the device performance were investigated. CISe solar cells with an intermediate amount of carbon demonstrated high efficiencies of up to 9.15 %, while the devices built with highly carbon-rich and carbon-free CISe films exhibited very low efficiency. These results could be successfully explained by the carbon-affected different diffusion behavior of the metal ions together with the impact of carbon on the film morphology. Finally, we suggest that future research studies on the solution-processed CIGSe thin film formation should focus mainly on better compositional optimization of top CIGSe layer instead of only on the removal of the CRL. Improvement of properties of top CIGSe films and hence device performance can be pursued regardless of presence of CRL.

4. Experimental Section


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4.1. Materials: Copper (II) acetate monohydrate (99.99 %), indium (III) acetate (99.99 %), and MEA (99.0 %) were purchased from Sigma Aldrich, 2-Methoxyethanol (99.8 %) from Junsei, and elemental Se (99.999 %, 3 mm shot) from Cerac.

4.2. Formation of precursor ink: Precursor ink was prepared by dissolving Cu(II) acetate monohydrate (3.6 mmol), In(III) acetate (4.2 mmol) into 15 mL of 2-methoxyethanol with 4 mL of MEA, and then stirring for 30 min, resulting in a deep-blue-colored solution. The Cu/In atomic ratio of the ink was 0.86.

4.3. Film deposition and selenization: The precursor ink was spin-coated onto a Mo (1 µm) coated soda-lime glass (SLG) substrate and then heated in air on a hot plate at 270 °C for 5 min to evaporate solvents and to partially decompose the precursor. This process was repeated several times until the required thickness of the precursor film is achieved. The precursor film was then selenized in Se-rich environment using a heat-treatment chamber, in which a graphite container is placed inside. The chamber was equipped with a temperature and pressure controller. The precursor samples were placed inside the graphite container along with 0.2 g of Se pellets, followed by covering the container with a graphite lid. After loading the samples in this way, the chamber was initially evacuated to a base pressure of 10-6 Torr using a turbo molecular pump, and then the background pressure was regulated by the injection of nitrogen gas to be 40 Torr. In a typical process, selenization was performed at 560 °C for 30 min.

4.4. Device fabrication: Solar cells were fabricated according to the conventional Mo/CISe/CdS/i-ZnO/n-ZnO/Al structure. A 50-nm thick CdS buffer layer was deposited onto the


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CISe films via chemical bath deposition (CBD), and i-ZnO (50 nm)/Al-doped n-ZnO (500 nm) were deposited onto the CdS layer using radio-frequency (rf) magnetron sputtering. An Al grid of 500 nm in thickness was deposited as a current collector using thermal evaporation. The active area of the completed cell was defined to be 0.44 cm2.

4.5. Characterization:

The morphology and composition of the selenized films were

investigated using high-resolution scanning electron microscopy (HRSEM, XL30SFEG Phillips Co., Holland at 10 kV) and energy dispersive spectroscopy (EDS, EDAX Genesis apex, acceleration voltage: 30 kV, collection time: 100 s with standard-less method), respectively. To obtain structural information, Raman scattering spectra of the selenized films were taken in the quasiback-scattering geometry by using the 514.5 nm line of an Ar-ion laser as the excitation source. The scattered light was filtered with a holographic edge filter, dispersed by a Spex 0.55 m spectrometer, and detected with a liquid-nitrogen-cooled back-illuminated charge-coupleddevice (CCD) detector array. The depth compositional profile of the precursor films and the selenized films were obtained via Auger electron spectroscopy (AES, PerkinElmer, SAM 4300) and secondary ion mass spectrometry (SIMS, Cameca IMS-7f, CAMECA ADDR.) without a standard. In this regard, the depth profile data were considered only semi-quantitatively. The optical transmission spectra were obtained using a UV-VIS-NIR spectrophotometer (UV3101PC, SHIMADZU, Japan) with a spectral range of 400-2000 nm. The AC impedance measurements were performed on selective devices at various forward biases in the frequency range of 0.1 to 106 Hz at an AC amplitude of 10 mV root-mean-square. Dark and light illuminated J-V characteristics of devices were obtained using a class AAA solar simulator (WXS-155S-L2, WACOM, Japan) under one-sun illumination (AM 1.5 G, 1000 mW/cm2) at 25


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°C. The voltage was swept in both forward and reverse directions with a scan speed of 0.5 V/s, and only the results being unaffected by the sweep direction were selected. Before measurement the equipment was calibrated with a National Renewable Energy Laboratory-certified crystalline silicon reference cell (PVM-524). All the device characterizations were performed in a special measurement laboratory, in which a primary reference solar cell calibration technique was recently established.40

ASSOCIATED CONTENT Supporting Information XRD pattern of CISe film and Nyquist and Bode plots of vacuum-processed CISe device are added in the supporting information file. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author Email: [email protected] (Ara Cho); [email protected] (SeJin Ahn)

ACKNOWLEDGMENT This work was supported by the New & Renewable Energy Core Technology Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) funded by the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20143030011950). This work was also conducted partly under the framework of the Research and Development Program of the Korea Institute of Energy Research (KIER) (B6-2419) and “Development of 25% Efficiency Grade


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Tandem CIGS Thin Film Solar Cell Core Technology” of MSIP (Ministry of Science, ICT and Future Planning) and ISTK (Korea Research Council for Industrial Science and Technology) of Republic of Korea. Further, the authors thank Prof. J. Kim at Incheon National University for Raman analysis. ABBREVIATIONS CISe, Copper Indium Diselenide; MEA, Monoethanolamine; CRL, Carbon-containing residue layer; SLG, Soda Lime Glass; XRD, X-ray diffraction. REFERENCES [1]

Cho, A.; Ahn, S.; Yun, J.; Gwak, J.; Song, H.; Yoon, K. A Hybrid Ink of Binary Copper

Sulfide Nanoparticles and Indium Precursor Solution for a Dense CuInSe2 Absorber Thin Film and its Photovoltaic Performance. J. Mater. Chem. 2012, 22, 17893-17899. [2]

Li, M.; Zheng, M.; Zhou, T.; Li, C.; Ma, L.; Shen, W. Fabrication and Characterization of

Ordered CuIn(1−x)GaxSe2 Nanopore Films via Template-based Electrodeposition. Nanoscale Res. Lett. 2012, 7, 675-680. [3]

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Powalla, M. Properties of Cu(In,Ga)Se2 Solar Cells with new Record Efficiencies up to 21.7 %. Phys. Status Solidi RRL 2015, 9, 28-31. [4]

Hibberd, C. J.; Chassaing, E.; Liu, W.; Mitzi, D. B.; Lincot, D.; Tiwari, A. N. Non-

Vacuum Methods for Formation of Cu(In,Ga)(Se,S)2 Thin Film Photovoltaic Absorbers. Prog. Photovoltaics 2010, 18, 434-452.


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Nakashima, M.; Fujimoto, J.; Yamaguchi, T.; Izaki, M. Cu2SnS3 Thin-Film Solar Cells

Fabricated by Sulfurization from NaF/Cu/Sn Stacked Precursor. Appl. Phys. Express 2015, 8, 042303. [7]

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Szczepaniak, S. M.; Carter, N. J. C.; Handwerker, A.; Agrawal, R. A

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of Self-Stabilized Aqueous Nanoinks for Cu2ZnSn(S,Se)4 Solar Cells. Chem. Mater. 2014, 26, 3573-3578. [9]

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M.; Werner, M.; Haass, S.; Stuckelberger, J.; Broussillou, C.; Grand, P. P.; Bermudez, V.; Tiwari, A. N. All Solution-Processed Chalcogenide Solar Cells – from Single Functional Layers Towards a 13.8 % Efficient CIGS Device. Adv. Funct. Mater. 2015, 25, 12-27. [10] Tian, Q.; Wang, G.; Zhao, W.; Chen, Y.; Yang, Y.; Huang, L.; Pan, D. Versatile and Low-Toxic Solution Approach to Binary, Ternary, and Quaternary Metal Sulfide Thin Films and Its Application in Cu2ZnSn(S,Se)4 Solar Cells. Chem. Mater. 2014, 26, 3098-3103.


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[11] McLeod, S. M.; Hages, C. J.; Carter, N. J.; Agrawal, R. Synthesis and Characterization of 15 % Efficient CIGSSe Solar Cells from Nanoparticle Inks. Prog. Photovoltaics 2015, 23, 15501556. [12] Jeong, S. H.; Lee, B. S.; Ahn, S.; Yoon, K. H.; Seo, Y. H.; Choi, Y.; Ryu, B. H. An 8.2 % Efficient Solution-Processed CuInSe2 Solar Cell Based on Multiphase CuInSe2 Nanoparticles. Energy Environ. Sci. 2012, 5, 7539-7542. [13] Uhl, A. R.; Fuchs, P.; Rieger, A.; Pianezzi, F.; Sutter-Fella, C. M.; Kranz, L.; Keller, D.; Hagendorfer, H.; Romanyuk, Y. E.; LaMattina, F.; Yoon, S. H.; Karvonen, L.; Friedlmeier, T. M.; Ahlswede, E.; VanGenechten, D.; Stassin, F.; Tiwari, A. N. Liquid-Selenium-Enhanced Grain Growth of Nanoparticle Precursor Layers for CuInSe2 Solar Cell Absorbers. Prog. Photovoltaics 2014, DOI: 10.1002/pip.2529 [14] Ahn, S.; Choi, Y. J.; Kim, K.; Eo, Y.; Cho, A.; Gwak, J.; Yun, J.; Shin, K.; Ahn, S. K.; Yoon, K. Amorphous Cu–In–S Nanoparticles as Precursors for CuInSe2 Thin-Film Solar Cells with a High Efficiency. ChemSusChem 2013, 6, 1282-1287. [15] Uhl, A. R.; Fella, C.; Chirila, A.; Kaelin, M. R.; Karvonen, L.; Weidenkaff, A.; Borca, C. N.; Grolimund, D.; Romanyuk, Y. E.; Tiwari, A. N. Non-Vacuum Deposition of Cu(In,Ga)Se2 Absorber Layers from Binder Free, Alcohol Solutions. Prog. Photovoltaics 2012, 20, 526-533. [16] Septina, W.; Kurihara, M.; Ikeda, S.; Nakajima, Y.; Hirano, T.; Kawasaki, Y.; Harada, T.; Matsumura, M. Cu(In,Ga)(S,Se)2 Thin Film Solar Cell with 10.7% Conversion Efficiency Obtained by Selenization of the Na-Doped Spray-Pyrolyzed Sulfide Precursor Film. ACS Appl. Mater. Interfaces 2015, 7, 6472-6479.


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[17] Hossain, M. A.; Tianliang, Z.; Keat, L. K.; Xianglin, L.; Prabhakar, R. R.; Batabyal, S. K.; Mhaisalkarab, S. G.; Wong, L. H. Synthesis of Cu(In,Ga)(S,Se)2 Thin Films using an Aqueous Spray-Pyrolysis Approach, and their Solar Cell Efficiency of 10.5 %. J. Mater. Chem. A 2015, 3, 4147-4154. [18] Park, S. J.; Cho, Y.; Moon, S. H.; Kim, J. E.; Lee, D. K.; Gwak, J.; Kim, J.; Kim, D. W.; Min, B. K. A Comparative Study of Solution-Processed Low- and High-Band-Gap Chalcopyrite Thin-Film Solar Cells. J. Phys. D: Appl. Phys. 2014, 47, 135105-135111. [19] Batuk, M.; Buffiere, M.; Zaghi, A. E.; Lenaers, N.; Verbist, C.; Khelifi, S.; Vleugels, J.; Meuris, M.; Hadermann, J. Effect of the Burn-Out Step on the Microstructure of the SolutionProcessed Cu(In,Ga)Se2 Solar Cells. Thin Solid Films 2015, 583, 142-150. [20] Ahn, S.; Kim, C. W.; Yun, J.; Gwak, J.; Jeong, S.; Ryu, B. H.; Yoon, K. H. CuInSe2 (CIS) Thin Film Solar Cells by Direct Coating and Selenization of Solution Precursors. J. Phys. Chem. C 2010, 114, 8108-8113. [21] Azimi, H.; Hou, Y.; Brabec, C. J. Towards Low-Cost, Environmentally Friendly Printed Chalcopyrite and Kesterite Solar Cells. Energy Environ. Sci. 2014, 7, 1829-1849. [22] Park, S. J.; Cho, J. W.; Lee, J. K.; Shin, K.; Kim, J. H.; Min, B. K. Solution Processed High Band-gap CuInGaS2 Thin Film for Solar Cell Applications. Prog. Photovoltaics 2014, 22, 122-128. [23] Lee, E.; Park, S. J.; Cho, J. W.; Gwak, J.; Oh, M. K.; Min, B. K. Nearly Carbon-Free Printable CIGS Thin Films for Solar Cell Applications. Sol. Energy Mater. Sol. Cells 2011, 95, 2928-2932.


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[24] Park, S. N.; Sung, S. J.; Son, D. H.; Kim, D. H.; Gansukh, M.; Cheong, H.; Kang, J. K. Solution-Processed Cu2ZnSnS4 Absorbers Prepared by Appropriate Inclusion and Removal of Thiourea for Thin Film Solar Cells. RSC Adv. 2014, 4, 9118-9125. [25] Mainz, R.; Walker, B. C.; Schmidt, S. S.; Zander, O.; Weber, A.; Alvarez, H. R.; Just, J.; Klaus, M.; Agrawal, R.; Unold, T. Real-time Observation of Cu2ZnSn(S,Se)4 Solar Cell Absorber Layer Formation from Nanoparticle Precursors. Phys. Chem. Chem. Phys. 2013, 15, 18281-18289. [26] Lee, D.; Yong, K. Non-Vacuum Deposition of CIGS Absorber Films for Low-Cost Thin Film Solar Cells. Korean J. Chem. Eng. 2013, 30(7), 1347-1358. [27] Wu, W.; Cao, Y.; Caspar, J. V.; Guo, Q.; Johnson, L. K.; Malajovich, I.; Rosenfeld, H. D.; Choudhury, K. R. Studies of the Fine-grain Sub-layer in the Printed CZTSSe Photovoltaic Devices. J. Mater. Chem. C 2014, 2, 3777-3781. [28] Ahn, S.; Son, T. H.; Cho, A.; Gwak, J.; Yun, J.; Shin, K.; Ahn, S. K.; Park, S. H.; Yoon, K. H. CuInSe2 Thin-Film Solar Cells with 7.72% Efficiency Prepared via Direct Coating of a Metal Salts/Alcohol-Based Precursor Solution. ChemSusChem 2012, 5, 1773-1777. [29] Lee, D. Y.; Yun, J. H.; Yoon, K. H.; Ahn, B. T. Characterization of Cu-poor Surface on Cu-rich CuInSe2 Film Prepared by Evaporating Binary Selenide Compounds and its Effect on Solar Efficiency. Thin Solid Films 2002, 410, 171-176. [30] Shin, Y. M.; Lee, C. S.; Shin, D. H.; Kwon, H. S.; Park, B. G.; Ahn, B. T. Surface Modification of CIGS Film by Annealing and its Effect on the Band Structure and Photovoltaic Properties of CIGS Solar Cells. Curr. Appl. Phys. 2015, 15, 18-24.


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[31] Baydogan, N. D. Evaluation of Optical Properties of the Amorphous Carbon Film on Fused Silica. Mater. Sci. Eng., B 2004, 107, 70-77. [32] Bisquert, J. A Variable Series Resistance Mechanism to Explain the Negative Capacitance Observed in Impedance Spectroscopy Measurements of Nanostructured Solar Cells. Phys. Chem. Chem. Phys. 2011, 13, 4679-4685. [33] Kim, M. S.; Chalapathy, R. B. V.; Yoon, K. H.; Ahn, B. T. Grain Growth Enhancement and Ga Distribution of Cu(In0.7Ga0.3)Se2 Film Using Cu2Se Layer on Cu–In–Ga Metal Precursor. J. Electrochem. Soc. 2010, 157 (1), B154-B158. [34] Soltz, D.; Dagan, G.; Cahen, D. Ionic Mobility and Electronic Junction Movement in CulnSe2. Solid State Ionics 1988, 28-30, 1105-1110. [35] Becker, K. D.; Wagner, S. Temperature-Dependent Nuclear Magnetic Resonance in CuInX2 (X=S,Se,Te) Chalcopyrite-Structure Compounds. Phys. Rev. B 1983, 27, 5240-5249. [36] Tell, B.; Wagner, S.; Bridenbaugh, P. M. Motion of p-n Junctions in CulnSe2. Appl. Phys. Lett. 1976, 28, 454-455. [37] Cho, A.; Song, H.; Gwak, J.; Eo, Y.; Yun, J.; Yoon, K.; Ahn, S. A Chelating Effect in Hybrid Inks for Non-Vacuum Processed CuInSe2 Thin Films. J. Mater. Chem. A 2014, 2, 50875094. [38] Kim, K. H.; Yoon, K. H.; Yun, J.; Ahn, B. T. Effects of Se Flux on the Microstructure of Cu(In, Ga0.3)Se2 Thin Film Deposited by a Three-Stage Co-evaporation Process. Electrochem. Solid-State Lett. 2006, 9 (8), A382-A385.


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[39] Ahn, S.; Kim, K.; Cho, A.; Gwak, J.; Yun, J.; Shin, K.; Ahn, S. K.; Yoon, K. CuInSe2 (CIS) Thin Films Prepared from Amorphous Cu−In−Se Nanoparticle Precursors for Solar Cell Application. ACS Appl. Mater. Interfaces 2012, 4, 1530-1536. [40] Ahn, S. K.; Ahn, S.; Yun, J.; Lee, D. H.; Winter, S.; Igari, S.; Yoon, K. H. Establishment of a Primary Reference Solar Cell Calibration Technique in Korea: Methods, Results and Comparison with WPVS Qualified Laboratories. Metrologia 2014, 51, 139-147.


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Table of Contents


41


(b)

(a)

CISe CRL

MoSe2 Mo

Mo 2 µm

2 µm

6000

80

(c)

(d)

CISe2 A1

CISe

4000

CRL

Se

40

CISe2 B2/E

OVC

C

In

CISe2 B2/E

20

Mo

2000

Cu 100

60

150

200

250

300

0.0

-1

Raman shift / cm

atomic conc. / %

Raman intensity / cps

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 42 of 53

O 0.5

1.0

0

depth / m

Figure 1. Characteristics of a typical CISe film; (a) planar SEM image, (b) cross-sectional SEM image, (c) Raman scattering spectra and (d) AES depth profile.


2

3

CISe

(b)

CuInSe2 A1

(a)

CISe2 B2/E 4

5

CRL

6

CISe2 B2/E

1 1 7

8

MoSe2

4900

9

4200

2 2

Mo

5600

3500

OVC

2 µm

100

CuInSe2

4900

2800

100

200

MoSe2

250

300

350

(d)

MoSe2 MoSe2

7 7

4 4 4200

3500

150

(c)

Cu2Se

2800

3 3

4900

MoSe2 4200

MoSe2 8 8

55

66

3500

2800

99 150

200

250

300

350

Raman shift (cm-1)

100

150

200

250

300

Raman Intensity / cps

1


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 43 of 53

350

Raman shift (cm-1)

Figure 2. (a) Locations where the Raman spectra are taken in a CISe/CRL/Mo sample after inclined polishing; Raman spectra obtained from (b) top CISe layer, (c) CRL and (d) MoSe 2 layer.



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 44 of 53

(b)

(a)

CISe CRL

CRL 2 µm

2 µm

Figure 3. Cross-sectional SEM images of a selenized samples on SLG (a) before and (b) after etching of the top CISe layer.


(i) CRL (ii) CISe + CRL (iii) CISe only : calculated

60

40

(i) (iii) (ii)

600

900

1200

20

(transmittance %)

80

(a)

0

1500

wavelength / nm

(b) CISe CRL (iii) %TCISe only (i) %TCRL

(ii) %TCISe+CRL 8

(c)

6

4

2

8

from (iii) of Fig. 4(a)

-2

Eg of CISe film

2

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


 x 10 / cm

Page 45 of 53

Eg = 0.96 eV 0.5

1.0

1.5

0 2.0

h / eV

Figure 4. (a) Transmittance curves and (b) conceptual schematic drawing of (i) CRL, (ii) CISe+CRL and (iii) CISe only cases. Note that (i) and (ii) are experimental data and (iii) is calculated using (i) and (ii). (c) Bandgap of CISe film estimated from (iii) of Figure 4(a).



Ag/CRL/MoSe2/Mo device

150

(c)

-2

(a)

Ag

100 CRL

Ag/CRL/MoSe2/Mo completed CISe device

MoSe2/Mo

(b)

completed CISe device

50 0

Ag CdS/ZnO

-50

CISe CRL MoSe2/Mo

-1.0

-0.5

0.0

0.5


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 46 of 53

-100 1.0

voltage / V

Figure 5. Schematics of (a) Ag/CRL/MoSe2/Mo device and (b) completed CISe device; (c) I-V curves of the two devices.


Page 47 of 53

0.1 V 0.2 V 0.3 V 0.4 V 0.5 V 0.6 V

(a)

-Z" / Kcm

2

6

4

2

0

0

2

4

6

8

10

2

Z' / Kcm

(b)

10

3

10

2

10

1

|Z| / cm

2

10

4

10 -80

 / degree

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


0.4 V 0.5 V 0.6 V

0.1 V 0.2 V 0.3 V

-1

10

0

1

10

10

2

10

3

10

4

5

10

10

6

frequency / Hz

-60

0.1 V 0.2 V 0.3 V 0.4 V 0.5 V 0.6 V

-40

-20

0 -1 10

0

10

1

10

2

10

3

10

4

10

5

10

6

10

frequency / Hz

Figure 6. (a) Nyquist and (b) Bode plots of a solution-processed CISe device measured at various forward biases.



(e)

(d) 560 sub. temp. / oC

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 48 of 53

(c)

400

(b)

300

200

(a)

30min

time / min

Figure 7. Thermal history of selenization, showing the points (a)-(e), where the process was interrupted and the samples were removed for analysis.


Page 49 of 53

1

2

3 60

(a) 2 µm

2 µm

atomic conc. / %

50 40

In

30

O

Cu

20

C

10

Mo

Se

0 0.0

0.5

1.0

depth / m

60

(b) 2 µm

2 µm

atomic conc. / %

50 40 30

In Cu

O 20

C 10

Mo

Se

0 0.0

0.5

1.0

depth / m

60

CRL

(c) 2 µm

2 µm

atomic conc. / %

50

Se 40

In

30 20

Cu

C

O

10 0 0.0

Mo 0.5

depth / m

(d) 2 µm

2 µm

atomic conc. / %

60

CRL

Se

50 40

In

30 20

C

Cu

Mo

10

O 0 0.0

0.5

1.0

depth / m

60 50

(e) 2 µm

2 µm

atomic conc. / %

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


Se

CRL

40

In

C

30 20

Cu Mo

10

O 0 0.0

0.5

depth / m

1.0

Figure 8. Column 1, 2, and 3, showing planar SEM images, cross-sectional SEM images, and AES depth profiles of elemental diffusion in the precursor film at points designated as (a), (b), (c), (d), and (e), in Figure 7.



100

100

(b)

80

Cu

60

60

O

O

40

40

In

20

0

80

Mo

Mo

20

C 0

2

4

C 0

6

2

sputter time / min

4

6

8

0 12

10

sputter time / min

Figure 9. AES depth profiles of (a) Cu-precursor film and (b) In-precursor film.

C-free film

CRL

100

(a)

80

In

C-free film

60

CRL

40

Cu 20

2 µm 0

10

20

atomic conc. normalized to (Cu+In)=100%

(b)

0 30

sputter time / min C-free film

100

(d) 80

In

C-free film

60 40

Cu 20

2 µm 0

10

20

sputter time / min


0 30

atomin conc. normalized to (Cu+In)=100%

(c)

atomic conc. / %

(a) atomic conc. / %

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 50 of 53

Page 51 of 53

Figure 10. (a) Cross-sectional SEM image and (b) Cu and In depth composition profile of “high carbon” CISe film. (c) Cross-sectional SEM image and (d) Cu and In depth composition profile of “low carbon” CISe film. The atomic concentration is normalized to (Cu+In)=100 %.

solution CISe/CRL vacuum CISe 10

intensity / counts

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


7

CRL

CISe

10

6

10

5

(MoSe2)/Mo

CISe

10

4

0

500

1000

1500

2000

sputtering time / s

Figure 11. Na SIMS depth profiles for a solution-processed CISe/CRL film and a vacuumprocessed CISe film. A three-stage co-evaporation technique was used to form the vacuumprocessed film.



-2

150


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 52 of 53

100

high carbon intermediate carbon low carbon

50

0

-50

-100

-1

0

1

voltage / V Figure 12. Dark I-V curves of “high carbon”, “low carbon” and “intermediate carbon” CISe devices.



40

2

Active area : 0.44 cm without AR coating

(a)

30

20

10

VOC : 0.43 V JSC : 33.8 mAcm-2 FF : 0.63 Eff : 9.15 %

0 0.0

0.1

0.2

0.3

0.4

0.5

voltage / V 100

(b) 80

EQE / %

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


-2

Page 53 of 53

60 40 20 0

400

600

800

1000

1200

1400

wavelength / nm

Figure 13. (a) Light illuminated I-V curve and (b) EQE curve of the best “intermediate carbon” CISe device


Double Laser for Depth Measurement of Thin Films of Ice.

Electrodeposition of antimony selenide thin films and application in semiconductor sensitized solar cells.

Solvent Vapor Treatment Effects on Poly(3-hexylthiophene) Thin Films and its Application for Interpenetrating Heterojunction Organic Solar Cells.

In-depth analysis of chloride treatments for thin-film CdTe solar cells.

Dye-sensitized solar cell using sprayed ZnO nanocrystalline thin films on ITO as photoanode.

Correlating the Polymorphism of Titanyl Phthalocyanine Thin Films with Solar Cell Performance.

Erratum: Phase transitions via selective elemental vacancy engineering in complex oxide thin films.

Potassium-induced surface modification of Cu(In,Ga)Se2 thin films for high-efficiency solar cells.

Phase transitions via selective elemental vacancy engineering in complex oxide thin films.

High-performance perovskite CH3NH3PbI3 thin films for solar cells prepared by single-source physical vapour deposition.

Atomically thin group v elemental films: theoretical investigations of antimonene allotropes.

Enhanced optoelectronic quality of perovskite thin films with hypophosphorous acid for planar heterojunction solar cells.

In situ intercalation dynamics in inorganic-organic layered perovskite thin films.

Heterogeneous Charge Carrier Dynamics in Organic-Inorganic Hybrid Materials: Nanoscale Lateral and Depth-Dependent Variation of Recombination Rates in Methylammonium Lead Halide Perovskite Thin Films.

Nano-structured CuO-Cu2O Complex Thin Film for Application in CH3NH3PbI3 Perovskite Solar Cells.

Thermal dewetting behavior of polystyrene composite thin films with organic-modified inorganic nanoparticles.

Cell death in the thin films of bursting bubbles.

molecular layer-deposited hybrid organic-inorganic zincone thin films.

Local control over nucleation of epitaxial thin films by seed layers of inorganic nanosheets.

Aerosol-assisted atmospheric cold plasma deposition and characterization of superhydrophobic organic-inorganic nanocomposite thin films.

Controlling the surface nanostructure of ZnO and Al-doped ZnO thin films using electrostatic spraying for their application in 12% efficient perovskite solar cells.

Perovskite solar cells based on nanocolumnar plasma-deposited ZnO thin films.

Magnetoelectric Fe2TeO6 thin films.

Thin hydrogel films for optical biosensor applications.