The path towards a high-performance solution ...

Solar Energy Materials & Solar Cells ] (]]]]) ]]]–]]]

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The path towards a high-performance solution-processed kesterite solar cell$ David B. Mitzi n, Oki Gunawan, Teodor K. Todorov, Kejia Wang, Supratik Guha IBM T.J. Watson Research Center, P.O. Box 218, Yorktown Heights, NY 10598, USA

a r t i c l e i n f o

Keywords: Cu2ZnSn(S,Se)4 CZTS Kesterite Earth-abundant Thin-film solar cell Solution processing

a b s t r a c t Despite the promise of thin-film Cu(In,Ga)(S,Se)2 (CIGSSe) chalcopyrite and CdTe photovoltaic technologies with respect to reducing cost per watt of solar energy conversion, these approaches rely on elements that are either costly and/or rare in the earth’s crust (e.g., In, Ga, Te) or that present toxicity issues (e.g., Cd), thereby potentially limiting these technologies in terms of future cost reduction and production growth. In order to develop a photovoltaic technology that is truly compatible with terawatt deployment, it is desirable to consider material systems that employ less toxic and lower cost elements, while maintaining the advantages of the chalcopyrite and CdTe materials with respect to appropriate direct band gap tunability over the solar spectrum, high device performance (e.g., 4 10% power conversion efficiency) and compatibility with low-cost manufacturing. In this review, the development of kesterite-based Cu2ZnSn(S,Se)4 (CZTSSe) thin-film solar cells, in which the indium and gallium from CIGSSe are replaced by the readily available elements zinc and tin, will be reviewed. While vacuum-deposited devices have enabled optimization within the compositional phase space and yielded selenium-free CZTS device efficiencies of as high as 6.8%, more recently a liquid-based approach has been described that has enabled deposition of CZTSSe devices with power conversion efficiency of 9.7%, bringing the kesterite-based technology into a range of potential commercial interest. Electrical characterization studies on these high-performance CZTSSe cells reveal some of the key loss mechanisms (e.g., dominant interface recombination, high series resistance and low minority carrier lifetime) that limit the cell performance. Further elucidation of these mechanisms, as well as building an understanding of long-term device stability, are required to help propel this relatively new technology forward. & 2010 Elsevier B.V. All rights reserved.

1. Introduction Photovoltaic (PV) electricity generation currently meets less than 0.1% of world-wide electricity demand [1], despite orders of magnitude higher potential, primarily because of the cost discrepancy between solar and more conventional carbon-based technologies. Within the PV arena, over 80% of the solar market is currently dominated by silicon-based technology [1,2], in part because of an entrenched and relatively mature silicon industry, as well as the attractive device stability and power conversion efficiencies offered by crystalline silicon (record efficiency for crystalline silicon cells is currently 25.0% [3]). However, crystalline silicon technology relies on an absorber with indirect band gap, thereby necessitating a thick layer to absorb an appreciable

$ This manuscript is for the Special SOLMAT issue: ‘Thin film and nano-structured solar cells’ in honor of Professor Marta Ch. Lux-Steiner on the Occasion of Her 60th Birthday. Prof. Dr. Ahmed Ennaoui Physical Chemistry and thin film solar cells, Group leader Helmholtz-Center Berlin for Materials and Energy Institute Heterogeneous Material Systems Glienicker Strasse 100, 14109 Berlin, Germany. Tel.: ( +49) 030 8062 3038; fax: (+ 49) 030 8062 3199, e-mail: [email protected], Web: http://www.helmholtz-berlin.de and Greg P. Smestad Associate Editor Solar Energy Materials and Solar Cells www.solideas.com. n Corresponding author. Tel.: + 914 945 4176; fax: +914 945 2141. E-mail address: [email protected] (D.B. Mitzi).

fraction of the incident solar radiation ( 4100 mm assuming no light trapping incorporated within the cell design). Additionally, since grain boundaries are active as recombination centers in Sibased technologies, nominally perfect single crystal substrates are required to build highest efficiency solar modules, thereby contributing to higher cost. By contrast, thin-film PV technologies rely on direct band gap materials such as CdTe, CuIn(S,Se)2 (CISSe) and more generally Cu(In,Ga)(S,Se)2 (CIGSSe). Because of the direct band gap (and corresponding high absorption coefficient of 104–105 cm 1), material utilization can be reduced, with 1–2 mm layer thickness generally being enough to absorb most of the incident solar radiation. Reduced requirements for film crystalline quality (grain boundaries are less active as recombination centers [4]) also enable lower-cost routes to be employed for layer deposition. Both CdTe and CIGSSe technologies yield champion cell efficiencies of greater than 15% (16.7% for CdTe and 20.1% for CIGSSe—not far behind crystalline Si technology [3]) and module production with average module power conversion efficiency of greater than 10% [3]. For CIGSSe, module efficiencies of as high as 16% have been demonstrated on the laboratory scale [5]. Manufacturing production growth rates for both technologies are also impressive, with First Solar (CdTe) being the first manufacturer to achieve greater than 1 GW/yr production and also the first to achieve o$1/W production (at $0.76/W level as of August 2010) [6].

0927-0248/$ - see front matter & 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.solmat.2010.11.028

Please cite this article as: D.B. Mitzi, et al., Sol. Energy Mater. Sol. Cells (2011), doi:10.1016/j.solmat.2010.11.028

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Industrial production of CIGSSe-based modules is also expected to reach the GW level by 2011 [5]. Despite the promise of CdTe and CIGSSe, both technologies rely on elements that are scarce in the earth’s crust [7,8]. The abundance of indium in the upper continental crust is estimated to be 0.05 ppm (compared with an abundance of 25, 71 and 5.5 ppm for copper, zinc and tin, respectively) [9]. Current world-wide production capacity for indium is of order 600 metric tons per year [10]. However, much of this capacity is required for transparent conductive coatings for the growing flat panel display industry. One estimate, based on current production trends, is that world-wide production of indium can support a CIGSSe production capacity of approximately 70 GW/yr, well below the desired TW level [7]. The abundance of tellurium is even lower (0.001 ppm) [9]. Although there is a great deal of uncertainty in the ultimate production capacity limitations for CIGSSe and CdTe manufacturing [7,8,11], depending upon whether recycling programs are instituted for In and Te, the ultimate thickness of the absorber layer, the achievable module efficiencies and on competition for raw materials with other technologies, if pervasive deployment were implemented for CIGSSe and CdTe photovoltaics, materials costs would likely impact competition with other technologies as production volumes increase. Besides the issue of abundance, the heavy metal cadmium has experienced resistance towards adoption in some countries because of the toxicity issue. One reason for this resistance is the mass cadmium poisoning incident in the early-to mid-20th century, in which cadmium-laden waste water from mining operations in the Toyama Prefecture in Japan resulted in an outbreak of a painful and sometimes fatal affliction commonly referred to as ‘‘itai–itai’’ (translated literally as ‘‘ouch–ouch’’) disease [12,13]. Given the above considerations, there is a need to identify thin-film materials that are composed of plentiful and less toxic elements, while still providing adequate device performance, in order to enable pervasive deployment of PV technology to meet ever growing energy needs. While materials availability and toxicity are two considerations in the search for appropriate absorber layer compositions that might contribute, along with CdTe and CIGSSe, to the thin-film PV market, the material must fulfill several additional fundamental criteria in order to enable the two key targets for an efficient solar cell of: (1) effective absorption of incident photons to generate electron–hole pairs and (2) ability to collect the photo-generated charges before they recombine. Optimal absorption of the solar spectrum and generation of electron–hole pairs requires the appropriate choice of band gap. Theoretical considerations predict maximum power conversion efficiency in the 20–30% range for absorber materials with a band gap of between 1 and 2 eV (optimal overlap with the solar spectrum should occur for a band gap of 1.5 eV). A direct lowest energy band gap also provides a large absorption coefficient (a 4104 cm 1) for photons with energy greater than the band gap. A direct band gap thereby allows effective absorption in a chalcogenide layer thinner than a few microns, thus reducing significantly the materials needs, in comparison with indirect band gap materials such as crystalline silicon. The ability to rely on a thin absorber layer also relaxes to some extent the constraint on the diffusion length (essentially the average distance photo-generated carriers diffuse before recombination), which poses in silicon technology costly requirements for materials purity and crystal quality. Nevertheless, in addition to an appropriate direct band gap semiconductor, the desire for a material with a relatively large photo-generated carrier lifetime (or long diffusion length) is a second target in identifying appropriate absorber materials for highest efficiency devices. The need to concurrently address the requirements of appropriate size direct band gap and carrier lifetime, along with the need for a conveniently deposited material that has low toxicity and

readily available elements, has led to a recent shift of research focus towards the kesterite-related family of thin-film chalcogenide materials.

2. History of kesterite-related PV materials 2.1. Kesterite-related compounds and structures Chalcogenide compounds with a Cu2(MII)(MIV)(S,Se)4 (MII ¼ Mn, Fe, Co, Ni, Zn, Cd, Hg; MIV ¼Si, Ge, Sn) stoichiometry have been of interest for many years because of their appearance as naturally occurring minerals and also suitable direct band gaps for application in solar cells and other optical devices [14–19]. The zincblende (or sphalerite)-related structures adopted by these compounds depend on the degree and type of metal cation ordering within the face-centered cubic (fcc) array of chalcogenide anions (with both metals and chalcogens adopting a tetrahedral coordination). One way of viewing these structures is by starting from the ternary chalcopyrite structure (Fig. 1), CuMIII(S,Se)2, and replacing the trivalent MIII ions with an equal number of divalent MII and tetravalent MIV metals. When the ordering of the metals is such that Cu and MIV atoms alternate on the z¼0 and ½ (z ¼fractional coordinate along the long c-axis of the structure) planes and Cu and MII atoms alternate on the z¼¼ and 3/4 planes, this is known as the kesterite structure, whereas when MII and MIV atoms alternate on the z ¼0 and ½ planes and only Cu resides on the z ¼¼ and 3/4 planes, this is known as the stannite structure. While the kesterite structure has the same basic Cu/S structure as chalcopyrite, the stannite structure requires some reorganization of the Cu sublattice. Note that determining whether a structure adopts a kesterite or stannite structure is difficult to do without careful single crystal structural analysis or Rietveld-type analysis using X-ray and/or neutron diffraction data, which has not been completed for most of

Fig. 1. Schematic representation of the chalcopyrite structure (drawn with MIII ¼In), and kesterite and stannite structures (drawn with MII ¼Zn, MIV ¼ Sn). The unit cell boundaries are denoted with dashed lines and the space group for each structural type is also provided.



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Table 1 Lattice constants and band gaps for selected families of kesterite/stannite materials for which both structural and band gap information have been reported. Compound

˚ a (A)

˚ c (A)

V (A˚ 3)

Cu2ZnSnS4

5.427 5.432 5.435 5.426

10.848 10.840 10.822 10.81

319.5 319.9 319.7 318.3

5.567 5.668 5.681 5.684 5.688

11.168 11.349 11.34 11.353 11.338

346.1 364.6 366.0 366.8 366.8

5.341a 5.344 5.622 5.606 5.618 5.586 5.593

10.509a 10.513 11.06 11.042 11.04 10.834 10.840

299.8 300.2 349.6 347.0 348.4 338.1 339.1

Cu2ZnSn(Se2.4S1.6) Cu2ZnSn(Se3.9S0.1) Cu2ZnSnSe4

Eg (eV)

1.45 1.45 1.51 1.21 1.03 0.94 1.017

Cu2ZnGeS4 Cu2ZnGeSe4

Cu2CdSnS4

Cu2CdSnSe4

5.832

11.389

387.4

Cu2CdGeSe4

5.829 5.832 5.814 5.657

11.418 11.392 11.47 10.988

388.0 387.5 387.7 351.6

1.63 1.52

1.06 1.16 0.96 0.89

1.29

References [19] [26] [27] [28] [29] [30] [26] [26] [14] [31] [23] [32] [33] [27] [14] [16] [34] [19] [27] [29] [35] [16] [36] [23] [27] [14] [19] [36]

a ˚ Also reported [19] as orthorhombic with lattice constants a¼ 7.504 A, ˚ c¼ 6.185 A. ˚ b¼ 6.474 A,

the compounds in the above series. Two counter-examples are Cu2FeSn(S,Se)4, which has been determined to be stannite [15], and Cu2ZnSnS4 (CZTS), which was originally described as stannite, but is now recognized as kesterite [15,20,21]. Cu2ZnSnSe4 (CZTSe) has also been described as stannite in the literature [16,22–24]. However, the data presented for this compound are largely powder X-ray diffraction (XRD) data, which cannot reliably be used to distinguish between the two possibilities. Theoretical calculations predict that the lowest energy configuration for CZTSe (as for CZTS) is kesterite [22,25], although the difference in energy between kesterite and stannite is much smaller for the selenide vs. the sulfide, suggesting a stronger tendency of the selenide to incorporate stacking faults involving alternative layering sequences. We will describe all compounds in the Cu2ZnSn(S,Se)4 (CZTSSe) family as kesterite, with the recognition that experimental verification of this assignment for the high Se-content samples must still be provided. Lattice constants and band gaps of several more thoroughly studied kesterites and stannites are given in Table 1. 2.2. Kesterite-related photovoltaic films Table 1 demonstrates that many of the kesterite-related phases exhibit a direct band gap in the optimal range for photovoltaic energy conversion (all are reported to be p-type semiconductors). Despite this fact, relatively few of these compounds have been reported in actual solar cell devices. In 1988, Ito et al. [29] examined the electrical and optical properties of CZTS, a material which has all earth-abundant and relatively non-toxic elements, as well as Cu2CdSnS4 (CCTS). Both compounds were reported to be direct band gap p-type semiconductors with band gaps of 1.45 and 1.06 eV for the zinc and cadmium compounds, respectively. The first CZTS solar cell was constructed from a heterostructure with cadmium tin oxide, yielding an open circuit voltage of 165 mV under AM1.5 illumination (no efficiency reported). Earlier work on

Fig. 2. Record CZTSSe thin-film device performance vs. year, showing consistent progress in power conversion efficiency. Circles represent data for pure-sulfurcontaining materials (CZTS), whereas the triangle represents data for kesterites in which Se has been introduced. The laboratories responsible for the record are noted. References for the data points are given in the text.

the CCTS system [35], employing a heterodiode of a CCTS single crystal with a CdS film, yielded a power conversion efficiency of 1.6% (reported to be limited by a large series resistance), an open circuit voltage of 500 mV and a short circuit current density of 7.9 mA/cm2. By far the most extensive subsequent work on kesterite-related PV devices has been performed on CZTSSe materials, which are the primary focus of this review, leading to continual device performance improvement (Fig. 2) [37]. In 1996, Katagiri et al. [28,37] reported on sequentially evaporated CTZS (pure sulfide) films, yielding a glass/Mo/CZTS/CdS/ZnO:Al device efficiency of 0.66% and open circuit voltage of 400 mV. Within a year, the Stuttgart University group had pushed the power conversion efficiency of a similarly prepared device to 2.3% [38]. Currently, the record power conversion efficiency for vacuum-deposited CZTS has improved to 6.8% [37]. For analogous selenide CZTSe devices, in 1997 Friedlmeier et al. [38] reported on vacuum-fabricated films, obtaining a device efficiency of 0.6%. By 2009 the efficiency for CZTSe devices had increased to 3.2% [31]. The current CZTSSe kesterite record of 9.7% (Fig. 2) was achieved using a hydrazinebased solution deposition approach with mixed sulfur/selenium anions [26]. More detailed discussion of the various deposition methods for CZTSSe films, as well as approaches for improving efficiency, will be given in the next section, while analogous discussion on in-depth device characterization will be given in Section 4.

3. Approaches for CZTSSe film deposition The growing perception of CZTSSe as a potentially ideal photovoltaic absorber material is reflected in the abundance and variety of routes that have been undertaken for thin-film deposition. These include a wide spectrum of vacuum and atmospheric-pressure approaches. For each of these deposition approaches one key barrier toward a reliable and low-cost process is the complex and incompletely understood nature of the multinary Cu–Zn–Sn– S–Se phase diagram and control over the phase progression during film formation [39–41], which presents a challenge for preparing single phase films. A second common theme generally encountered


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during process optimization involves the volatility upon heating of Sn (and perhaps Zn) materials [42], which makes compositional control a challenge during film fabrication. Despite these challenges, reasonably successful film deposition and device fabrication has been demonstrated for the full range of S:Se ratios in CZTSSe and for both vacuum and solution-based deposition approaches. 3.1. Vacuum-based approaches The potentially high uniformity of vacuum deposition methods, coupled with their capacity to deliver precisely timed elemental fluxes, provides good opportunity for fabrication of high-quality thin-film devices. However, the transfer of these approaches to CZTSSe deposition has encountered significant challenges. Of particular concern is the issue of element volatility and consequent reevaporation from deposited films, which is further enhanced relative to atmospheric pressure approaches because of the vacuum conditions during heating. One study, for example, has found that the evolution of Sn from the film, which is believed to be in the form of SnS, decreases for compounds in the order SnSCu2SnS3-Cu4SnS4-Cu2ZnSnS4 [42]. Vacuum-based kesterite deposition can be roughly sub-divided into two categories, sputtering- and evaporation-based approaches. 3.1.1. Sputtering Ito and Nakazawa’s first report [29] on CZTS thin film deposition in 1988 was based on argon beam sputtering from pressed targets of the quaternary materials. Remarkably, no post-deposition sulfurization treatment was required using this approach and crystalline material was reported even for substrate temperatures during deposition of as low as 90 1C. No efficiency data of these devices was provided, although an open circuit voltage of 165 mV was demonstrated. More recent work has focused on radio frequency (RF) magnetron sputtering of sequentially deposited metal layers, which are subsequently treated in a sulfur-containing atmosphere at elevated temperature to yield the desired phase [30,43]. In one example, stacked Cu/Sn/Zn (Zn on bottom of stack) metal layers were sputtered and sulfurized using a flux of elemental sulfur in a flowing nitrogen atmosphere for 1 h at 300 1C and 3 h at 550 1C [30]. The initial low-temperature pretreatment was performed to minimize metal reevaporation at the subsequently applied elevated temperatures. In an analogous study, RF-sputtered metal multilayers were sulfurized with elemental sulfur in a two-zone (one zone for the sulfur source and one zone for the multilayer metal film, with the zone for the film heated to 570 1C within 20 min) vacuum furnace, with no carrier gas used for the S-vapor [44,45]. In each of these cases, metal compositional control for the final CZTS film is provided by varying the thickness of the individual metal precursor layers. The most successful sputtering route to CZTS has been developed by the group of Prof. Katagiri, after extensive pioneering work on the same material employing evaporation-based approaches (see Section 3.1.2). Using RF co-sputtering of Cu and binary ZnS and SnS materials, a 5.74% efficient solar cell was demonstrated in 2007 [46]. Fine intermixing and uniformity of the precursors was achieved by substrate rotation during the sputtering process and the samples were subjected to a high-temperature (580 1C) final anneal for 3 h using 20% H2S in nitrogen. Detailed investigation of different elemental ratios for CZTS and their effect on device performance resulted in highly useful composition maps for efficient kesterite absorbers, refining the earlier Cu-poor, Zn-rich prescriptions and demonstrating that high power conversion efficiency is achieved in a relatively narrow range of the compositional phase diagram ([Zn]/[Sn]E1.25, [Cu]/([Zn] +[Sn]) E0.9) [47]

(Fig. 3). In addition, in this study, the H2S concentration was treated as a variable during the sulfurization treatment, ranging from 5% to 20%, and little variation in film/device properties was noted, thereby suggesting that the lower H2S content processing should be more desirable in order to limit wear on the reactor chamber. Another important refinement in the Katagiri process was the observation that soaking the absorber in DI water (10 min at room temperature) prior to CdS deposition leads to significant device performance improvement [48]. The described preferential etching of surface oxide contamination (Fig. 4) is believed to lead to improvement in the FF and Jsc by decreased series resistance. These developments have led to the current record-performance sputtered kesterite device (Fig. 5), reaching 6.77% efficiency (0.15 cm2 active area), which also represents one of the highest efficiency values reported so far for a pure sulfide CZTS film prepared by any method. Analogous CZTSe films were prepared by selenization of a DC magnetron sputtered Sn/Zn/Cu stack (Cu on bottom) using elemental selenium in the vacuum chamber (avoiding the use of highly toxic H2Se) [49]. Heating of the substrate after metal deposition was performed in a multistep process: 150 1C for

Fig. 3. Conversion efficiency of CZTS-based thin-film solar cells vs. [Cu]/([Zn] + [Sn]) and [Zn]/[Sn] ratios [47]. High conversion efficiency devices are found in a relatively narrow region in this phase diagram. [Figure reproduced with permission from H. Katagiri, K. Jimbo, M. Tahara, H. Araki, K. Oishi, ‘‘The influence of the composition ratio on CZTS-based thin-film solar cells,’’ in Thin-Film Compound Semiconductor Photovoltaics—2009, edited by A. Yamada, C. Heske, M. Contreras, M. Igalson, S.J.C. Irvine (Mater. Res. Soc. Symp. Proc. Volume 1165, Warrendale, PA, 2009), M04-01].

Fig. 4. In-plane distribution of oxygen in CZTS layer as measured by electron probe X-ray micro analysis (EPMA) [48] (a) before 4 hr deionized-water (DIW) soaking and (b) after DIW soaking. Bright point-like areas in (a) correspond to higher oxygen concentration areas. [Figure reproduced with permission from H. Katagiri, K. Jimbo, S. Yamada, T. Kamimura, W. S. Maw, T. Fukano, T. Ito, T. Motohiro, ‘‘Enhanced conversion efficiencies of Cu2ZnSnS4-based thin film solar cells by using preferential etching technique,’’ Appl. Phys. Express 1, 41201, 2008. Copyright The Japan Society of Applied Physics].



Fig. 5. J–V characteristic of a top-performing sputtered CZTS-based thin-film solar cell under AM 1.5 and 100 mW/cm2 illumination and after 5 min of light soaking: Z ¼ 6.77%, Voc ¼ 610 mV, Jsc¼ 17.9 mA/cm2, FF¼ 62% [48]. Inset: Cross-sectional SEM image of a sputtered CZTS absorber layer on a Mo electrode layer. [Figure reproduced with permission from H. Katagiri, K. Jimbo, S. Yamada, T. Kamimura, W. S. Maw, T. Fukano, T. Ito, T. Motohiro, ‘‘Enhanced conversion efficiencies of Cu2ZnSnS4-based thin film solar cells by using preferential etching technique,’’ Appl. Phys. Express 1, 41201, 2008. Copyright The Japan Society of Applied Physics].

30 min and at the final temperature (200, 255, 300, 350 or 400 1C) for 30 min with Se flux. The vacuum (10 6 mbar) employed during the selenization reaction led to significant Sn and moderate Zn loss at higher temperatures ( 4400 1C), while lower temperatures ( o350 1C) led to secondary phases. Further studies of varying the pressure during the selenization reaction (10 5–10 2 mbar range) indicated that higher pressure may be effective in coping with reevaporation losses at 375 1C, but not at temperatures as high as 500 1C [50]. CZTSe devices with 3.2% efficiency (0.23 cm2 device area) were fabricated by sequential sputtering of Cu, Zn and Sn layers, with the selenization treatment provided by elemental selenium in an argon atmosphere at 500 1C for 30 min [31]. To ensure intimate mixing of the Cu, Zn and Sn layers, few-nanometerthick layers of each element were alternatively stacked (rather than single thicker layers for each element). As for CZTS, a slightly Cupoor and Zn-rich absorber was targeted and the optical band gap for the deposited CZTSe film was determined to be 0.94 70.05 eV from external quantum efficiency measurement. The devices were prepared on Mo-coated soda lime glass substrates and finished with chemical bath deposited CdS (70 nm) followed by 50 nm of RF sputtered i-ZnO and 400 nm of indium tin oxide (ITO). Single-step preparation of CZTSe films has also been attempted by RF magnetron sputtering from pressed binary selenide mixtures, with substrates held at various temperatures. Near stoichiometric compositions and dense films were obtained using a substrate temperature of 150 1C, while significant deviation from the targeted composition due to reevaporation was found for higher temperature deposition [51]. The nominally stoichiometric films nevertheless yielded a larger than expected band gap [52], 1.56 eV, as determined from the plot of the square of the absorption coefficient vs. photon energy.

3.1.2. Evaporation Evaporation was the principal deposition method used in much of the successful early work on kesterite absorbers, a natural choice based on the previous success of CIGSSe evaporated materials [53]. The first functional evaporated kesterite device (0.66% power conversion efficiency) was reported by Katagiri et al. [28]. Multilayer

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elemental Cu/Sn/Zn (Zn on bottom of stack) structures were evaporated at a substrate temperature of 150 1C and subjected to a subsequent sulfurization at 500 1C in a Pyrex glass tube using nitrogen with 5% H2S atmosphere. Further refinements of the sequential evaporation and reactive anneal approach included introduction of ZnS, as opposed to Zn, for the bottom precursor metal layer, and increase in the substrate temperature to 400 1C during precursor layer deposition, leading to 2.62% efficiency [54]. Improvement in the heat treatment chamber (stainless steel chamber and a turbo pump vs. a quartz glass tube furnace with a rotary pump), replacement of CdSO4 during CdS deposition with CdI2, and introduction of controlled Na-doping using a Na2S layer between the Mo and the ZnS precursor layer, yielded a significantly improved efficiency of 5.45% for a device with active area of 0.11 cm2 [55]. Open circuit voltage and fill factor were found to improve with use of CdI2 instead of CdSO4 for chemical bath deposition, while Jsc improved with the introduction of a Na2S layer. Other areas of investigation included different sulfurization temperatures and metal ratios [56] and substitution of elemental Sn with SnS2 in the precursor stack, as well as introducing multiple periods of Cu/ SnS2/ZnS (to enhance the interdiffusion/sulfurization process and improve reproducibility) [57]. Finally, devices in which sulfur vapor was used for sulfurization instead of H2S resulted in 1.79% efficiency devices [58]. Shortly after the first report on evaporated CZTS devices deposited using the stacked layer approach, the Stuttgart University group reported on a 2.3% efficient co-evaporated kesterite device [38]. Both selenide (0.6% power conversion efficiency) and sulfide (2.3% power conversion efficiency) devices were fabricated. For CZTS, Cu, ZnS, SnS2 or Sn, and S served as the sources. For CZTSe, Cu, ZnSe, Sn and Se were employed. Substrate temperatures varied between 300 and 600 1C. For temperatures above 400 1C, reevaporation of Sn from the sample became a serious issue, a point further amplified in the work by Ahn et al. [52]. The relatively poor performance for the CZTSe device (one of the first reports of devices with this material) was proposed to arise from poor material quality or secondary phases. At the Hahn-Meitner-Institut (Helmholtz-Zentrum Berlin), studies continued on targeting the development of co-evaporation for the CZTS system [39]. While a promising approach for reaction between ternary Cu–Sn–S and binary ZnS precursors was investigated, challenges related to reevaporation of Zn and Sn were difficult to overcome, resulting in 1.1% efficiency [39]. An alternative approach was subsequently introduced that relied on fast co-evaporation of Cu-rich material from ZnS, Sn, Cu and S sources onto a substrate held at 550 1C and subsequent KCN etching to remove the highly conductive CuS phase [59]. Significantly improved efficiency (4.1%) was thus achieved, with Voc, Jsc and fill factor of 541 mV, 13.0 mA/cm2 and 59.8%, respectively. Similar conditions for co-evaporation (substrate temperature was 550 1C) were explored by Tanaka et al. [60], with a focus on examining the influence of film composition on film structural properties. Larger grain size was noted for increasing [Cu]/([Zn]+[Sn]) ratio for 0.82r[Cu]/([Zn]+[Sn])r1.06. Both Cu-rich and Sn-rich films showed very low resistivity (not suitable for PV device fabrication). Stoichiometric films were achieved in these studies (i.e., the effects of reevaporation of Zn and Sn were minimized), despite high substrate temperatures, perhaps because of the high Sn pressure used during the deposition [61]. Recent work on co-evaporated Cu2ZnSnS4 at IBM [62] employed Cu, Zn, Sn and S sources and a relatively low (110 1C) substrate temperature, followed by a reactive anneal on a 540 1C hot plate for a few minutes, to yield the highest performance to date for a coevaporated pure sulfide kesterite (Fig. 6)—6.8% total area efficiency (Voc¼587 mV, Jsc¼ 17.8 mA/cm2, FF¼65%, device area¼0.45 cm2), on par with the best device performance achieved by a sputtered precursor layer (Z ¼6.8%, Voc¼610 mV, Jsc¼17.9 mA/cm2, FF¼62%,


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Fig. 6. J–V characteristic of top-performing co-evaporated CZTS-based thin-film solar cell under AM1.5 and 100 mW/cm2 illumination: Z ¼6.8%, Voc ¼587 mV, Jsc¼ 17.8 mA/cm2, FF ¼65% [62]. Inset: Cross-section SEM image of IBM CZTS solar cell device with 650-nm-thick absorber, prepared by co-evaporation.

device area¼0.15 cm2; see Section 3.1.1) [48]. More detailed device characterization for the record evaporated CZTS films will be provided in Section 4. Notably, the co-evaporated device employed only a 660-nm-thick absorber layer. The very short reaction times may be instrumental in reducing metal reevaporation from the film, as well as being more attractive for commercial production relative to the hours required for sulfurization in earlier reports [28,54,56,57]. The rapid reaction rate is consistent with recent in situ X-ray diffraction studies that show that, at a temperature of 500 1C, a 2 mm thick kesterite layer should be formed within minutes from a stack of component binary sulfides [63]. Further evaporation-based reports focused on the Se-based CZTSe and include a report on stacked Cu/Zn/Sn (Sn on bottom) materials and subsequent selenization with elemental Se in a sealed quartz ampoule at temperatures of between 250 and 520 1C [64]. Selenization at temperatures greater than 370 1C always resulted in multiphase films that consisted of high quality CZTSe crystals of size up to 2 mm and of a separate ZnSe phase. Similar studies were conducted on co-evaporated (Cu, ZnSe, Sn, Se sources) CZTSe films, with substrate temperature held at between 250 and 400 1C and with a postdeposition treatment in Se vapor at 450 1C [65]. Films processed above 350 1C contained ZnSe as a secondary phase. Other evaporation-related approaches that have been explored for CZTSSe deposition include pulsed laser deposition using an approximately stoichiometric target [24,66,67] and a hybrid sputtering (for Cu/Sn)–evaporation (for Zn/S) approach [68]. For the pulsed laser deposition, recent work has demonstrated a CZTS device efficiency (0.11 cm2 device area) of 0.64% [69].

3.2. Non-vacuum approaches The combination of a successful low-cost deposition approach, as offered by the majority of non-vacuum routes, and the abundant and/or readily available constituents of CZTSSe materials, offers the potential to bring unprecedented growth in photovoltaic manufacturing. Below we detail some of the approaches that are being pursued to demonstrate high-performance devices using a solution-processing approach.

3.2.1. Electrodeposition and other bath-based techniques Electrodeposition is an attractive industrial approach for largescale application and one that is well-established in the electronics industry (e.g., copper interconnection technology in microelectronics). Early work on electrodeposition of CZTS included stacked electroplated Zn/Sn/Cu (Cu on bottom) layers that were subsequently sulfurized with elemental sulfur in a quartz tube furnace (550 1C, 2 hr, argon atmosphere), leading to a power conversion efficiency of 0.8% [70]. The primary limitations on device performance in this study were given as the high series resistance of 10 O cm2, a high shunt conductance of 7 mS cm 2 in the dark (20 mS cm 2 under AM 1.5 conditions), and substantial recombination in the space charge region. In another work, where similar stacks were sulfurized in a nitrogen atmosphere at temperatures of up to 600 1C, issues with poor layer adhesion to the substrate were resolved with a thin layer of Pd on the Mo surface, leading to devices with 0.98% efficiency [71]. An improved process, employing Zn/Cu/Sn/Cu stacks, a rotating disk electrode (RDE) for metal layer deposition to improve the large scale uniformity, and etching of the CTZS films with aqueous KCN to remove copper-rich phases, yielded an efficiency of 3.2% (0.23 cm2 device area) [72]. The problem with unstirred solutions and the impact on lateral film uniformity was further discussed in [73], with the suggestion that the RDE allowed improved uniformity in the precursor layers and consequently in the resulting CZTS films. The desirability of Cupoor composition (possibly [Cu]/([Zn]+ [Sn]) as low as 0.7) was confirmed for electrodeposited films by photoelectrochemical measurements, employing the Eu3 +/2+ redox couple. In the same study, better crystallization was observed by use of H2S for sulfurization in comparison with elemental sulfur [74]. In addition to CZTS, large-grained CZTSe films were obtained by electroplating binary brass and bronze (Cu/Zn and Cu/Sn) stacks followed by selenization in quartz ampoules at 530–560 1C [75]. Substitution of the stacked elemental layer approach with coelectrodeposited metal precursors (i.e., one-step process for the metals), obtained by precise tuning of the bath composition (i.e., copper(II) sulfate pentahydrate, zinc sulfate heptahydrate, tin(II) chloride dehydrate and tri-sodium citrate dehydrate) and plating conditions yielded a 3.16% CZTS device (0.15 cm2 device area) [76]. In situ X-ray diffraction crystallization studies [77] of co-electroplated metal precursors found that the electrodeposited Cu–Zn–Sn precursor exhibited two different formation reactions of Cu2SnS3, depending on the metal ratios in the as-deposited films (Cu-rich vs. Cu-poor). In either case, CZTS formation was completed by the solid-state reaction of Cu2SnS3 and ZnS, which commenced at the relatively high temperature of 570 1C. Some oxide (SnO2 and CuO) impurities were also detected during the in situ study, with the oxygen perhaps being introduced into the system during the electrochemical deposition process. A power conversion efficiency of 3.4% (Voc¼563 mV, Jsc¼14.8 mA/cm2, FF¼41%; device area¼0.5 cm2) was obtained by use of a similar process, which still represents the record efficiency for an electrodeposited CZTS device [78]. Heat treatment for this device was in 5% H2S in argon, with a total processing time of 8 h, including ramping at 2 1C/min and a 2 h hold at the maximum temperature of 550 1C. Although the films were Cu-poor, an additional KCN etching step was carried out before the buffer deposition. Device performance may be limited in these devices by the presence of Zn-rich precipitates (probably ZnS), Zn-poor regions (presumably Cu2SnS3) and voids close to the Mo/CZTS interface. Co-electrodeposition of Cu–Zn–Sn from a choline-based ionic liquid, followed by sulfurization at 450 1C for 1.5 h using Ar as the carrier gas, has also been demonstrated [79]. The choline-based ionic liquid is both water and air stable with negligible vapor pressure up to 130 1C (enabling film deposition at elevated temperature), has an electrochemical potential window of 2.5 V



( 1.2 to –1.25 V), and has a high conductivity. No oxides or hydroxides were detected in the resulting film using X-ray diffraction (in contrast to some reports for aqueous deposition). The measured band gap and absorption coefficient for deposited films were 1.49–1.5 eV and 104 cm 1, respectively, consistent with expected values for CZTS (see Table 1). No solar cell device results have yet been reported. A truly single-step co-electrodeposition of the metals, as well as the sulfur, has also been recently demonstrated [80]. An aqueous room-temperature electrolytic bath (no stirring) was employed, containing CuSO4, ZnSO4, SnSO4 and Na2S2O3 (the sulfur source), and using a conventional three-electrode electrochemical cell assembly. No post-deposition sulfurization step was required, with the as-deposited film being subjected only to a 550 1C anneal in argon for 1 h. For films processed using the 550 1C treatment, a band gap of 1.5 eV was determined from optical absorption measurement. Finally, an alternative bath-type approach has been pursued, based on photochemical deposition of Cu2ZnSnS4 by destabilization of thiosulfate ions in acidic solutions under UV irradiation [81,82]. Under deep-UV irradiation sulfur atoms are released from S2O23 ions in an acidic solution 2H + + S2O23 -S+H2SO3

(1)

The evolved sulfur can react with metal ions to form a metal sulfide Mn + +n/2S+ ne -n/2MS

(2)

The electrons in Eq. (2) are reportedly supplied, for example, by photoexcitation of S2O23 ions 2S2O23 +hu-S4O26 + 2e

(3)

and SO23 + S2O23 +hu-S3O26 + 2e

(4)

The above Eqs. (3) and (4) imply that the process can be controlled by light. Using CuSO4, ZnSO4, SnSO4 and Na2S2O3 for the metal ion and S2O23 sources, H2SO4 to control the pH, and a post-deposition sulfurization in N2 + 5% H2S, Zn-poor CZTS films have been reported [81,82]. Further optimization is required to make films suitable for devices. 3.2.2. Spray pyrolysis Some of the earliest attempts to prepare non-vacuum CZTS employed spray pyrolysis. In 1996, Nakayama and Ito [83] explored phase purity and the effect of the metal ratios using metal chlorides and thiourea as metal and sulfur sources, respectively. The substrates were heated to between 280 and 360 1C during deposition. As-deposited films were considerably sulfur-deficient, using water as the solvent. However, annealing of the films at 550 1C in an argon+5% H2S atmosphere yielded a nominally stoichiometric film. Thermal decomposition of related thiourea-metal halide complexes was further investigated by Madara´sz et al. [84]. The proposed chemical reaction that occurs on a heated glass substrate with the described precursor solution formulation was given as [85] 2CuCl+ ZnCl2 + SnCl4 + 4SC(NH2)2 + 8H2OCu2ZnSnS4 + 4CO2 + 8NH4Cl

(5)

The process was further optimized to yield nominally single phase CZTS in the as-deposited state (no annealing or postdeposition sulfurization treatment) [86]. Near-stoichiometric CZTS films were achieved using a solution containing 0.009 M cupric chloride (rather than CuCl), 0.0045 M zinc acetate, 0.005 M stannic

7

chloride, 0.05 M thiourea and a substrate temperature of 370 1C. The direct optical band gap for these films was measured to be 1.43 eV [86], consistent with established values for CZTS (see Table 1).

3.3. ‘‘Ink’’-based approaches Direct liquid deposition approaches, including solution, particle and mixed solution-particle precursors, are particularly attractive for large-scale manufacturing due to their compatibility with ultrahigh-throughput deposition techniques, such as printing and casting, which are well-established in industry [87]. A sol–gel-like approach for spin-coating CZTS films was reported in 2007 [88]. Copper(II) acetate monohydrate, zinc(II) acetate dihydrate and tin(II) chloride dihydrate were used as sources of the metals and 2-methoxyethanol and monoethanolamine were used as the solvent and stabilizer, respectively. After deposition, the films were dried in air at 300 1C and sulfurized at 500 1C in an N2 +5% H2S atmosphere for 1 h. The resulting films had nearly stoichiometric composition, although with a slight deficiency of sulfur, and the band gap was 1.49 eV (from optical absorption measurement). A similar approach later yielded 1.01% power conversion efficiency for a device with a structure of soda lime glass/Mo/CZTS/CdS/Al:ZnO/Al and in which the transparent conducting oxide (Al:ZnO) layer was also deposited by a sol–gel approach [89]. As the resistivity of the sol–geldeposited Al:ZnO layer was several orders of magnitude higher than is typical for sputtered Al:ZnO, this likely represents one reason for the relatively low device performance. Nevertheless, this represents the first report of a CTZS device in which the absorber, the buffer layer and TCO are all deposited from solution. The same group later improved the above-described device performance to 1.61% efficiency [90]. In addition to pure solution-based approaches, quaternary particle-based precursors have also been examined. As a first example, suspensions, synthesized by reacting metal salts (complexed with triethanolamine) with dissolved elemental sulfur in ethylene glycol, were printed on glass and annealed in S vapor at 550 1C for 10 min to yield Cu2ZnSnS4 films [91]. As-synthesized precursors, prior to heat treatment, comprised reasonably uniform particles with approximate diameter of 200 nm. An issue faced by this approach was the necessity to employ organic binders for crack-free films, which impeded crystal growth. Later work reported synthesis of uniformly size-distributed CZTS nanocrystals using high-temperature arrested precipitation in the coordinating solvent oleylamine, which contained the constituent metal salts and elemental sulfur. The average diameter for the resulting nanoparticles was approximately 10 nm, with composition Cu2.08Zn1.01Sn1.20S3.70. The efficiency for a device prepared with no heat treatment, after nanoparticle deposition from a toluene solution, reached 0.23% [92]. Using a very similar CZTS nanoparticle approach, but subjecting the nanoparticle films to a 500 1C heat treatment in a selenium atmosphere (thereby generating a sintered CZTSSe film) yielded a power conversion efficiency of 0.74% [93]. The low device performance in this latter case was in part attributed to the copper-rich nature of the nanoparticles. A solvothermal approach also yielded (Cu2Sn)x/3Zn1 xS particles with tunable composition, employing reaction between metal chlorides and thiocarbamide in ethanol [94]. Analogous selenide CZTSe nanocrystal synthesis has also been reported [95,96]. Moving to larger particle size, direct ball-milling of Cu2ZnSnS4 sintered material and screen-printing of the resulting isopropanol slurry formulations, followed by hot pressing at 195 1C to remove organic materials, yielded 0.49% efficient devices on molybdenumcoated polyimide substrates [97]. Monograin layer (MGL) solar cells were recently developed, based on growth of high-quality CZTSSe crystals in KI melts (followed by post growth annealing in controlled


8


atmosphere), thus bypassing the major issues encountered by other approaches related to phase purity, grain structure and elemental losses. The monograins were distributed in a single layer of organic resin and specific device processing enabled individual contacting of each CdS coated grain [98,99]. Active area device efficiencies of 2.16% have been reported for a solar cell based on a graphite/CZTSe/CdS/ ZnO/Al:ZnO structure [99]. Recently, the performance of a monograin CZTSSe device (S:Se ratio is 3:1; device area defined by 0.04 cm2 graphite contact spot) has been improved to Z ¼5.9%, Voc¼622 mV, Jsc¼15.9 mA/cm2 and FF¼60% [100]. While CZTSSe inks can be comprised of pure solutions or quaternary particles dispersed in a carrier fluid, mixed inks based on a combination of dissolved and solid chalcogenides were developed at the IBM T.J. Watson Research Center [26]. The symbiosis of a dissolved component with binding action (i.e., no need for an organic binding agent) and solid particles, allowing the introduction of insoluble materials proved to be a versatile approach for high-purity ink formulations. As a demonstration of this approach for CZTSSe, many metal chalcogenides, as well as elemental sulfur and selenium, can be effectively dissolved in hydrazine solutions [101–105]. For example, copper(I) sulfide and tin(II) sulfide/selenide readily dissolve in hydrazine with extra chalcogen added [102,104]. However, zinc sulfide/selenide does not readily dissolve in hydrazine under ambient conditions. Therefore hybrid solution-particle inks were prepared by dissolving Cu2S and SnS(Se) in hydrazine and adding elemental zinc powder to the solution [26]. The reaction between the zinc and the metal chalcogenide solution resulted in the in situ formation of ZnS(Se)N2H4 particles dispersed in the Cu–Sn–S–Se solution. Upon spin coating or doctor blading of precursor films using the hybrid slurries, followed by a short heat treatment at 540 1C, a uniform single phase CZTSSe film with good grain structure (Fig. 7, inset) was achieved. Using this approach, CZTSSe devices based on a soda lime glass/Mo/CZTSSe/CdS/ZnO/ITO structure have been prepared with total area (0.44 cm2) power conversion efficiency of 9.7% (Fig. 7; device results certified by Newport Corp.), which currently represents the highest demonstrated efficiency for a thin-film kesterite-based PV device (see Fig. 2). Note that the data in Fig. 7 is on a separate sample from the original one used for the NREL certification (which yielded essentially the same efficiency) [26].

Fig. 8. Cross section TEM images of hydrazine-processed Cu2ZnSn(Se1 xSx)4 devices with (a) x E0.1 and (b) xE0.4 [26]. In each case, the device structure consists of glass substrate/Mo back contact (800 nm)/spin-coated CZTSSe/CdS (60 nm)/ZnO (80 nm)/ITO (130 nm). The Mo–Se interfacial layer between the glass/Mo substrate and the CZTSSe is noted. (c) Transmission EDX compositional profiling showing uniform metal composition and higher S content in the x E0.4 sample. [Figure reproduced with permission from T. K. Todorov, K. B. Reuter, D. B. Mitzi, ‘‘High-efficiency solar cell with earth-abundant liquid-processed absorber, Adv. Mater., 22. E156-E159, 2010. Copyright Wiley-VCH Verlag GmbH & Co. KGaA.].

Control over the S:Se ratio, and consequently the band gap of the CZTSSe films, can be achieved by changing the ratio of these elements in the initial precursor solution, as well as by introducing elemental S or Se during the high temperature heat treatment (Fig. 8). Also, recent work demonstrated successful transfer of this approach from pure hydrazine solutions to more industry-friendly aqueous-based (diluted hydrazine) solvent vehicles, with initial device efficiencies already reaching 8.1% (0.45 cm2 device area) [106, 107]. Further discussion of device characterization on the record CZTSSe devices prepared from hydrazine solutions will be given in the next section.

4. Characterization of kesterite materials and devices While there have been numerous studies recently on the fabrication of CZTSSe films prepared in various ways, there are relatively fewer reports providing detailed electrical/optical characterization of these films (beyond simple light/dark J–V and external quantum efficiency) in an effort to enable the level of device understanding currently available for CIGSSe devices. Many of the discussions below on electrical characterization will therefore focus on a specific set of CZTSSe devices processed at IBM using solution- and vacuum-based approaches [26,62,106,107]. 4.1. Structural characterization

Fig. 7. Certified (by Newport Corp.; certification #0122) CZTSSe light J–V curve, under simulated AM1.5 light, for a hydrazine-processed solar cell (0.44 cm2 device area): Z ¼9.7%, Voc ¼0.448 V, Jsc¼ 32.2 mA/cm2, and FF¼ 67%. Inset: Cross-section SEM image of IBM CZTSSe solar cell device with absorber prepared by spin-coating from a hydrazine-based slurry.

Demonstration of a single phase CZTSSe film generally represents the first task in any physical property characterization study and X-ray diffraction (XRD) is typically the first tool used for this purpose [40,63]. Fig. 9 shows how the diffraction pattern for the CZTSSe kesterite differs from that for the analogous CIGSSe chalcopyrite [107]. The refined lattice constants for the particular two solution-processed films with comparable Se content (Fig. 9a ˚ c ¼11.30 A˚ for CZTSSe vs. and b) are very similar (a ¼5.67 A, ˚ c¼11.35 A˚ for CIGSSe). However, the different space a¼5.68 A, group for CZTSSe (I-4 vs. I-42d), arising from the ordering of the



9

Besides X-ray analysis, Raman spectroscopy study is useful to investigate the phase purity and composition of kesterite and stannite materials [27]. Fig. 10 shows the Raman spectra of CZTSSe thin films with varying S:Se ratio. The full-sulfur CZTS data is obtained from a thermally evaporated film [62], while the other data sets are obtained from hydrazine-processed films [107]. The major peaks in the Raman spectra arise from the A1 vibration mode of the lattice, where the group VI atoms (S or Se) vibrate while the restp offfiffiffiffiffiffiffiffiffiffiffiffiffi atoms remain fixed and the vibration frequency is given by v ¼ k=MVI , where k is the lattice vibration spring constant and MVI is the atomic mass of the group VI element. For the pure sulfide sample (a), the major peak appears at 338 cm 1, which is consistent with previous reports for CZTS (338 cm 1 major peak with additional peaks at 287 and 368 cm 1) [27,108]. For the high selenium sample (c), the major peaks shift to lower wave numbers, 196, 172 and 238 cm 1, in agreement with those reported in Ref. [108] for monograin powder samples (major peak at 196, medium peak at 173 and a weak peak with position between 231 and 253 cm 1). The ratio of the A1 peak frequencies in samples (a) and (c), nS =nSe ¼1.69, is approximately equal to the ratio of the square ffi pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi root of the atomic masses, MSe =MS ¼ 1:57, as expected from the formula above, with any discrepancy attributed to a slightly different spring constant. The Raman profile for the sample with intermediate value of [S]/ ([S]+ [Se]) ratio exhibits a bimodal behavior [108], similar to that observed in CuInSe2 xSx, with a linear combination of the 332 and 196 cm 1 peaks and some noticeable peak broadening and slight shifts in peak position (Fig. 10, top panel) [107]. Taking the ratios of

Fig. 9. XRD profiles (Cu Ka radiation) of typical (a) CIGSSe and (b) CZTSSe hydrazineprocessed absorber layers with [S]/([S] + [Se])E 0.1 on Mo-coated glass [26,105,107]. The data is plotted with a logarithmic intensity scale to highlight any potential impurity phases (data for CZTSSe is shifted along the y-axis for clarity). Weak broad features at approximately 2y ¼321 and 561 arise from the MoSe2 interfacial layer. Peaks from the Mo back contact are also noted. The indices enclosed within boxes correspond to peaks that are allowed for the I-4 space group but not for I-42d. In (c), the X-ray diffraction pattern for a hydrazine-processed CZTS film (no Se) is presented. The shift in the diffraction peaks to higher 2y, with very similar diffraction pattern, upon substitution of Se with S is highlighted by the dashed lines between b) and c).

metals, allows additional peaks to be observed in the XRD profile for the kesterite (as highlighted with boxes around the indices in Fig. 9). In both cases, other than the Mo from the substrate and a MoSe2 interfacial phase, no impurity peaks are observed in the diffraction patterns for these hydrazine-processed devices. XRD has also been used to track the shift in lattice constants as Se is replaced by S in the CZTSSe series, with no evidence of structure change or secondary phases being formed over this range (e.g., see Fig. 9c for the diffraction pattern for CZTS) [26]. Simple examination of a powder XRD pattern, however, cannot determine unambiguously whether a film is single phase, since many of the binary and ternary intermediates adopt an analogous zinc-blende-related structure and have peaks that overlap with the CZTSSe XRD pattern (e.g., Zn(S,Se), Cu2Sn(S,Se)3) [40,63].

Fig. 10. Raman spectra for CZTSSe films with varying [S]/([S] +[Se]) ratio (x) prepared on Mo-coated glass: (a) x¼ 1 (b) xE0.4 and (c) xo 0.1. The pure sulfide CZTS data (a) was taken on a co-evaporated film as described in [62], while the selenide-containing CZTSSe data was collected from hydrazine-processed devices [26,107]. The Raman spectra are taken using 625 nm laser excitation. Top panel: Bimodal behavior plot of A1 peak as a function of x.


10


the peak heights for the sulfide-based ( 330 cm 1) and seleniumbased ( 210 cm 1) peaks (background subtracted), leads to a [S]/([S] +[Se]) ratio of 0.4 for sample (b), which is similar to the value determined by secondary ion mass spectroscopy (SIMS) profiling and TEM -EDS studies. Furthermore, these spectra also indicate the absence of any major impurity phases, such as Cu2S or Cu2Se [109]. Finally, transmission electron microscopy (TEM) studies and secondary ion mass spectrometry (SIMS) depth profiling can prove useful to demonstrate uniform film composition between grains and also to evaluate the film interfaces (Fig. 8). TEM is particularly useful in characterizing the thickness of the Mo–Se-based interfacial layer that forms between the CZTSSe film and the Mo-coated glass substrate. Given the close structural relationship between the final CZTSSe compound and a number of the expected intermediary phases (e.g., ZnS, Cu2SnS3), it is preferable to confirm phase homogeneity and purity by using a combination of approaches, such as those mentioned above. From the reasonably uniform metal compositions seen in Fig. 8 and the X-ray and Raman data shown in Figs. 9 and 10 for similarly prepared samples, it appears that nominally single phase CZTSSe samples can be prepared using the hydrazine-based process [26,107]. 4.2. Electrical properties The CZTSSe system has the distinct advantage of enabling tuning of the direct band gap, by varying the S:Se ratio, from close to 1 eV for the pure selenide compound to approximately 1.5 eV for the pure sulfide material (very useful range for terrestrial solar illumination; see Table 1). From band structure calculations, the upper valence band in CZTSSe consists of antibonding Cu 3d and S(Se) 3p(4p) orbitals, Table 2 Reported carrier density, mobility and resistivity values of selected CZTSSe films, as determined by Hall measurement. Compound

Carrier density /cm3

Mobility cm2/Vs

Resistivity X cm

References

Cu2ZnSnS4 Cu2ZnSnS4 Cu2ZnSnS4 Cu2ZnSnS4 Cu2ZnSnSe4 Cu2ZnSnSe4

8.2 1018 8.0 1018 45 1019 3.8 1018

6.3 6.0 o 0.1 12.6

2.1 1017

39.7

0.16 0.13 41.3 – 0.1–0.8 0.74

[30] [68] [29] [97] [112] [24]

whereas the bottom of the conduction band consists of antibonding orbitals of Sn 5s and S(Se) 3p(4p). In these calculations, the Zn atom does not affect the valence band maximum or the conduction band minimum [25]. However, recent calculations predict that Sn substitution (due to disorder or non-stoichiometry) on the Zn site can create a detrimental deep level due to a +IV-+II transition (i.e., multivalency) inside the band gap [110]. Note that some early reports for the band gap of the pure selenide compound indicated band gaps of Z1.4 eV, a value very similar to the band gap value for CZTS [16,24,65]. Presumably these earlier high values of the band gap for CZTSe can be attributed to ZnSe or other impurity phases, which can be difficult to distinguish by X-ray diffraction [51]. Selected reported electrical properties of CZTSSe compounds are shown in Table 2 for films prepared using various methods, such as sputtering, co-evaporation and pulsed laser deposition (see Section 3). The listed CZTSSe Hall mobilities are in the range of those reported for CIGSSe materials (3–22 cm2/Vs) [111]. Even though minority carrier mobility is a more relevant parameter for photovoltaics performance (see Section 4.4.3), the reported Hall mobilities are still expected to reflect film quality. Generally these films are prepared on insulating substrates to prevent parallel conduction in the Hall measurements. Thus the reported electrical properties may not reflect the characteristics of photovoltaicsquality films grown on metal-coated substrates due to differences in grain structure and incorporated impurities. Electrical characterization, including mobility and carrier density, are still therefore a need in the CZTSSe field, with special care being taken that the films be single phase and deposited in the same configuration as that used in devices.

4.3. Device characteristics As an emerging thin-film technology with a promising demonstrated efficiency of 9.7% [26], it is crucial to understand the key device issues that limit the performance of the current generation of CZTSSe cells. We present in this section data from the highestperformance (defined here as 46% power conversion efficiency) CZTSSe cells reported in the literature (Table 3) and compare them with a solution-processed CIGSSe cell [105,113,114] with power conversion efficiency of 13.8%, which can serve as a high-efficiency benchmark. Samples that have undergone more thorough evaluation are also given labels to facilitate further discussion. CZTSSe cells ZA and ZB were fabricated using a hydrazine-based process

Table 3 Device parameters reported for high performance CZTSSe cells with various [S]/([S] +[Se]) ratio. Cell

ZA ZB ZD ZC ZR GR

Comp

CZTSSeb CZTSSe CZTSSec CZTSSeb CZTSSe CZTSSe CZTS co-evap. CZTS sputtering CIGS

g

Note

IBM hydr.-based IBM hydr.-based IBM hydr.-based IBM hydr.-based IBM hydr.-based IBM mixed solv. IBM x¼ 1 NNCT x ¼1 IBM hydr.-based

xo 0.1 xo 0.1 xE 0.4 xE 0.4 xE 0.4

P g (%)

FF (%)

Voc (mV)

Eg (eV)

Eg/q Voc (mV)

Jsc (mA/cm2)

RS,L (Ocm2)

GS,L (mS/cm2)

A

(%)

J0 (mA/cm2)

9.30b 8.74 9.7 9.66 9.50 8.13 6.81 6.77 13.8

– 10.0 11.0d 11.3d 11.2 10.1 7.81 – 14.9

62.4 58.8 67.2 65.4 64.3 60.0 65.0 62.0 72.4

412 380 448 516 499 563 587 610 578

1.03 1.04 1.10d 1.21d 1.21 1.29 1.45 1.45 1.14

618 660 652 694 711 727 863 840 562

36.4 39.1 32.2 28.6 29.6 24.1 17.8 17.9 33.1

– 1.23 1.56d 2.42d 2.24 4.42 3.40 – 1.33

– 7.32 1.81d 2.69d 1.75 1.94 1.90 – 0.30

– 1.49 1.32d 1.39d 1.32 1.41 1.70 – 1.33

– 1.80e-3 6.14e-5d 1.44e-5d 1.37e-5 4.18e-6 7.3e-5 – 1.31e-6

Ref.

[26] a a

[26] [114] [106] [62] [48] [114]

Notes: All device areas are 0.45 cm2, except for device ZR, which has an active area of 0.15 cm2. P Z is the pseudo-efficiency taken from the Jsc–Voc data shifted by Jsc at 1 sun, Eg is the band gap as determined from the external quantum efficiency data, RS,L is series resistance measured using the light-J–V and Jsc–Voc data [115], GS,L is shunt conductance from light J–V. A and J0 are the diode ideality factor and reverse saturation current determined from the Jsc–Voc data. a

Data presented in this study. NREL certified results. c Newport certified result. d Additional measurement entries performed at IBM after receiving certified results. b



11

Fig. 11. Light and dark J–V data of the high performance CZTSSe cells ZA, ZB and ZC. The cross-over points between the light and dark J–V curves are marked with solid circles. The dashed curve represents the 10% efficiency contour.

analogous to that described in Ref. [26], ZD is fabricated by an aqueous diluted-hydrazine process [106,107] and ZC is prepared by thermal co-evaporation [62]. They have varying [S]/([S] +[Se]) compositions ranging from o0.1 to 1, spanning a band gap range of 1.03–1.45 eV. The reference CIGSSe cell is labeled GR. In each case, the device structure is glass/Mo/CZTSSe (or CIGSSe)/CdS/ZnO/ ITO, with Ni–Al contact grid and MgF2 anti reflective (AR) coating on top (except for the sputtered film ZR, which has no AR coating). Fig. 11 presents the electrical characterization on devices ZA, ZB and ZC, consisting of light J–V and dark J–V for all three cells. The external quantum efficiency (EQE) spectra in Fig. 12a can be used to determine the absorber layer band gap by fitting a plot of [E ln(1 EQE)]2 vs. E near the band edge, where E¼hc/l (h is Planck’s constant, c is the speed of light and l is the wavelength of light), as shown in Fig. 12b. Fig. 12c shows the open circuit voltage plotted as a function of band gap. From these basic device characteristics several shortcomings of the CZTSSe cell become apparent by comparing their device performance with the reference CIGSSe cell [114]. First, the open circuit voltages (Voc) of the CZTSSe cells are low compared to their respective band gaps (Fig. 12c). A good thin-film solar cell should have the difference between Voc and Eg/q nominally close to 0.5 V, where Eg is the band gap and q is the electron charge [116]. For the Reference CIGSSe cell, GR, Voc–Eg/q is 0.56 V. However as shown in Table 3, Eg/q–Voc for the CZTSSe cells are all larger than 0.6 V, and this difference becomes larger for higher CZTSSe band gap. This trend is also apparent in the Fig. 12c plot, where the increase in Voc is less than the increase in band gap. Second, we observe that the fill factors of the CZTSSe cells are also relatively low, which can be caused by high series resistance as indicated by the series resistance values (RS,L) in Table 3. Third, from Fig. 12a, we observe that the quantum efficiency response in the long wavelength region near the band gap edge is relatively low for the CZTSSe cells. Overall, the best CZTSSe cell (ZB) used for more comprehensive characterization has a band gap of 1.2 eV, with room temperature efficiency of 9.7%. It also has a good ideality factor of A ¼1.39 (determined using the Jsc–Voc method [117] utilizing a continuous neutral density filter to obtain dense data points near the 1 sun Jsc– Voc point), which is in the range of the ideality factors of high performance CIGSSe solar cells [118]. Note that, in general, the sputtered CZTS cell (ZR) has very close performance parameters with the co-evaporated cell ZC (Table 3).

Fig. 12. (a) Quantum efficiency of the high performance CZTSSe cells at zero bias. Dashed-curve: Reflectance curve of cell ZB, which is typical for all cells coated with a MgF2 anti reflective coating. (b) Band gap determination plot using EQE data. (c) Voc vs. band gap. The dotted line is the Eg/q—0.5 V line that indicates a nominally good Voc for thin film solar cells. The star point represents the data of the reference CIGS cell (GR).

4.4. Factors limiting device performance Detailed CZTSSe device characterization studies on cells with varying [S]/([S] + [Se]) content from 0% to 100% have been performed to pinpoint the major loss mechanisms that limit the device performance [62,114]. Below we summarize the key findings from these studies with respect to the device shortcomings mentioned above. 4.4.1. Open circuit voltage To investigate the cause of deficiency in Voc, temperature dependence data has been collected to obtain information on the dominant recombination process in the CZTSSe cells. The relationship between Voc and temperature is generally given as [116] Voc ¼

EA AkT J00 ln q q JL

ð6Þ

where EA, A, J00 and JL are the activation energy of the dominant recombination mechanism, diode ideality factor, reverse saturation current prefactor and the photocurrent, respectively. The Voc vs. T data should yield a straight line and extrapolate to the activation energy EA/q at T¼0 K, as long as A, J00 and JL are temperature independent. For more rigorous determination of EA, without assuming EA, A, J00 and JL to be temperature independent, one could recast Eq. (6) into Aln J0 ¼ EA =kT þ A ln J00 ,


ð7Þ

12


conduction band edge of the absorber layer is higher than that of the buffer layer [119]. Note that Krustok et al. [98] observed a Voc vs. T plot that has a 0 K intercept at the band gap value in monograin CZTSSe solar cells, indicating that differences in the absorber layer synthesis may contribute to the quality of the buffer/absorber interface.

where A ln J0 is the ‘‘corrected saturation current density.’’ EA can then be extracted from the plot of A ln J0 vs. 1/T. In a well-behaved solar cell, the determined activation energy should equal to the band gap of the absorber layer, indicating that the main recombination mechanism arises in the space charge region (SCR) region of the absorber layer. The Voc vs. T plots are shown in Fig. 13a for three samples, ZA, ZB and ZC, and each has an intercept lower than the respective band gap value. This fact is confirmed from the A ln J0 vs. 1/T plots (Fig. 13b), which yield EA values of 0.86, 1.05 and 0.96 eV for cell ZA, ZB and ZC, respectively, and which are low compared to their respective band gap values (1.06, 1.21 and 1.45 eV). This fact suggests that the main recombination mechanism in these CZTSSe cells is dominated by interface recombination, thereby suggesting an important limiting factor for Voc. In contrast, the reference CIGSSe cell (GR) has T¼0 K intercept at the band gap value, indicating a normal SCR-dominated recombination [114]. This dominant interface recombination for the CZTSSe devices could occur due to a defective buffer/absorber interface, surface traps in the interface or a ‘‘cliff’’-type band alignment, where the

4.4.2. Series resistance We observed earlier that the CZTSSe cells exhibited relatively low fill factors, which is normally associated with high series resistance. A full temperature dependence sweep (light and dark J–V and Jsc–Voc) elucidates this problem. ZB is used as an example, while the other cells (ZA and ZC) show similar behavior [62,114]. Fig. 14 shows the light J–V, pseudo J–V and dark J–V at three different temperatures. The pseudo-J–V data are obtained from Jsc–Voc data shifted down by Jsc at 1 sun [120]. The pseudo J–V data is essentially the intrinsic J–V of the solar cell, free from the effect of series resistance. The Jsc–Voc or pseudo J–V data are used to extract pseudo efficiency (P Z), pseudo fill factor (PFF), diode ideality factor (A), reverse saturation current (J0) and

Fig. 13. (a) Temperature dependence data of the open circuit voltage and its linear extrapolation to T¼ 0 K. Data for the CZTSSe devices have 0 K intercepts that do not reach the band gap value. (b) The corrected saturation current density vs. inverse temperature plot. The slope yields the activation energy of the recombination process, EA.

Fig. 14. (a) Temperature dependent light J–V measurement for cell ZB. Dashed curve: Pseudo J–V (shifted Jsc–Voc), which represents the intrinsic J–V of the cell free from series resistance. (b) Temperature dependent dark J–V curves for cell ZB.



13

Fig. 15. The series resistance problem in a CZTSSe cell, exhibited in a temperature dependence experiment [114]. Comparison data from a high performance reference CIGSSe cell (GR) is also shown. Data is shown for cell ZB, with all other CZTSSe cells (ZA and ZC) showing similar behavior. Temperature dependence of (a) device efficiency, (b) short circuit current, (c) fill factor and (d) dark series resistance.

series resistance RS,L (together with light-J–V data). We observe that the light J–V curve deteriorates at lower temperature (lower fill factor) while the shape of pseudo J–V curves remain intact. This observation indicates that intrinsically the cell still has good performance, but the series resistance increases and quenches the fill factor. This trend can be quantified further as shown in Fig. 15, where a comparison is made between CZTSSe cell ZB and the high-performance solution-processed CIGSSe cell (GR). Fig. 15a shows that the efficiency of the GR cell increases almost monotonically towards low temperature and saturates at 20% below 160 K. However the CZTSSe cell (ZB) shows a dramatically different behavior, as its efficiency collapses towards zero at lower temperature. Similar behavior happens in the short circuit current data and fill factor as shown in Figs. 15b and c. We know from Fig. 13a that the open circuit voltage is steadily increasing (or saturates for the CZTSSe cells) at low T. Thus, this efficiency collapse is caused primarily by the collapse in fill factor followed by short circuit current in the progression towards lower temperature. The collapse in fill factor effect can in turn be related to the dramatic change in series resistance as shown in Fig. 15d. This dark series resistance RS,D is extracted from the dark J–V data, using a standard diode analysis of a thin-film solar cell, as described in Refs. [121,122]. Indeed, the dark series resistance RS,D diverges towards lower temperature for the CZTSSe cells. RS,D of the ZB cell increases by as much as 5000 from 340 to 120 K, compared to only 3 for the GR cell. One probable cause of this diverging series resistance at low temperature is the presence of a blocking back contact at the interface between the Mo and CZTSSe [114], which can suppress the majority carrier (hole) transport. In fact, a similar problem exists in many high band gap chalcogenide thin-film solar cells such as CdTe [123,124]. Another observation consistent with the presence of a blocking back contact is the cross-over behavior between the light and dark J–V data, as indicated by solid circles in Fig. 11 [123]. The barrier in these CZTSSe cells could be associated with a relatively thick MoSe2 layer ( 300 nm) found in hydrazineprocessed CZTSSe cells (e.g., see TEM images in Fig. 8) or with the fact that the back contact is simply not optimized for CZTSSe (as opposed to CIGSSe). The exact band alignment at this back contact region still requires further study, both using experimental study

Fig. 16. Ratio of EQE at voltage biases of 1 and 0 V. Collection efficiency becomes worse for higher band gap CZTSSe solar cells.

and modeling, in order to understand the limitations imposed by this component of the device.

4.4.3. Carrier collection efficiency The EQE data from Fig. 12a show a relative lack of long wavelength response for the CZTSSe cells. To examine the carrier collection efficiency of these cells, the ratio of the EQE at reverse bias ( 1 V) and zero bias is shown in Fig. 16. This ratio increases to more than one for higher band gap CZTSSe (ZB and ZC), which indicates that a larger depletion width improves the collection efficiency. Similar behavior has been observed in larger band gap CIGSSe devices [125]. This observation implies that the collection efficiencies in these cells are most likely limited by very low carrier lifetime. Low minority carrier lifetime also limits the Voc in these cells. This low lifetime could arise from a high defect density in the absorber layer or could simply be a consequence of high recombination loss at the back contact or at the front interface, as has been discussed previously (see Section 4.4.1). Note that time-resolved photoluminescence data have also been collected on these samples and indeed we observe that the CZTSSe cells (ZA and ZB) have a


14


short minority carrier lifetime ( 1.2 ns) compared to the 6 ns lifetime of the reference CIGSSe cell (GR) [114].

5. Future prospects and conclusion Kesterite-related materials, and in particular CZTSSe compounds, represent an important opportunity to combine both appropriate electronic properties and abundant and/or readily available elements in a high volume and low-cost technology. A key enabler to make this vision a reality is the achievement of high power conversion efficiencies, on par with other leading thinfilm technologies. With the progressive improvement that has been achieved [37] and the most recent demonstration of efficiencies near 10% (and pseudo-efficiencies, removing the influence of series resistance, above 11%—see Table 3) [26], performance metrics are entering a range that can be considered interesting for commercialization. In addition, studies are beginning to be conducted (both experimental and theoretical) that may provide the necessary detailed understanding of these materials to build a successful technology [114]. Finally, many approaches have been attempted to deposit CZTSSe thin films, with the major impediments to success being the complex and not fully elucidated phase diagram (leading to difficulty in preparing single phase films), volatility of Sn and Zn (in addition to S and Se), and the need to optimize device interfaces. Despite these difficulties successful film deposition and device fabrication are beginning to appear for both vacuum-based and solution-based deposition. There are still numerous hurdles that must be overcome before a truly competitive kesterite technology can be validated. Device efficiencies must preferably be raised to values more on par with CdTe and CIGSSe based devices (at least 414% power conversion efficiency). In order to achieve this, a more detailed understanding of the inter- and intra-layer interfaces (e.g., grain boundary, buffer layer-CZTSSe and back contact-CZTSSe) must be achieved, as well as a corresponding reduction in the generally high device series resistance and an elucidation of the nature of defects in the absorber layer (whether electronically active or not). A complete picture of the band alignment at the buffer/absorber layer interface and at the back contact is also critical to understanding the Voc limitation and the possible back contact issue. The long-term stability of CZTSSe and other kesterite PV devices under heat, light and moisture should be established and compared with the CdTe and CIGSSe counterparts. Finally, a truly low-cost and highthroughput approach for deposition needs to be established, without sacrificing on performance, in order to meet a o$1/W price target. The initial results on CZTSSe films and devices presented in this review point to successes in the early stages of addressing these issues and consequently suggest a promising future for this class of materials. References [1] 2008 Solar Technologies Market Report, US Department of Energy, NREL Report no. TP-6A2-46025, DOE/GO-102010-2867, 2010, pp. 1–131. [2] Surprise–surprise. Solar cell production for 2009 hits 12 GW, Photon International, March 2010, 2010, pp. 176–199. [3] M.A. Green, K. Emery, Y. Hishikawa, W. Warta, Solar cell efficiency tables (version 36), Prog. Photovolt.: Res. Appl. 18 (2010) 346–352. [4] I. Visoly-Fisher, S.R. Cohen, A. Ruzin, D. Cahen, How polycrystalline devices can outperform single-crystal ones: thin film CdTe/CdS solar cells, Adv. Mater. 16 (2004) 879–883. [5] Y. Chiba, S. Kijima, H. Sugimoto, Y. Kawaguchi, M. Nagahashi, T. Morimoto, T. Yagioka, T. Miyano, T. Aramoto, Y. Tanaka, H. Hakuma, S. Kuriyagawa, K. Kushiya, Achievement of 16% milestone with 30cmx30cm-sized CIS-based thin-film submodules, in: Proceedings of the 35th IEEE Photovoltaic Specialist Conference, 2010, pp. 164–168. [6] First Solar press release, Second Quarter 2010 Financial Results, /http:// investor.firstsolar.comS (accessed 08.23.2010).

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