Compound Copper Chalcogenide Nanocrystals - ACS Publications

Review pubs.acs.org/CR

Compound Copper Chalcogenide Nanocrystals Claudia Coughlan,† Maria Ibáñez,‡ Oleksandr Dobrozhan,‡,§ Ajay Singh,∥ Andreu Cabot,*,‡,⊥ and Kevin M. Ryan*,† †

Department of Chemical Sciences and Bernal Institute, University of Limerick, Limerick, Ireland Catalonia Energy Research Institute - IREC, Sant Adria de Besos, Jardins de les Dones de Negre n.1, Pl. 2, 08930 Barcelona, Spain § Department of Electronics and Computing, Sumy State University, 2 Rymskogo-Korsakova st., 40007 Sumy, Ukraine ∥ Materials Physics & Applications Division: Center for Integrated Nanotechnologies, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States ⊥ ICREA, Pg. Lluís Companys 23, 08010 Barcelona, Spain ‡

ABSTRACT: This review captures the synthesis, assembly, properties, and applications of copper chalcogenide NCs, which have achieved significant research interest in the last decade due to their compositional and structural versatility. The outstanding functional properties of these materials stems from the relationship between their band structure and defect concentration, including charge carrier concentration and electronic conductivity character, which consequently affects their optoelectronic, optical, and plasmonic properties. This, combined with several metastable crystal phases and stoichiometries and the low energy of formation of defects, makes the reproducible synthesis of these materials, with tunable parameters, remarkable. Further to this, the review captures the progress of the hierarchical assembly of these NCs, which bridges the link between their discrete and collective properties. Their ubiquitous application set has cross-cut energy conversion (photovoltaics, photocatalysis, thermoelectrics), energy storage (lithium-ion batteries, hydrogen generation), emissive materials (plasmonics, LEDs, biolabelling), sensors (electrochemical, biochemical), biomedical devices (magnetic resonance imaging, X-ray computer tomography), and medical therapies (photochemothermal therapies, immunotherapy, radiotherapy, and drug delivery). The confluence of advances in the synthesis, assembly, and application of these NCs in the past decade has the potential to significantly impact society, both economically and environmentally.

CONTENTS 1. Introduction 1.1. General Overview 1.2. Outline of Review 2. Crystal Phases and Stoichiometries 2.1. Binary Compounds 2.1.1. Cu−S System 2.1.2. Cu−Se System 2.1.3. Cu−Te System 2.2. Ternary Compounds 2.2.1. Cu-III-VI 2.2.2. Cu-IV-VI 2.2.3. Cu-V-VI 2.3. Quaternary Compounds 2.3.1. Cu-III-VI 2.3.2. Cu-II-IV-VI 3. Functional Properties 3.1. Electronic Properties 3.1.1. Band Structures 3.1.2. Defects 3.1.3. Surface Defects/Composition 3.2. Photoluminescence 3.3. Plasmonics 3.4. Nonlinear Optics 3.5. Magnetism © 2017 American Chemical Society

4. Solution Synthesis Approaches 4.1. Colloidal Synthesis Method 4.1.1. LaMer Classical Nucleation Theory 4.1.2. Size Focusing/Defocusing Concepts 4.1.3. Newer Considerations in Nucleation 4.1.4. Hard-Soft-Acid−Base (HSAB) Theory 4.2. Solvothermal/Hydrothermal Method 4.3. Template-Directed Synthesis Method 4.4. Kirkendall Effect-Induced Method 4.5. Cation Exchange Method 5. Synthesis of Copper Chalcogenide Nanocrystals 5.1. Binary Semiconductor Nanocrystals 5.1.1. Copper Sulfide (Cu2−xS) 5.1.2. Copper Selenide (Cu2−xSe) 5.1.3. Copper Telluride (Cu2−xTe) 5.2. Ternary Semiconductor Nanocrystals 5.2.1. Copper Indium Sulfide (CuInS2) 5.2.2. Copper Indium Selenide (CuInSe2) 5.2.3. Copper Gallium Sulfide (CuGaS2) 5.2.4. Copper Gallium Selenide (CuGaSe2) 5.2.5. Copper Germanium Selenide (Cu2GeSe3)

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Received: June 14, 2016 Published: April 10, 2017 5865

DOI: 10.1021/acs.chemrev.6b00376 Chem. Rev. 2017, 117, 5865−6109

Chemical Reviews 5.2.6. Copper Tin Sulfide (Cu2SnS3) 5.2.7. Copper Tin Selenide (Cu2SnSe3) 5.2.8. Copper Antimony Sulfide (CuSbS2) 5.2.9. Other Ternary Compositions 5.3. Quaternary Semiconductor Nanocrystals 5.3.1. Copper Indium Gallium Sulfide (CIGS) 5.3.2. Copper Indium Gallium Selenide (CIGSe) 5.3.3. Copper Zinc Tin Sulfide (CZTS) 5.3.4. Copper Zinc Tin Selenide (CZTSe) 5.3.5. Other Quaternary Compositions 5.4. Semiconductor Nanocrystals with Two Chalcogens 5.4.1. Copper Sulfur Selenide (Cu2−xSySe1−y) 5.4.2. Copper Indium Sulfur Selenide (CuIn(S1−xSex)2) 5.4.3. Copper Zinc Tin Sulfur Selenide (Cu2ZnSn(SxSe1−x)4) 6. Nanocrystal Interactions and Assembly Strategies 6.1. Assembly Introduction 6.2. Interactions between Nanocrystals 6.2.1. van der Waals Interactions 6.2.2. Dipole−Dipole Interactions 6.2.3. Electrostatic Interactions 6.2.4. Depletion Interactions 6.2.5. Capillary Interactions 6.3. Role of Surface Ligands in Nanocrystal Assembly 6.4. Strategies of Nanocrystal Assembly 6.4.1. Drying Mediated Assembly 6.4.2. Self-Assembly at the Interface (Liquid−Liquid and Liquid−Air) 6.4.3. Assembly in Solution 6.4.4. Directed Assembly Process (Electrophoretic Deposition) 7. Photovoltaic Applications 7.1. Thin Film Solar Cells Based on Sintered NCs 7.1.1. Deposition Technologies 7.1.2. Ink Composition 7.1.3. Sintering Atmosphere 7.1.4. Graded Solar Cells 7.1.5. Surface Ligands/Carbon 7.1.6. Other Architectures/Components 7.1.7. Other Absorbers 7.2. Sintered Nanostructured Solar Cells Based on NCs 7.2.1. Outlook of Sintered Solar Cells 7.3. Nanocrystal Thin Film Solar Cells 7.3.1. Cu Chalcogenides in NCSCs 7.3.2. Layer Processing 7.3.3. Device Architecture 7.3.4. Ligands 7.3.5. Limiting Device Parameters 7.3.6. Single-NC Devices and Cu Chalcogenide Arrays 7.3.7. Combinations with Other Materials 7.3.8. Outlook of NCSCs 7.4. Semiconductor Sensitized Solar Cells 7.4.1. Composition 7.4.2. Deposition Process 7.4.3. Multilayer/Barrier Layer 7.4.4. Photocathode 7.5. Counter Electrodes in Photoelectrochemical Cells

Review

7.5.1. 7.5.2. 7.5.3. 7.5.4. 7.5.5.

Composites and Nanostructured Arrays Ligand Removal Processing Influence of Stoichiometry Combination with Other Materials and Architectures 7.5.6. Other Cu Chalcogenides 7.5.7. Applications beyond SSSCs 7.6. Hybrid Organic−Inorganic Solar Cells 7.7. Third Generation Concepts 7.7.1. Multiple-Exciton Generation Solar Cells 7.7.2. Intermediate Band Solar Cell 7.7.3. Tandem Solar Cells 7.8. Optical Enhancement 7.8.1. QD Luminescent Concentrator 7.8.2. Down-Converter 7.8.3. Plasmonic Enhancement of Light Absorption 8. Lighting/Displays 8.1. Electroluminescent QD-LED 8.2. Down Conversion QD-LED 9. Catalytic Applications 9.1. Photocatalysis 9.1.1. Water Splitting 9.1.2. Photocatalytic Degradation of Pollutants 9.1.3. CO2 Photoreduction 9.1.4. Other Photocatalytic Applications 9.2. Other Catalytic Applications 9.2.1. Oxygen Evolution and Oxygen Reduction Reactions (OER/ORR) 9.2.2. Other Catalytic Reactions 10. Energy Storage Applications 10.1. Batteries 10.1.1. Cu−S 10.1.2. Cu−Se 10.1.3. Cu−Te 10.1.4. Cu−Sn−S 10.1.5. Cu-Zn-Sn-S 10.1.6. Cu-In-Zn-S 10.1.7. Cu-Bi-S 10.1.8. Cu-Sb-S 10.1.9. Cu-Mo-S 10.1.10. Cu-Fe-S 10.2. Supercapacitors 11. Thermoelectric Applications 11.1. Quaternary Copper Chalcogenides 11.2. Ternary Copper Chalcogenides 11.3. Binary Copper Chalcogenides 12. Sensors 12.1. Fluorescence-Based Sensors 12.2. Chemiluminescence-Based Sensors 12.3. Electrochemical-Based Sensors 12.4. Other Sensing Methods 13. Bioapplications 13.1. Photothermal Therapy 13.2. Photodynamic Therapy/Photochemotherapy 13.3. Chemotherapy/Drug Delivery 13.4. Immunotherapy 13.5. Radiotherapy 13.6. Photoacoustic Imaging 13.7. Fluorescence Imaging

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Chemical Reviews 13.7.1. Quantum Dot Fluorescence 13.7.2. Up-Conversion Luminescence (UCL) 13.8. Dark-Field Microscopic Imaging 13.9. Ultrasound Imaging 13.10. Magnetic Resonance Imaging 13.11. X-ray Computed Tomography (CT) 13.12. Positron Emission Tomography 13.13. Single-Photon Emission Computed Tomography 13.14. Toxicity Studies 14. Summary, Challenges, and Outlook 14.1. Synthesis 14.2. Heterostructures/Material Doping 14.3. In Situ Monitoring of NC Growth 14.4. Upscaled Synthesis Approaches 14.5. Composition Control 14.6. Characterization 14.7. Unexplored Compositions 14.8. Surface Engineering 14.9. Challenges in Assembly 14.10. Challenges in Applications Author Information Corresponding Authors ORCID Notes Biographies Acknowledgments List of Abbreviations References

Review

understanding and classification of the metal ions and ligands from molecular chemistry, which allows for the design of experiments to optimize their compatibility during the nucleation and growth process. This has allowed for a level of synthetic control to be attained, whereby the NC size, shape, crystal structure, and composition can be accurately defined. The synthetic strategies are as extensive and diverse as the compositions attainable, comprising bottom-up protocols such as hydrothermal, solvothermal, and colloidal hot-injection approaches to postsynthetic modifications such as cation exchange. A commonality in all approaches, that involve NC nucleation and growth in solution, is the requirement of coordinating reagents (i.e. surfactants), where subtle differences in their binding behavior to the facets of a growing crystal permit size and shape selectivity. The catalogue of potential morphologies is unlimited, in that it is possible to control solid shapes from 0D to 3D, in addition to porosity in the form of tubes. The synthetic parameters to define each of these morphologies is a very finely balanced combination of control factors, from the choice of precursors and their decomposition kinetics, the reaction temperature, the ligand functionality and combination, in addition to the choice of coordinating or noncoordinating solvents. The resultant colloidal dispersions enable wet chemical processing, where the collective properties of these NCs can be harnessed for game-changing innovations that impact a vast range of technological applications. Secondary shells can be grown on Cu chalcogenide NCs to passivate the surface states, thus resulting in core−shell quantum dots (QDs) with enhanced optical properties. Through the use of deposition protocols directly from the colloidal NC dispersions, it is also possible to exert complete control of the position and orientation of the nanoscale building blocks, to form superstructures of highly organized NCs. This ultimately allows for bulk properties that are not just defined by the composition, size, and shape of the functional unit, but also by the order attainable in the final NC assembly. Advances in NC assembly by directed electric field protocols have reached the point that NC assemblies can be grown layer by layer, directly from a substrate, thus allowing for the rational integration of Cu chalcogenide NC assemblies into functional devices. There is an interesting evolution in properties and applications as the number of elements increases, from the simplest binary elemental composition, of Cu and a chalcogen (i.e. S, Se, Te), to the more complex multinary compositions. In transition metal Cu chalcogenides, there is a strong mixing of the chalcogen valence s and p orbitals with the outer s- and p-orbitals from the metals, which results in more covalent rather than ionic character. In addition to this, the low size and electronegativity differences between the metals and chalcogens, the ability to form chalcogen−chalcogen bonds, and the possibility of metal−metal bonds within the structure, allows for a large diversity in stoichiometry and crystal structures and the resultant functional properties. The high defect concentration influences their charge and heat transport properties, which greatly affects their electronic, thermoelectric, and optoelectronic properties. For example, the binary Cu chalcogenides have been long investigated for their interesting plasmonic properties, which allow them to be used as novel probes for surface-enhanced Raman spectroscopy (SERS) or hyperthermia. Copper sulfide (Cu2−xS), copper selenide (Cu2−xSe), and copper telluride (Cu2−xTe) are regarded as self-doped materials and can support localized surface plasmon resonances (LSPRs) in the near-infrared (NIR) region, due to Cu

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1. INTRODUCTION 1.1. General Overview

Copper chalcogenide nanocrystals (NCs) and their multinary derivatives, encompassing additional transition metal (Zn, Fe, Cr, Zn, Cd, Hg), main-group metal (In, Ga, Sn, Bi, Sb), and semi-metal (Ge) ions, are an extraordinarily interesting class of materials with a wide range of stable crystal phases attainable. The expansive developments in this field represent a paradigm of materials engineering, as the entire gamut of physical, electronic, optical, and magnetic properties are controllable by the elemental composition and the stoichiometry of their distribution. The high level of interest in copper (Cu) chalcogenide NCs resides in three key aspects: (i) their abundance, low cost, and reduced environmental and health impact, compared with cadmium- and lead-based compounds; (ii) their excellent intrinsic functional properties, including appropriate direct band gaps for solar light absorption, plasmonic properties, notable charge carrier mobilities, potential high carrier concentrations, and low thermal conductivity; and (iii) their structural, compositional, and stoichiometric versatility, including abundant non-stoichiometric phases, a wide range of solid solutions, and the related low energy of formation of defects. Even the simplest binary compounds, Cu2−xA (where A = S, Se, Te), present over 20 reported binary stoichiometries, polymorphs, and defect phases, and several dozens of ternary and multinary compounds have also been reported. The extensive catalogue of precursor metal salts and chalcogen sources, combined with a relatively low nucleation temperature in solution, makes their stabilization in nanocrystal (NC) form particularly amenable. The most successful synthetic protocols have been derived from a fundamental 5867


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photocatalytic (PC) or photoelectrochemical (PEC) splitting of water, or for the degradation of pollutants and carbon dioxide via photocatalysis. Their catalytic activity in the absence of light has also been demonstrated, in that they can serve as electrocatalysts for oxygen evolution and oxygen reduction reactions (OER/ORR), which are promising for lithium−air reactions. The optoelectronic properties can also be tuned for emission with suitable surface capping, with applications of multinary Cu chalcogenide QDs demonstrated as light emitting diodes (LEDs), as electroluminescent devices, and as down-converters for displays and solid-state lighting. The ubiquity of their application has also been extended to sensors, thermoelectrics, electrodes for lithium-ion batteries, and supercapacitors, where nanoscale sizes, high surface areas, tunability of NC shape and crystal phase, and the different electrode potentials of the metals and chalcogens allow for designer material characteristics. In this comprehensive review, we give a detailed account of the synthetic strategies and protocols available for the formation of multinary Cu chalcogenide NCs, collating all experimental conditions and interpreting the critical aspects of these protocols, which have allowed for progression to the point where NC size, shape, and crystal phase control are routine. This treatise is structured in such a way that the evolution in complexity is explained, particularly as the compositions progress from binary to ternary and higher order forms, so that the rationale for the choice of key reaction conditions and parameters is apparent to attain the desired structures. The assembly pathways (both self- and directed-protocols) and the intrinsic and controllable properties, that can be used to manipulate assembly in this system, are extensively detailed within this review. The hierarchical organization of NCs serves as a link to a full discourse on the wide application set for these materials, including photovoltaics, optoelectronics, thermoelectrics, photocatalysis, fluorescent biological imaging, and photothermal therapy. This complete review builds on more focused reviews1−23 and on particular aspects of these materials in the last decade. The essence of this review is to capture the confluence of advances that have emerged in the synthesis, properties, assembly, and application of multinary Cu chalcogenide NCs in a single resource, to serve as a benchmark and foundation for the next generation of advances.

vacancies in the material. This composition-dependent property is particularly attractive for biological applications because the LSPR is spectrally located at a suitable window for applications such as photothermal therapy and biomedical imaging. The reversible tunability of the LSPR is one of the strongest advantages of Cu chalcogenide semiconductor NCs, compared to noble metals, which are restricted to a fixed carrier density. This intrinsic doping strategy to control the semiconductor electronic properties is especially interesting in the bottom-up processing of NCs, where the introduction of extrinsic dopants still represents a major challenge. The control of plasmonic properties is the focus of extensive efforts, where the possibility to manipulate the plasmonic frequency by changing the crystal phase, composition, and geometry of the NCs is very attractive. While the most common Cu chalcogenide phases have been widely investigated for plasmonic applications, less studied systems such as Cu3BiS3 are recently finding application in synergistic therapies, where a combined near-infrared (NIR) and magnetic resonance imaging (MRI) response allows for detailed and exact information on tumors to be extracted, at a very early stage of development. The introduction of p-block metals/semi metals (Ga, In, Sn, Ge) or transition metals (Zn, Fe, Cd) to form the ternary and quaternary Cu chalcogenide compositions adds a high degree of complexity in the NC synthesis, but it further widens their range of technological applications. The numerous possibilities for chemical substitutions and structural modifications in the ternary and quaternary compositions provide significant scope to tune the fundamental material properties for the desired end application. For example, copper indium disulfide (CuInS2, CIS), copper indium diselenide (CuInSe2, CISe), and copper indium gallium diselenide (CuInGaSe2, CIGSe, known as chalcopyrite) are important compound semiconductor representatives, which have been studied for more than two decades as p-type absorber materials in thin film photovoltaics. These materials are renowned for their high energy conversion efficiencies, high optical absorption coefficients, good photostability under long term radiation, and relatively low toxicity, in comparison to the conventional cadmium-based materials. A material currently of specific interest is copper zinc tin sulfide (Cu2ZnSnS4, CZTS), as all the elements are in high natural abundance, which makes this material very attractive for sustainable large-scale roll-out of photovoltaic devices. Notably, the best efficiencies for this material to-date have been achieved by solution-processed ink-based technologies, which offer a lower energy production pathway than the conventional vacuum-based fabrication approaches. As the NC composition can be tuned synthetically, this allows for band-gap tuning either as a function of the metal (cation) ratios or the chalcogen (anion) ratios, if sulfur and selenium are interchanged. The versatility of Cu chalcogenide NCs has allowed for their use, not just in thin film photovoltaics, but also in semiconductor sensitized solar cells (SSSCs), hybrid organic-inorganic cells (with conducting polymer blends), and third generation multiple-exciton cells, where external quantum efficiencies greater than 100% are possible. In addition, the ability to tune the majority carrier concentration and energy levels makes multinary Cu chalcogenide compounds particularly attractive, as they can function either as tandem absorbers for multi-juntion cells or as counter electrodes for electron and hole extraction in photoelectochemical and perovskite solar cell devices. These multinary Cu chalcogenide NCs can also be used to harvest solar energy for the direct conversion of chemicals in the

1.2. Outline of Review

The review is divided up into 14 different sections to provide the reader with a comprehensive review of Cu chalcogenide semiconductor NCs, with details on their crystal phases and stoichiometries (section 2), functional properties, such as electronic properties, photoluminescence, plasmonics, nonlinear optics, and magnetism (section 3), solution synthesis approaches (section 4), focused synthesis protocols for each material (section 5), and assembly strategies (section 6). The latter half of the review covers the broad range of applications of Cu chalcogenide NCs (sections 7−13), specifically photovoltaics (section 7), lighting/displays (section 8), catalysis (section 9), energy storage (section 10), thermoelectrics (section 11), sensors (section 12), and bio-applications (section 13). A summary of this review, with details on the challenges and future outlook in the field of Cu chalcogenide NCs, is provided in the final section (section 14). Figure 1 shows the wide variety of Cu chalcogenide NCs that are covered in this review, spanning from binary compositions such as Cu2−xS, Cu2−xSe, and Cu2−xTe to ternary compositions that contain an additional group III (In, Ga), group IV (Sn, Ge), 5868


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Figure 1. Schematic figure illustrating the range of Cu chalcogenide semiconductor NCs that are covered in this review.

The large diversity in the transition metal chalcogenide stoichiometries and structures, including abundant nonstoichiometric phases and a wide range of solid solutions, is due to a series of factors such as (i) the low size and electronegativity differences between transition metals and chalcogens; (ii) the fact that the formal chalcogen oxidation state is not necessarily 2, as in oxygen; (iii) the chalcogen ability to form chalcogen−chalcogen bonds; (iv) the variable valences of transition metals; and (v) the higher ability of metal atoms to form metal−metal bonds within the transition metal chalcogenide compounds.

or group V (Sb, Bi) atom to form Cu-III-VI, Cu-IV-VI, and Cu-V-VI compounds (where VI = S, Se, Te), respectively, right through to the quaternary compositions that are typically comprised of either two group III atoms (In, Ga) or one group II (Zn, Cd, Hg) and one group IV (Sn, Ge) atom to form Cu-III-VI (e.g. CIGS, CIGSe) and Cu-II-IV-VI (e.g. CZTS, CZTSe) compounds, respectively. Other quaternary compositions have also been obtained via the replacement of Zn or Sn with Co, Mn, Ni, and Fe. It is also possible to form NCs that are not only restricted to one chalcogen compound, but contain two chalcogens, and this forms NC compositions such as Cu2−x(SySe1−y), CuIn(SySe1−y)2, CuInGa(SySe1−y)2, and Cu2ZnSn(SySe1−y)4, which further permit band-gap tuning by variation of the (SySe1−y) ratio.

2.1. Binary Compounds

Cu chalcogenide compounds are a rich family of materials. In particular, the simplest binary compounds (Cu2−xA, where A = S, Se, Te) present several different stoichiometries and polymorphs, which exhibit numerous defect phases and quite complex atomic arrangements. These binary compounds share a common feature in their high temperature phases, in that the chalcogen atoms form regular lattices such as face-centered cubic (fcc) or hexagonal close-packing (hcp) and the Cu atoms are in disorder, occupying various interstitial sites in the chalcogen lattice. All these compounds show a wide deviation from stoichiometry, as expressed by the formula Cu2−xS, Cu2−xSe, and Cu2−xTe, and have the characteristic features of a p-type semiconductor because of Cu vacancies within the lattice. However, the experimental determination of the stoichiometric crystal structures of Cu2S, Cu2Se, and Cu2Te is limited by the commonly observed Cu deficiency, which typically makes the synthesized samples non-stoichiometric and further complicates the crystal structure determination.27 2.1.1. Cu−S System. Copper sulfide (Cu2−xS) is probably the most complicated binary compound, in terms of evaluating its crystal phase, and the nomenclature used is somewhat confusing. From an applications point of view, it is critical to finely tune the crystal phase because the properties of the Cu2−xS compounds are sensitive to both the Cu−S stoichiometry and the crystal structure. The binary Cu−S system exhibits a rich phase diagram28 and is known for its diversity of crystal phases, which are comprised of numerous nonstoichiometric compositions over a wide range. There are eight dominant crystal phases (shown in Figure 2a−h) in the Cu−S

2. CRYSTAL PHASES AND STOICHIOMETRIES An important distinction between transition metal chalcogenides and oxides is that transition metal chalcogenides have more covalent character, whereas transition metal oxides are more ionic in character and resemble fluorides more than sulfides. Transition metal chalcogenides possess pronounced covalent character because of the low electronegativity differences between transition metals and chalcogens, which leads to a strong mixing of the s and p orbitals of the chalcogen with the outer s and p orbitals of the metal. This mixing results in broad valence and conduction bands and narrower energy gaps because the valence band of chalcogenides lies at higher energy than in oxides.24−26 The differences between oxides and chalcogenides are reflected in their strikingly different crystal structures, in their chemical and physical properties, and markedly in their functional properties. In contrast to the abrupt dissimilarities between oxides and chalcogenides, related to the electronegativity jump from oxygen (3.5) to sulfur (2.5), selenium (2.4), and tellurium (2.1), the structure of transition metal sulfides generally resembles that of selenides and tellurides, compared to that of oxides. In terms of their physical properties, a trend toward more covalent bonding, more delocalization of electrons, and increasing metallic behavior is observed in the chalcogenide compounds, as we move down through the chalcogen group. 5869


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Figure 2. Crystal structures for (a) low-chalcocite, (b) high-chalcocite, (c) cubic-chalcocite, (d) djurleite, (e) digenite, (f) anilite, (g) roxbyite, and (h) covellite. Blue and yellow spheres are copper and sulfur atoms, respectively. (i) Table of the different mineral names and their corresponding compositions, sulfur (S) packing, and crystal system. (j) Calculated heat of formation of CuxS per formula unit as a function of x. Panels (a−f) are reprinted with permission from ref 30. Copyright 2012 AIP Publishing. Panel (g) is reprinted with permission from ref 31. Copyright 2016 American Chemical Society. Panel (h) is reprinted with permission from ref 32. Copyright 2014 American Chemical Society. Panel (j) is reprinted with permission from ref 30. Copyright 2012 AIP Publishing.

2.1.1.1. Instability of Cu2S. An important point to address in Cu2S is its instability. Previous experimental and theoretical investigations have found that Cu2S is intrinsically unstable at ambient conditions and rapidly degrades into copper-deficient Cu2−xS phases, most notably djurleite, which has been determined to be the thermodynamically stable stoichiometry under ambient conditions.6,33,38,39 Djurleite is a slightly copperdeficient phase, which exists at Cu/S ratios between 1.94 and 1.97,40 and has a lower free energy than Cu2S at room temperature, owing to its lower crystallographic symmetry; thus, djurleite is more stable in the bulk36,41 and in NCs.42 In addition, it was deduced that Cu2S NCs without Cu defects were nearly impossible to synthesize or even store long term.43 This is because Cu2S (high chalcocite) has a thermodynamic propensity toward a Cu deficiency, owing to the low chemical potential of Cu0.33,41 A review of the literature on Cu2S NCs raises an important fundamental question regarding the correct crystal structure assignment of as-synthesized NCs with the standard JCPDS database files. While chalcocite (Cu2S) and djurleite (Cu1.94S) have distinguishable electronic and crystal structures, they have very similar XRD patterns at first glance. However, Burda and co-workers provided important insights into assigning the correct phase and determined that, after careful study, the XRD patterns are also distinguishable.42 For example, the XRD pattern of Cu1.94S (JCPDS 023-0959) can be distinguished from Cu2S (JCPDS 033-0490) because Cu1.94S has very narrow double peaks at 37.6°, 46.3°, and 48.6° 2θ, whereas Cu2S has strong peaks at 45.9° and 48.5° 2θ but it does not have a peak at 46.3° 2θ. Lotfipour et al. also pointed out that although most

system, which are distinguished from each other depending on the Cu content: low-chalcocite (monoclinic Cu2.0-1.997S), high-chalcocite (hexagonal Cu2.0-1.94S at T > 103.5 °C), cubicchalcocite (Cu2S at T > 436 °C), djurleite (Cu1.97S−Cu1.94S), digenite (Cu9S5 or Cu1.8S), anilite (Cu7S4 or Cu1.75S), roxbyite (Cu58S32 or Cu1.81S), and covellite (CuS).28 These crystal structures are characterized by either hcp or fcc (also known as cubic close-packing) of sulfur (S) atoms, with Cu atoms occupying various interstitial sites.29 The nature of packing of S atoms to the corresponding mineral name and composition is shown in the table in Figure 2i. The complexity of these crystal structures has led to a longstanding and challenging problem in elucidating the structure of the various mineral forms of Cu2−xS, principally because the positions of the Cu atoms within the close-packed sublattice of S atoms are not well-defined.33 At elevated temperatures, the Cu atoms are unusually mobile, thus making Cu2S a partially ionic conductor.34,35 Moreover, transformations involving the arrangement of S atoms from cubic to hexagonal packing (and vice versa) are extremely slow (particularly at low temperatures), which leads to a number of metastable Cu2−xS phases,28 all of which exhibit p-type electronic character as a result of Cu vacancy doping. These vacancies can form by the loss of Cu to either oxygen or carbon dioxide at a free surface or at a grain boundary.36,37 The vacancy (i.e. hole) concentration in Cu2−xS is generally understood to be stoichiometrydependent and increases with the number of Cu vacancies (x), with reported carrier densities typically on the order of 1021 cm−3.37 5870


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of the small peaks in the Cu1.94S pattern are coincident with peaks in the Cu2S pattern, the peaks at 26.2° and 55.75° 2θ are indicative of the formation of Cu1.94S.43 The XRD pattern of Cu1.94S is also similar to Cu1.8S (JCPDS 047-1748), but the key differences between these phases are the peaks at ∼37.6° and 48.6° 2θ, which are present in Cu1.94S but are absent in Cu1.8S. Despite some inconsistencies in the literature between the reported phase assignments of Cu2−xS NCs, it is interesting to note that NCs that were previously assigned as highchalcocite (Cu2S) actually have a better match with the djurleite phase (Cu1.94-1.97S).42 The first reliably assigned Cu2S NCs were reported by the Alivisatos group,44 based on NCs that were synthesized and stored under inert gas, but the XRD pattern of perfect Cu2S NCs quickly transformed to Cu1.97S once exposed to air. Through a high resolution TEM study, Alivisatos and co-workers observed the dynamic structural transformation of monoclinic Cu2S (low-chalcocite) into the hexagonal Cu2S (high-chalcocite) in Cu2S nanorods.45 In this report, irradiation from the electron beam was used to induce the transformation of low-chalcocite to high-chalcocite, where the latter structure formed at temperatures between 104 and 436 °C. Temperatures higher than 436 °C would have been required to visualize the transformation of high-chalcocite into cubic-chalcocite. Moreover, it was also determined that the phase transformation temperature of Cu2S NCs was dependent on the particle size and that particles >6 nm are required to maintain the low-chalcocite structure at room temperature.46 2.1.1.2. Relationship between Cu2−xS Composition and Crystal Structure. While the names for the different crystal phases are not very meaningful, it is noteworthy to highlight the different crystal structures possible in the Cu−S system, given that there is a great deal of disorder in all these compositions. The crystal structure of low-chalcocite was fully determined to be monoclinic (P21/c space group) by Evans.47 It consists of a distorted hexagonal framework of 12 S atoms with 24 Cu atoms occupying mainly triangular interstices, where the 36 crystallographically distinct Wyckoff sites are fully occupied. Of the 24 different Cu atoms, 21 form triangular CuS3 groups and one is in a distorted CuS4 tetrahedron.48 In each unit cell, there are a total of 48 S and 96 Cu atoms (forming a 144-atom unit cell, Cu96S48) because every Wyckoff site represents 4 equiv atomic positions.30 Low-chalcocite is closely related to high-chalcocite, in that it is essentially a superstructure of the hexagonal lattice but the Cu arrangements in the two phases are entirely different and the Wyckoff sites in the high-chalcocite and cubic-chalcocite structures are not fully occupied. The crystal structure of lowchalcocite was initially assigned as orthorhombic because a prevalent twin law resulted in nearly perfect superposition of monoclinic lattices to produce an apparent end-centered orthorhombic lattice.49 For the high-chalcocite structure, the unusually mobile nature of the Cu atoms makes it difficult to determine their exact positions and so, only a statistical distribution of Cu atoms over the various Wyckoff sites in the hexagonal lattice can be determined. This mobility is possible because the Cu atoms can assume tetrahedral, trigonal, and linear coordination.50 It is interesting to note that the disorder in the high-chalcocite phase leads to reduced mobility for holes and electrons, in what amounts to the higher symmetry phase.51 However, different authors report somewhat different distributions of Cu atoms, which are based on the best fit of the generated XRD data from the corresponding relative occupancy of the Wyckoff sites.33

Cu1.94-1.97S (djurleite) is similar to low-chalcocite, in that it has a monoclinic structure with a slightly reduced Cu concentration. Both low-chalcocite and djurleite are very common species and are often intermixed or intergrown.48,52 The existence of the djurleite phase was initially discovered in an X-ray study of the Cu−S system by Djurle53 and later became established as the mineral djurleite in his honor.54 The unit cell content of djurleite corresponds to Cu62S32 or Cu1.94S, and of the 62 Cu atoms in the structure, 52 are in a threefold, triangular coordination with S, 9 are in a tetrahedral coordination, and 1 is in a linear coordination.52 If an extra Cu atom is added, the unit cell content becomes Cu63S32 or Cu1.97S, thus explaining the lower and higher limit of Cu1.94-1.97S for djurleite. Cu1.8S (digenite) is a separate phase and can exist in cubic, rhombohedral, and hexagonal modifications depending on the temperature.33,55 In the early literature, digenite is sometimes referred to as Cu9S5, but since the Cu atoms are statistically distributed over several sites, the formula Cu1.8S is preferred when referring to digenite. It should be noted that djurleite and digenite undergo a phase transformation to structures with higher symmetry, in the presence of high temperatures51 or under electron beam exposure.56 Cu7S4 or Cu1.75S (anilite) has an orthorhombic crystal system (Pnma space group), in which the S atoms approximate the fcc arrangement and the Cu atoms are ordered in the interstices.57 Of the 28 Cu atoms in the orthorhombic unit cell, 20 are in triangular coordination and 8 are in tetrahedral coordination.49 The crystal structure of roxbyite (Cu58S32 or Cu1.81S) is just as complex in its entirety as the crystal structures of lowchalcocite and djurleite.58,59 Mumme et al. described roxbyite as having a triclinic unit cell (90 atoms), which consists of a distorted hcp framework of 32 S atoms with 58 independent Cu atoms occupying these layers, all having triangular coordination.59 Other layers sandwiched between the closepacked S layers consist purely of double or split layers of Cu atoms, some of which have linear coordination to the S atoms, but mostly they have triangular or tetrahedral coordination. Roxbyite is uncommon in bulk syntheses but can be readily synthesized in NCs.31,60−63 An interesting solid−solid phase transformation of roxbyite (Cu1.81S) phase NCs into a more thermodynamically stable, copper-rich phase (djurleite (Cu1.94S)/ low-chalcocite (Cu2S)) was recently reported, which occurred as a result of Cu ions (that were displaced by Zn ions) occupying the intrinsic Cu vacancy sites in the roxbyite phase.60 However, the minimization of strain energy induced a second phase transition back to roxbyite, even though djurleite (Cu1.94S)/low-chalcocite (Cu2S) was predicted to be more thermodynamically stable than roxbyite from density functional theory (DFT) calculations. CuS (covellite) has a peculiar layered structure for such a simple stoichiometry, in that its Cu atoms are in two different environments: CuS3 units (triangular planes) are positioned between two layers of CuS4 units (tetrahedral) along the c-axis, with the trilayers stitched together by S−S covalent bonds.32 Another description of CuS in the literature is that it consists of alternating CuS3−CuS3−CuS3 layers and S−S layers along the c-axis.64 It is interesting to note that CuS forms this unusual structure, rather than one of the simpler available structures, because it is more energetically favorable for the S atoms to covalently bond to each other, than to bond with Cu atoms.65 Of the 6 formula units (12 atoms) in the unit cell, 4 of the Cu atoms have tetrahedral and 2 have triangular coordination, whereas 4 of the S atoms form disulfide S2 groups and 2 are 5871


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single sulfide ions.66 Covellite CuS has a unique metallic-like character accounting for its elevated, anisotropic p-type conductivity.67−69 The peculiar crystallographic feature in CuS allows for the intercalation and reaction of alkali metal ions (i.e. Li) between the S−Cu−S layers, which makes CuS NCs promising electrode materials for Li-ion battery applications.70 Our discussion on the crystal systems in Cu−S would not be complete without mentioning other less researched phases such as Cu1.6S (geerite), Cu1.39S (spionkopite), and Cu1.12S (yarrowite) and CuS2 (villamanitite), which are possible in the bulk but are somewhat uncommon in NCs. While the optoelectronic properties are known to greatly depend on the crystal phase, there is a lack of robust experimental data for Cu2−xS because it is nontrivial to identify the low symmetry phases and to measure the exact stoichiometry, which varies between x = 0 and 0.2 (i.e. Cu2S−Cu1.8S). Xu et al. employed first-principles methods to systematically search for the most energetically favorable Cu2−xS structure and to understand the general phase stability (Figure 2j).30 Of the three chalcocite phases, Cu2S is more stable in the lowchalcocite structure than in high-chalcocite and cubic-chalcocite structures. However, as the Cu concentration increases, the formation energy also increases, which destabilizes the structure and indicates that Cu2S is not stable against the formation of Cu vacancies. The calculations suggested that anilite (Cu1.75S) is one of the most stable forms of copper sulfide under Cu-rich conditions because it has the lowest heat of formation.30 To cover the complete picture, Robinson and co-workers calculated the phase stability for the two phases missing in Xu et al.’s study (i.e. Cu1.81S (roxbyite) and CuS (covellite)) and revealed that CuS has an even lower formation energy than anilite.60 However, it is difficult to extract further conclusions given the limited degree of phase stability data; thus, further studies are warranted in this area. 2.1.1.3. Oxidation State of Cu. By far, one of the biggest unexpected discoveries in Cu2−xS is that the oxidation state of Cu does not change and remains close to Cu+, regardless of the stoichiometry. It had been assumed that Cu was changing oxidation states, and in CuS (covellite), in particular, the valence of Cu is still highly debated. The differentiation between Cu+ and Cu2+ has important consequences if these materials are used as magnetic resonance imaging (MRI) contrast agents because Cu+ is diamagnetic, meaning that it will be repelled by an externally applied magnetic field. Some studies have found Cu(I) and suggested a (Cu+)3(S2−)(S2−) valency formalism69 or (Cu+)3(S22−)(S−).71 Other studies have indicated the presence of both Cu(I) and Cu(II), proposing a (Cu2+)(S22−)(Cu+)2(S2−) formulation,72,73 and recent calculations have set the valency of Cu in CuS as 1.33.74 However, recent measurements have shed some new insights on this topic. X-ray photolectron spectroscopy (XPS) and electron paramagnetic resonance (EPR) measurements have provided strong evidence that the valency of Cu remains close to +1 in Cu-rich covellite (Cu1.1S) and in all NCs with stoichiometry close to Cu2S, and that the sulfur (S) valency changes from −1 in Cu1.1S to −2 in Cu2S; i.e. sulfur is progressively reduced.75 This finding is in agreement with previous XPS studies on bulk Cu2−xS.76−78 Manna and coworkers highlighted that, in Cu1.1S, the incorportion of Cu(I) species resulted in S−S covalent bonds being progressively broken to make room for the intercalated metals; thus, the valency of the S anion is strongly affected.75 The intercalation of metals in Cu1.1S is sustained by a change in the redox state of

the anion framework. Specifically, the S anion framework is reduced and this is paralleled by an oxidation process in solution, where a fraction of the Cu(I) species that remained in solution were oxidized to Cu(II). This data was also corroborated with EPR measurements and SQUID (superconducting quantum interference device), which can measure extremely subtle magnetic fields and, thus, can be used to differentiate Cu+ (diamagnetic) from Cu2+ (paramagnetic). Previous studies on the oxidation of bulk Cu2−xS in water have also shown that the process involves the oxidation of surface Cu+ to Cu2+.79 In terms of Cu2−xS NCs, it was proven by EPR spectroscopy that the reaction of the NCs with a Cu(I) complex and the resultant increase in the Cu stoichiometry in the NCs were matched by the formation of an equimolar amount of Cu(II) species in solution.75 Thus, EPR spectroscopy has proven to be a powerful technique in terms of elucidating the valency of Cu, where the presence of a signal indicates the presence of Cu2+ and no signal (i.e. EPR silent) indicates that Cu exists as Cu+. A similar observation was noted in a separate study, where the presence of an EPR signal in chemically etched Cu2−xS NCs definitively established that the removed Cu from the NCs was oxidized to Cu2+ in solution and that the chalcogens were removed in the form of phosphine chalcogenides.80 In this particular report, it was concluded that alkylphosphines chemically etched the Cu2−xS NCs in the presence of oxygen, as the analysis of a similar sample, which was kept in an N2 glovebox instead of in the presence of oxygen, showed no EPR signal. The absence of an EPR signal is consistent with earlier studies, which confirmed that Cu+ is the only oxidation state present in Cu2−xS, regardless of stoichiometry.75,81,82 The stoichiometric variations in Cu2−xS can be correlated to changes in the underlying electronic structure, particularly where the valency of the cations and anions in the lattice is concerned. In Cu chalcogenides, the top of the valence band has a strong contribution from the chalcogenide 3p orbitals and the bottom of the conduction band mainly has contributions from the Cu 4s and 4p orbitals (in the case of Cu2−xSe).15,75 It is assumed that each Cu atom contributes to bonding with one 4s electron and each chalcogen atom contributes with six p electrons.15 In the fully stoichiometric Cu2S material, the valence band is completely filled and the material would behave as an intrinsic semiconductor. However, when Cu vacancies are created (e.g. from oxidation), this leaves a hole in the top of the valence band. This hole creation mainly affects the valency of the chalcogen (S) because the valence band has a major contribution from the chalcogen, whereas the valency of Cu remains close to +1. 2.1.2. Cu−Se System. Copper selenide (Cu2−xSe) has a similar tendency to crystallize in a large number of phases, with a significant Cu deficiency allowed in the chemical stoichiometry of Cu2−xSe. It is a self-doped p-type semiconductor, which is dominated by Cu vacancies (holes), and this makes the carrier concentration and transport properties strongly dependent on composition. The binary Cu−Se phase diagram was first reported by Heyding83 and is comprised of various compounds and structural forms, ranging from cubic berzelianite (Cu2Se, Cu1.8Se), to tetragonal umangite (Cu3Se2), hexagonal klockmannite (CuSe, Cu0.87Se), orthorhombic athabascaite (Cu5Se4), and orthorhombic marcasite (CuSe2).84−88 All these compounds exhibit near metallic conduction at room temperature.89 However, the thermal stability of these compounds is highly dependent on the stoichiometric composition. 5872


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arrangement and each Se atom is bonded to three Cu neighbors and one Se neighbor in a distorted tetrahedral configuration.84,126,127 The difference between the structures is related to the Se octahedra, where they are connected by shared corners in pyrite and by shared edges to form chains in the marcasite structure. XPS measurements have determined that Cu is monovalent (+1) in many of these selenides, whereas the average oxidation state of the chalcogen (Se) varies from −2 in Cu2Se to −3/2 in Cu3Se2, and from −1 in CuSe to −1/2 in CuSe2.76 However, similar to CuS, there are some discrepancies in the literature regarding the oxidation states of Cu and Se in Cu2−xSe. While the oxidation of Cu and Se surface sites has little impact on the electronic properties of a bulk material, these effects become significant on the nanoscale and can drastically alter the electronic properties, as evidenced by Riha et al. in their work on Cu2−xSe NCs.128 In this work, changes in both the Cu and Se oxidation states were detected in the XPS analysis of Cu2−xSe NCs upon air exposure. The initial Cu XPS spectrum detected the presence of Cu+ in Cu2Se NCs, but a dramatic change in the spectra occurred after the oxidation of Cu2Se to Cu1.8Se, with peaks attributed to mixed valent Cu+ and Cu2+ present in the oxidized Cu1.8Se NC sample. The Se XPS spectra was also evaluated, where Se2− was initially the only species present in Cu2Se, but upon oxidation, it was converted to an intermediate species with an oxidation state more positive than Se2− and a smaller percentage being converted to Se4+, tentatively identified as SeO2. The oxidation of Cu+ and Se2− ultimately led to a solid-state conversion of the NC core from monoclinic Cu2Se to cubic Cu1.8Se, which resulted in a significant enhancement in the electronic properties and a change from semiconducting to ohmic behavior, respectively. Some studies on the magnetic properties of Cu2−xSe compounds in the bulk have shown that Cu3Se2 exhibits the largest susceptibility to paramagnetism, consistent with the presence of d9 ions, and may be anti-ferromagnetic at low temperature.89 Paramagnetic susceptibility has also been observed in phases with lower Cu content (i.e. CuSe and CuSe2), whereas phases with higher Cu content (Cu2Se, Cu1.8Se) are deemed to be diagmagnetic or show weak temperature independent paramagnetism.89,129 Moreover, CuSe2 has been reported to be a superconductor at low temperatures (2.4 K) and is weakly ferromagnetic below 31 K, implying the possible co-existence of ferromagnetism and superconductivity in this compound.130−132 2.1.3. Cu−Te System. Copper telluride (Cu2−xTe) can exist in a wide range of compositions and phases. The binary Cu−Te phase diagram is the most complex among the Cu chalcogenide systems, with the most accurate phase diagram proposed by Pashinkin and Fedorov.133 It is comprised of different crystal phases and structures depending on the relative stoichiometry, such as Cu2Te (weissite, hexagonal Nowotny’s model134), Cu7Te4 (trigonal), Cu4Te3, Cu3Te2 (rickardite, tetragonal), CuTe (vulcanite, orthorhombic), and CuTe2 (pyrite-type compound). All of these phases in the Cu−Te system consist of a rigid Te framework and mobile Cu ions in different valence states, Cu(I) or Cu(II), with the arrangement of the Cu ions depending on the composition.133 The phase sequence in Cu2Te is more complicated than in Cu2Se because it has five successive phase transitions between room temperature and 900 K.135−137 The phases are described as α (< 548 K), β (593 K), γ (638 K), δ (848 K), and ε, respectively.96 The normal (prototypic) phase corresponds to the high temperature δ-phase, which upon cooling transforms

According to the phase diagram, Cu2Se undergoes a solidstate phase transformation from the low-temperature, α-phase to the high-temperature, β-phase at 400 K.83 Both α- and β-phases are essentially a modification of the anti-fluorite structure, in which the Se atoms occupy the fcc lattice to form a rigid framework and the Cu atoms are distributed randomly on both trigonal and tetrahedral sites in the Se sublattice.84,90,91 The structure of the enigmatic α-phase is complicated because at least three types of crystal symmetry have been reported: monoclinic, tetragonal, and cubic.92,93 In the low-temperature α-phase, the Cu atoms are ordered and are not superionic, thus, the α-phase displays a very low electronic conductivity.94−97 However, at higher temperatures (>400 K), the Cu atoms become kinetically disordered throughout the fcc structure (Fm3̅m space group), which leads to superionic conductivity and high electronic conductivity in the β-phase.94,98 The electrical conductivity increases with decreasing Cu concentration, and this is assumed to result in an equivalent concentration of holes, as supported by transport measurements.99 It is generally accepted that the disordered β-phase is constructed by statistically distributing Cu atoms over the 8c tetrahedral interstitial sites in the fcc matrix, formed by Se atoms.100 This extraordinary “liquid-like” behavior of Cu ions, in the cubic lattice in the β-phase, results in an intrinsically low lattice thermal conductivity, which enables a high thermoelectric figure of merit (zT) in this otherwise simple semiconductor.98 The large positive effect of this phase transformation in Cu2Se has spurred great interest among the thermoelectric community because of the unique transport properties associated with this phase transition.92,98,101−108 Unlike Cu−S and Cu−Te, no hexagonal phase has been reported so far in the bulk for Cu2Se but a hypothetical unknown hexagonal modification has been proposed in various Cu−Se clusters.109−111 In general, the major crystal phases in both the bulk and in NCs are cubic berzelianite Cu2Se or Cu1.8Se. The Cu3Se2 phase (umangite) is stable only at a stoichiometric composition below 408 K and disproportionates to Cu2−xSe and CuSe above this temperature.83 Cu3Se2 has a tetragonal unit cell (P421m space group), which is composed of a tetrahedral network of Cu atoms, with Se atoms situated at atypical positions coordinated by six Cu atoms.84,112 It is noted that the structure of Cu3Se2 is derived from Cu1.8Se by rearrangement of the Cu atoms and compression of the lattice to the c-direction. In addition, a quasi-tetragonal phase associated with Cu3Se2 has been observed in some NC reports to-date.113−118 The CuSe phase (klockmannite) has a similar structure to that of CuS (covellite).119,120 In particular, CuSe has a hexagonal unit cell (P63/mmc space group) and an atomic arrangement that consists of alternating planar hexagonal hcp layers of [CuSe] and double layers of CuSe tetrahedra [Cu2Se2] in the c direction.121 Specifically, Cu atoms in the [CuSe] layers are triangularly coordinated by Se atoms, and Cu atoms in the [Cu2Se2] layers are surrounded by four tetrahedrally arranged Se atoms.122 However, unlike CuS, CuSe has a hexagonal superstructure and also undergoes phase transitions from the α-phase (room temperature, hexagonal unit cell), to the β-phase (327 K, orthorhombic unit cell) and the γ-phase (410 K, hexagonal unit cell).121,123−125 Copper diselenide (CuSe2) occurs in two different modifications, which are dependent on both temperature and pressure; marcasite is the low-temperature, low-pressure form, and pyrite is the high-temperature, high-pressure form.126 In both CuSe2 structures, each Cu atom is coordinated to six Se atoms in a distorted octahedral 5873


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into incommensurate γ- and β-phases and finally to the α-Cu2Te phase, which is composed of several commensurate structures.136 The highest temperature ε-phase has a cubic fluorite structure, while hexagonal-based superstructures are reported for the other phases (α, β, γ, and δ).138 Despite many years of study, the crystal structure of the binary Cu2Te compound has still not been well-determined. The hexagonal Nowotny structure was proposed in 1946 by H. Nowotny,134 where the Te atoms form trihedral prisms stacked along the c-axis and pairs of Cu atoms fill every second prism, located on lines parallel to the c-axis. However, recent theoretical studies have determined that the Nowotny model structure is energetically unfavorable and that the trigonal139 and monoclinic structures140 have much lower energies, thus forming more stable crystal structures. Moreover, Yu et al. predicted that the formation of Cu vacancies in Cu2−xTe induces a structural transition from monoclinic (most stable when x = 0), to trigonal (0.125 < x < 0.625) and then to the hexagonal Nowotny structure (0.75 < x < 1).27 Despite theoretical calculations showing that the hexagonal Nowotny structure has a low possibility of existing in stoichiometric samples, some reports on Cu2−xTe NCs have assigned the crystal structure to hexagonal weissite.141−144 Thus, it was proposed that the previous refinement of Cu2−xTe XRD spectra (using the Nowotny structure) should be re-examined by using the trigonal structure.27 The Cu7Te4 compound was found to crystallize in a layered structure of trigonal symmetry, where the Te atoms form a distorted hexagonal sublattice.145 The Te layers are connected by Cu atoms in two different tetrahedral holes, where a third Cu site within every second layer leads to an alternating Cu-rich and Cu-poor sandwich structure to form a layered-like network of Cu−Cu and Cu−Te bonds. There are some reports of hexagonal structured Cu7Te4 nanostructures in the literature to-date. 146−151 The thermal behavior of the Cu 3−x Te 2 compound was investigated in a separate study, and it was found to be orthorhombic at room temperature, tetragonal between 413 and 623 K, and cubic above 633 K.152 Cu3−xTe2 is regarded as having an incommensurately modulated crystal structure. The average structure of orthorhombic Cu3−xTe2 is a distorted defective Cu2Sb-type structure, in which the Cu atoms are in two different environments and can occupy the tetrahedral sites and pseudo-octahedral sites in the Te lattice.153 The origin of the incommensurable modulation can be given in terms of two competing interactions, between the ordering of Cu sites and the polarizability of the Te lattice, where the strongest displacive modulation is observed for the Cu atoms along the a direction and for the Te atoms along the c direction. CuTe (vulcanite) was determined to have an orthorhombic crystal structure. The arrangement of atoms in this structure can be described as a typical layered structure, where one Cu atom is distorted and is tetrahedrally coordinated to four Te atoms, and one Te atom is one-side coordinated to four Cu atoms.154 There is very little data regarding the oxidation states of Cu and Te in the literature. XPS measurements of Cu2Te NCs detected the presence of Cu+ and Te2−,155 whereas Cu2+ and Te2− were identified in CuTe nanostructures.156 Among the Cu chalcogenides, Cu2Te has gained interest as an attractive material for thermoelectrics and, in fact, many of the state-ofthe-art thermoelectric materials are based on tellurides, i.e. PbTe, Bi2Te3.157−159 Recently, Cu2S160 and Cu2Se98 were shown to be excellent thermoelectric materials, which possessed

exceptionally low thermal conductivities and high thermoelectric figures of merit (zT). Cu2Te is expected to possess a much lower lattice thermal conductivity, particulary when the character of the heavy element Te and the complicated Cu2Te crystal structure are considered.161 Moreover, Te is also less electronegative than S and Se, and so, the chemical bonds for Te should be less ionic than those for S and Se, and the carrier mobility should be large. Some studies have already highlighted the potential of bulk Cu2Te for thermoelectrics;96,161−163 thus, it is highly likely that Cu2Te NCs may contribute to new directions in this field. 2.2. Ternary Compounds

Cu chalcogenides are often found to be constituents of multinary adamantine compounds, as they prefer tetrahedral coordination due to their tendency for sp3 bonding.164 Adamantine ternary chalcogenides are conceptually derived from binary adamantine, by replacing a divalent cation (e.g. Zn) with one monovalent (e.g. Cu) and one trivalent (e.g. In) cation. Successive atomic mutations starting from the group-IV diamond structure result in a rich family of ternary Cu chalcogenide compounds that retain local tetrahedral bonding.165,166 This is the case of the Cu-III-VI, Cu-IV-VI, and Cu-V-VI families, which include several of the current, most interesting ternary Cu chalcogenides for technological applications. Some of these structures are vacancy assisted, and some contain lone electron pairs on the corresponding anions, which play a key role in thermoelectric applications.167 While the main multinary Cu chalcogenides crystallize in a diamond-like structure, Cu chalcogenide compounds with ordered chalcogen defective structures are also common. This is because defect pairs such as 2 VCu− + InCu2+ and 2CuIn2− + InCu2+ have particularly low formation energies in the ternary Cu-III-VI compounds, which explains the existence of unusual, ordered compounds such as CuIn5Se8, CuIn3Se5, Cu2In4Se7, and Cu3In5Se9.165 Heavy main group elements such as Sb and Bi generally result in strongly defective compounds, with quite different band structures. For example, Cu3SbSe3 is characterized by larger band gaps than its diamond-like counterparts.168 A main feature of defective ternary structures like Cu3SbSe3 is that the local void spaces near the main group ions, which are induced by the stereochemically active s2 lone pairs, result in ultralow thermal conductivity.169 2.2.1. Cu-III-VI. From a technological point of view, a particularly attractive family of Cu chalcogenides are the Cu chalcopyrite compounds. These compounds have the general stoichiometric formula Cu-III-VI2 (where III = Al, Ga, In; and VI = S, Se, Te) and form a large group of semiconducting materials, with diverse optical, electrical, and structural properties that are suitable for numerous applications. These properties are dominated by the presence of intrinsic defects (i.e. vacancies, interstitials, and anti-site defects), which occur when a compound deviates from the 1:1:2 (Cu-III-VI2) stoichiometry.170,171 The fact that Cu-III-VI2 chalcopyrite compounds can tolerate a large range of anion and cation offstoichiometry (unlike their binary II-VI analogs) allows them to induce different doping defects, where either p- or n-type materials can be produced via stoichiometry control.165 These off-stoichiometry effects can be used to influence the conductivity,172 band gap,173 and luminescence properties.12,174−176 The most explored ternary compounds are CuInS2 and CuInSe2, which exhibit a chalcopyrite phase (named after the mineral chalcopyrite, CuFeS2) at room temperature. Depending 5874


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Figure 3. Unit cell structures of (a) chalcopyrite, (b) zinc blende, and (c) wurtzite CuInS2, with (d) their simulated XRD patterns. Reproduced with permission from ref 183. Copyright 2016 American Chemical Society.

polarities of the crystal facets in this structure can be exploited for the growth of anisotropic NCs (e.g. nanorods). Until recently, it was believed that the Cu and In cations were randomly distributed (i.e. cation disordered) in the wurtzite phase. However, a recent fundamental study discovered a new form of crystalline order in wurtzite CuInS2 NCs and showed that the wurtzite structure is cation ordered.186 These “interlaced crystals” have a perfect, uninterrupted anionic lattice, in which Cu and In have an infinite set of possible ordering patterns within the cation sublattice and form interlaced domains and phases. This ordering has important implications for phonon and electron transport, in that electronic motion is unaffected by interlaced phase boundaries but phonon mobility is likely to be significantly reduced.183,186 Thus, interlaced crystals should have high electron conductivites and low phonon conductivities, which make them ideal candidates for thermoelectrics. An increase in the thermal resistivity of up to two orders of magnitude has already been predicted in these interlaced crystals.187 In terms of their XRD patterns (Figure 3d), the chalcopyrite and zinc blende phases cannot always be easily distinguished from one another, especially if the reflections are broadened from the small NC sizes. They also show similar lattice fringing in high-resolution TEM (HRTEM) analysis, but the main difference is evident through close examination of the XRD patterns, where the low-intensity (101), (103), and (211) reflections are present in the chalcopyrite phase but are absent in the zinc blende phase.183 In contrast, the wurtzite phase can be easily identified by both its XRD pattern and the lattice fringes in HRTEM imaging. From the phase diagram of the Cu2S−In2S3 system, a second semiconductor compound, CuIn5S8, was determined to have a cubic spinel-type structure over the whole temperature range from 20 to 1085 °C.177 This In-rich compound is typically an n-type semiconductor that has a band gap of 1.3 eV. However, this n-type semiconductor has proven to be difficult to prepare, with most methods yielding the CuInS2 compound with p-type conductivity. In terms of CuIn5S8, almost all of the reported methods to-date are mainly focused on thin films188−190 and some studies have investigated the electronic and optical properties of CuIn5S8.191−194 In comparison to CuInS2, there are very few literature reports on the preparation of CuIn5S8 NCs.183,195,196 CuInSe2 is another Cu-III-VI2 compound semiconductor that can exist in different crystal structures. The earliest published

on the Cu/In ratio, these compounds can display n-type (Cu/In ratio 1) conductivity. The phase diagram of the Cu2S−In2S3 system was determined by Binsma et al., where it was revealed that CuInS2 exists in three polymorphic modifications: (i) in the chalcopyrite structure from ambient temperature up to 980 °C; (ii) in the zinc blende structure from 980 to 1045 °C; and (iii) in the wurtzite structure from 1045 to 1090 °C.177 In contrast to the bulk, all three crystal structures can be stabilized in NCs at room temperature.1 In the chalcopyrite phase, the Cu and In atoms are ordered within the unit cell (Figure 3a), and each S atom is tetrahedrally coordinated to two Cu atoms and two In atoms. The structural complexity of this phase results in a lower bandgap energy178 and an abundance of intrinsic defects,171 in comparison to II-VI compounds. The Cu-III-VI2 compounds with the chalcopyrite phase (tetragonal structure) can be envisioned as ternary analogs of the II-VI zinc blende binary compounds,179−181 with cations of higher valency (In3+) and lower valency (Cu+) occupying the cation sublattice in an ordered manner. It was noted that Zn0.5Cd0.5S is the binary analog of the CuInS2 compound, whereas ZnS is the binary analog of CuGaS2.182 The unequal valence of the Cu and In cations and the corresponding unequal bond lengths (RCu−S ≠ RIn−S) cause anion displacement away from a close-packed arrangement, which leads to reduced crystal symmetry and tetragonal distortion of the lattice, where the c-axis equals approximately twice the length of the a-axis. The chalcopyrite to zinc blende phase transition corresponds to a disordering of the cation sublattice, where a random distribution of the Cu and In cations (Figure 3b) becomes thermodynamically favored. Thus, the zinc blende phase is analogous to the chalcopyrite phase with related lattice parameters (cCP = 2cZB), but the key difference is that it is cation disordered.183 For example, in the zinc blende structure, cations can be easily exchanged for one another, leading to various stoichiometries, ranging from Cu3InS3 to CuIn2S3.5.184 In the wurtzite structure (Figure 3c), the sulfur anions form a hexagonal close packed (hcp) arrangement, whereas the Cu and In cations occupy the tetrahedral interstices of the sulfur framework. It is a hexagonal analogue of the zinc blende phase, but its anion stacking sequence (abab) is distinct from that of the zinc blende (abcabc) phase, along the [0001]WZ and [111]ZB directions.185 Moreover, the wurtzite phase is of particular interest, as it is anisotropic in the c-direction; thus, the different 5875


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study of the Cu−In−Se phase diagram197 was restricted to a segment of the Cu2Se−In2Se3 pseudobinary section and centered on the CuInSe2 compound.198 However, since then, there have been many studies on the phase relations in this ternary system to understand the thermochemistry.199−201 The CuInSe2 compound has two polymorphic modifications, which are separated by a first-order solid−solid phase transition between the chalcopyrite and sphalerite phases at 810°C.199 In the chalcopyrite phase, the Cu and In cations are ordered in the Se anion lattice and each Se has two Cu and two In ions as its nearest neighbors. The chalcopyrite structure has a specific cation ordering and requires a tetragonal unit cell, with Cu occupying the (0,0,0) sites and In occupying the (1/2,1/2,0) sites of the cation sublattice.202 However, these cations become disordered at the elevated temperature, resulting in the cubic sphalerite phase. Since the CuInSe2 chalcopyrite and sphalerite XRD patterns are almost identical, a low intensity, characteristic chalcopyrite (211) peak can be used to differentiate between them.203 In NCs, an additional hexagonal wurtzite phase of CuInSe2 has been reported.204,205 Similar to CuInS2, a large degree of off-stoichiometry can be tolerated in the Cu−In−Se compounds, which leads to a series of compounds with different stoichiometries, such as CuIn5Se8, CuIn3Se5, Cu2In4Se7, and Cu3In5Se9.165 Since these In-rich phases are often difficult to distinguish by XRD, elemental analysis is generally employed to aid in their differentiation from CuInSe2. These small deviations in composition can produce large changes in the electrical properties, as well as in the band gap and electronic band structure.206 Moreover, Cu rich films are found to mostly exhibit p-type conductivity, but as the films become In-rich, a conversion to n-type conductivity occurs.207 Some of these In-rich structures likely result from the reduced formation energies of cation vacancy sites in the chalcopyrite phase and others as the result of energetically stabilized defect pairs; thus, defects are an important topic in the Cu−In−Se system.165,203 At present, there is no agreement between the numerous studies of these In-rich materials on the phase boundaries’ compositions, on the various phases that lie between CuInSe2 and CuIn5Se8, or on their crystal structures.198 2.2.2. Cu-IV-VI. The Cu-IV-VI compounds (where IV = Sn, Ge) have attracted significant research interest because of their interesting properties and their potential applications for nonlinear optics, photovoltaic cells, energy storage, and thermoelectrics. The most explored Cu-IV-VI compounds are in the Cu−Sn−S system, of which Cu2SnS3208 is the most commonly explored compound. Other nonstoichiometric compounds, such as Cu4SnS4,209 Cu4Sn7S16,210 CuSn3.75S8,211 and Cu3SnS4,212 have also been reported in this system. The Cu-IV-VI compounds (IV = Sn, Ge) have a tendency to crystallize in a large number of phases and structural forms, such as a cubic sphalerite-like phase (F43̅ m space group); a monoclinic sphalerite superstructure; an orthorhombic structure (Imm2 space group);208 and a metastable wurtzite phase, which is exclusive to NCs and cannot be stabilized in the bulk.213−215 In order to describe the differences between these structures, we have focused on the Cu2GeSe3 system. In this system, three possible crystallographic phases/structures have been reported (Figure 4): an orthorhombic structure (Imm2 space group),216 a cubic zinc blende structure (F4̅3m space group),217 and a hexagonal wurtzite structure (P63/mc space group) in NC form.215 These phases are intimately related and

Figure 4. Unit cells of the different crystallographic phases in the Cu2-IV-VI3 system. Adapted with permission from ref 215. Copyright 2012 American Chemical Society.

are similar to those found in the group III-V and group IV semiconductors. In the orthorhombic phase, the cations are ordered such that all of the cation positions in every plane are occupied by the same element and follow an ordered sequence of two planes with Cu cations, plus one plane with Ge cations. On the contrary, in the cubic zinc blende structure, the Cu and Ge cations are randomly distributed in the cation positions, with occupation factors of 2/3 and 1/3 for Cu and Ge, respectively. The orthorhombic and cubic zinc blende structures can be understood as an abcabc stacking along the respective planes, whereas the hexagonal wurtzite structure has an abab stacking sequence. By visualizing either the orthorhombic or the zinc blende structure, the pure hexagonal wurtzite structure can be envisioned as a sequence of consecutive twin defects in those structures [corresponding to a 180° rotation of the structure along the corresponding axes]. 2.2.3. Cu-V-VI. The Cu-V-VI (V = Sb, Bi) compounds are the least explored family of ternary Cu chalcogenide semiconductors, despite their high technological interest. There are many crystallographic phases and structural forms possible in the Cu−Sb−S system, such as CuSbS2, Cu3SbS4, Cu3SbS3, and Cu12Sb4S13 (Figure 5), and the structural differences between these compounds give rise to different functional properties and potential applications. CuSbS2 crystallizes in the orthorhombic structure (Pnma space group) and has a layered structure, consisting of SbS2 and CuS3 chains along the b-axis, which are formed by linkage of Sb square pyramids and CuS4 tetrahedral links.219 These two infinite chains are interconnected to produce layers that are perpendicular to the c-axis, with 2.051 Å separation. This separation distance allows for the intercalation of small atoms, ions, or molecules; thus, CuSbS2 could be envisaged as a potential electrode material for lithium/sodium ion batteries.220 Cu12Sb4S13 forms part of the tetrahedrite family and has a cubic sphalerite-like structure (I4̅3m space group).221 In particular, six of the 12 Cu atoms occupy trigonal planar 12e sites, and the remaining Cu atoms are distributed on tetrahedral 12d sites.222 Moreover, four of the six tetrahedral sites are thought to be occupied by monovalent Cu, and the other two sites are occupied by Cu2+ ions, whereas the trigonal planar sites are occupied exclusively by monovalent Cu.223 The Sb atoms also occupy a tetrahedral site but are bonded to only three S atoms, leading to a void in the structure and a lone pair of electrons, as in Cu3SbSe3. A combination of factors, such as the large number of atoms per unit cell, the large Gruneisen parameters (large anharmonicity),222 and the low energy vibrations of the Cu atom out of the [CuS3] trigonal planar unit, result in very low thermal conductivity in this material.224 In terms of the Cu-IV-VI compounds comprised of Se, the Cu3SbSe4 compound crystallizes in a zinc blende-type tetragonal 5876


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Figure 5. Schematic representations of (a) orthorhombic CuSbS2, (b) orthorhombic Cu3SbS3, (c) tetragonal Cu3SbS4, and (d) cubic Cu12Sb4S13. Adapted with permission from ref 218. Copyright 2016 Royal Society of Chemistry.

superstructure (I42̅ m space group),225 in which two different Cu positions with different Cu−Se bond lengths are present. Each Cu and Sb atom has four Se neighbours, which are tetrahedrally coordinated. This configuration can be envisioned as a threedimensional Cu−Se framework of distorted [CuSe4] tetrahedra, with inserted one dimensional arrays of [SbSe4] tetrahedra.226 The valence band maximum is mainly formed by the hybridization of Cu 3d and Se-4p states, while the conduction band minimum consists of mainly Sb-5s and Se-4p hybridization. The Cu3SbSe4 compound has a small direct band gap (0.29 eV), a large effective mass (1.5 me),227 and a defect-related carrier density on the order of 1018 cm−3.228 The high carrier concentration in this compound is associated with the high concentration of intrinsic point defects, where the Cu vacancy has a low formation energy with a tetrahedrally coordinated Cu. These Cu-vacancies, with quite small ionization energy, act as shallow acceptors and release holes; thus, a stable carrier concentration within a large temperature range is maintained.226 From an applications point of view, the Cu3SbSe4 compound is a promising thermoelectric (TE) material because of its relatively high hole mobility (associated with the Cu−Se framework) and the simultaneous low phonon mean free path, induced by a rather complex crystal structure. In addition, it has an appropriate electronic band structure that includes a large degeneracy at the valence band maximum, which makes it interesting for TE applications (as discussed in section 11).229 The Cu3SbSe3 compound typically crystallizes in an orthorhombic structure (Pnma space group)230 and consists of SbSe3 triagonal-pyramidal units, in which the Sb ion is in a trivalent state. Compared to Cu 3 SbSe 4 , the Cu 3 SbSe 3 compound has a more complicated crystal structure, in that it is comprised of two inequivalent Cu atoms (Cu1 and Cu2) and two inequivalent Se atoms (Se1 and Se2). Specifically, Cu1 has three Se neighbors (Cu−Se bond length: 2.41 Å) and one Cu (Cu−Cu bond length: 2.60 Å); Cu2 is additionally coordinated with another Cu atom; Sb has three Se neighbors (Sb−Se bond length: 2.64 Å); and Se1/Se2 is neighbored by three Cu and one Sb. Furthermore, Sb exists in the Sb3+ state in the Cu3SbSe3 compound, whereas it occurs in the Sb5+ state in the Cu3SbSe4 compound. This results in two 5s valence shell electrons being in an unbonded state (i.e. lone pairs), which cause nonlinear repulsive forces and strong anharmonicity, thus leading to strong phonon−phonon scattering.231 This, coupled with the complexity of its crystal structure, results in intrinsically low thermal conductivity in this material.232

in the attainment of quaternary chalcogenide semiconductors that retain local tetrahedral bonding.165,166 This is the case of the Cu-III-VI, Cu-II-IV-VI, Cu-II-III-VI, and Cu-III-IV-VI compositions. The octet rule plays an important role in the energy stability of the different structural configurations and assumes that configurations with all anions in an eight-electron closed-shell state have lower energy. Quaternary chalcogenides have more complicated electronic and structural properties than binary and ternary chalcogenides because of the increased number of elements and structural configurations. For example, in Cu2ZnSnS4 (CZTS), this manifests in an intense competition between its binary components, such as CuS, ZnS, and SnS, or its ternary Cu2SnS3 compositions.233 Furthermore, point defects are easily formed in quaternary chalcogenides, due to the similar radii of each ionic species. These include vacancies, antisites, and interstitials (e.g. CuZn, VCu, VZn, ZnSn, CuSn), which can result in trap states deep within the band gap.233,234 Since quaternary chalcogenides are derived from their binary counterparts (i.e. ZnS), they also share the same basic crystalline structures, such as the zinc blende (F43m space group) and wurtzite (P63mc space group) phases. 2.3.1. Cu-III-VI. Quaternary Cu-III-VI compounds are derived from the ternary Cu-III-VI chalcopyrite structures by the additional occupation of a group III atom. This results in the formation of the CuInGaS2 (CIGS) and CuInGaSe2 (CIGSe) compositions, which have the general stoichiometric formula Cu-III-VI2. Similar to the ternary I-III-VI2 chalcopyrite structure, there are three crystallographic phases possible for the quaternary Cu-III-VI2 compositions: (i) chalcopyrite; (ii) zinc blende; and (iii) wurtzite. The most stable crystal structure of CIGS and CIGSe is the chalcopyrite structure (I4̅2d space group), in which the Cu, In, and Ga ions are ordered in the cation sublattice.235 This is different from the zinc blende structure, where the cations are randomly distributed. Further disordering of these cations yields the kinetically stable, wurtzite structure, which has a hexagonal crystal system. The wurtzite structure differs from zinc blende in that it has a noncentrosymmetric structure, where the Cu, In, and Ga cations occupy the lattice in a disordered manner. Thus, the wurtzite structure offers more flexibility for stoichiometry control, than the chalcopyrite structure, and the ability to tune the Fermi energy over a wide range in this structure allows for the attainment of exciting optical properties and optimal device performance.236 2.3.2. Cu-II-IV-VI. Quaternary Cu-II-IV-VI compounds (where II = Zn, Cd, Mn, Fe, Co, Ni; and IV = Sn, Ge) are formed when the chalcopyrite Cu-III-VI structure mutates exclusively into different structural configurations that obey the octet rule. As early as the 1950s, efforts had been made to

2.3. Quaternary Compounds

Successive atomic mutations starting from the group-IV diamond structure, and transitioning through the binary group II-VI and ternary I-III-VI2 chalcopyrite structures, result 5877


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is absent in PMCA, which makes it primitive tetragonal, and this allows it to be differentiated from stannite. There are also different tetragonal structural modifications of the kesterite structure that can be considered,243 which are restricted to the exchange of two ions (Cu and Zn) in the cation sublattice and belong to the P4̅2c, P4̅21m, and P2 space groups.240 In all these structural modifications, the sulfur fcc lattice determines the unit cell dimensions. Despite XRD being the most prevalent technique used for crystal phase determination, it is not sensitive to the metal cation arrangement in Cu-II-IVVI compounds because the occupation of the atomic sites by Cu, Zn, and Sn results in similar X-ray scattering factors, which makes the ordering difficult to determine. For example, the XRD pattern of kesterite is strikingly similar to that of stannite. This explained why neutron scattering was the technique employed to differentiate between the different structures.240 Thus, it is not only challenging to differentiate between the kesterite, stannite, and PMCA structures using XRD but also to distinguish between CZTS and secondary phases such ZnS and Cu2SnS3. However, resonant Raman spectroscopy has been shown to be an effective tool to distinguish between the CZTS, ZnS, and CTS phases.242 Thus, Raman is commonly employed as an additional characterization technique to determine if the CZTS composition is formed and if it exists as a single phase. While the crystal structure of CZTS is similar to ZnS, its band gap is reported to be 1.5 eV, which is closer to Cu2S rather than to the wide-band-gap ZnS.6 An investigation on the band structure of CZTS revealed that its conduction band is dominated by the Sn 5s/S 3p orbital and that its valence band is dominated by the Cu 3d orbital hybridized with S 3p states, which is the same as the valence band of Cu2S.243 In stoichiometric CZTS, the dominant defect is the p-type CuZn antisite, which has an acceptor level deeper than the Cu vacancy. This leads to a dominant self-compensated [Cu−Zn + Zn+Cu]0 defect pair, which results in poor carrier separation. Based on this, it was predicted that CZTS should have a Cu-poor/Zn-rich composition so that Cu vacancy defects and Cu−Zn antisites dominate,244 and this composition has so far produced the highest PV efficiencies in CZTS devices to-date.245 While the zinc-blende structures (kesterite, stannite, PMCA) have been covered for the Cu-II-IV-VI compounds, wurtzite and wurtzite-derived structures (Figure 7) can also exist for this family of compounds. In the wurtzite structure, the Cu+, Zn2+, and Sn4+ cations are randomly distributed in the lattice. The wurtzite phase is particularly exciting, as it offers more flexibility for stoichiometry control, entailing band-gap engineering by the ability to tune the Fermi energy over a wide range. While it is not stable in the bulk, it can be stabilized in NC form as a metastable phase, which presents a low barrier for rapid conversion to the kesterite phase with composition control.246 The orthorhombic structure is another metastable that exists for CZTS and is based on a double-wurtzite unit cell (2 × 2 × 1 supercell), where the arrangement of the anions is hexagonal close packed. The orthorhombic unit cell for these derived structures was originally based on AgInS2 compound, which was experimentally found to crystallize in an orthorhombic wurtzite-like phase (Pna21 space group) at temperatures higher than 620 °C.247 This configuration was described as having a wurtzite-chalcopyrite structure because it has a similar cation arrangement to the chalcopyrite structure. Wurtzite (WZ)derived orthorhombic structures have already been reported for the Cu2ZnGeS4, Cu2ZnGeSe4, and Cu2CdGeS4 compositions. A wurtzite-CuAc-like (PMCA) structure (Pmc21 space group)

predict the design of quaternary chalcogenide semiconductors by the cross-substitution of cations in the ternary I-III-VI2 compounds.237,238 The replacement of two group III atoms by one II and one IV atom generates the Cu2-II-IV-VI4 family of compounds, which gives rise to Cu2ZnSnS4, Cu2ZnGeS4, Cu2CdSnS4, Cu2CdGeS4, Cu2FeSnS4, and Cu2CoSnS4 compounds, amongst others. Another possible configuration is the replacement of one group I atom and one III atom by two II atoms, which generates the Cu-II2-III-VI4 compounds. This constitutes the CuZn2GaS4 or CuZn2GaSe4 compounds, which are an alloy of the II-VI and I-III-VI2 compositions.239 Among these materials, Cu2ZnSnS4 (CZTS) and its selenium analog, Cu2ZnSnSe4 (CZTSe), have become the subject of intense research interest because they are made of earth abundant and non-toxic elements and their band gaps cover the optimal energy range of photovoltaic applications. Thus, to explain the crystal structures in the Cu2-II-IV-VI4 compounds, we will focus on the CZTS system. In CZTS, there are three zinc-blende derived cell structures: (i) kesterite (I4̅ space group); (ii) stannite (I4̅2m space group); and (iii) a primitive mixed CuAu-like structure (P4̅2m space group). These three structures are closely related, as shown in Figure 6, but they are assigned to different space groups due to

Figure 6. Unit cell structures of the kesterite, stannite, and primitive mixed CuAu-like (PMCA) phases in the CZTS system. Reprinted with permission from ref 242. Copyright 2012 American Institute of Physics.

the different arrangements of the Cu+, Zn2+, and Sn4+ cations within the tetrahedral voids. Specifically, the kesterite structure (Figure 6a) is a 1 × 1 × 2 expansion of the cubic zinc blende lattice and is characterized by two alternating cation layers, each containing Cu and Zn or Cu and Zn in the direction of the c-axis at z = 0, 1/4 , 1/2, and 3/4, respectively.240 In general, bulk CZTS favors the tetragonal kesterite structure because this is the lowest energy, ground state structure according to firstprinciple, total-energy calculations.241 The most important structural modification of kesterite is stannite, which is typically observed for the closely related Cu2FeSnS4 compound.167 Compared to kesterite, stannite has a different cation layer distribution, in that a layer of Zn and Sn alternates with a Cu layer in the crystallographic c direction at z = 0, 1/4, 1/2, and 3/4, respectively. In the stannite structure (Figure 6b), the Zn and Sn atoms on each layer switch their positions every other layer. While the primitive mixed CuAulike (PMCA) structure (Figure 6c) also contains an alternating ZnSn and Cu layer, the location swapping between Zn and Sn 5878


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chalcogenides in Table 1, with details of the crystal phase, band gap (Eg), band alignment, light absorption coefficient, conductivity type, exciton Bohr radius, their potential applications, and whether these measured properties are associated with the NC form or not. In the following subsections, a brief overview of the functional properties of Cu chalcogenides is provided, with special emphasis on those particular to NCs. 3.1. Electronic Properties

The functional properties of Cu chalcogenides are mainly determined by the energy band structure, including band gap, degeneracy of near-edge bands, band alignment, and density/type of defects in the Cu chalcogenide compound. Surface composition/defects can also play a key role in Cu chalcogenide NCs. 3.1.1. Band Structures. The main parameter that differentiates chalcogenides and oxides is the higher covalence of the chalcogenides, which results in broader bands and narrower energy gaps than in oxides. The small size and electronegativity differences between transition metals and chalcogens also results in higher concentrations of off-stoichiometries and electronically active defects, compared to oxides. A particularity of Cu chalcogenides is the strong effect of Cu 3d states in its electronic properties. Electronic hybridization in chalcogenides is usually quite strong, due to the highly delocalized p orbitals in the chalcogens. In Cu chalcogenides, the hybridization can be even stronger because the atomic levels of Cu 3d states are just slightly higher than the p orbitals of the chalcogens. A table of the relevant atomic levels of s, p, and d states is provided in Table 2. The degree of involvement of the d orbitals of the transition metal plays a key role in defining the functional properties of transition metal chalcogenides.24,26,390,391 While the strength of hybridization between two electronic states is determined by their wave function overlap and energy separation, the strong localization of 3d states in transition metals usually prevents them from undergoing strong chemical hybridization with other states. Although Cu 3d electrons are intrinsically rather localized in Cu, the Cu 3d states are quite high in energy and are strongly mixed with the delocalized chalcogen-p states, which forms a strong p-d hybridization. The involvement of shallow Cu 3d semicore states plays a key role in defining the band structure of Cu chalcogenides, including the band gap, degeneracy of near-edge bands, and band alignment.168 In particular, the strong p-d hybridization results in a very significant contribution of the Cu 3d states to the near-edge valence bands. For example, Cu 3d states are shallower than Zn 3d states, compared to ZnS. Thus, the p-d hybridization is much stronger in Cu chalcogenides (than in ZnS) and there is a more substantial contribution of Cu 3d components to the valence band maximum states (VBM). For example, in the CuGaS2 compound, the Cu 3d electrons in the t2 suborbitals strongly hybridize with the S 3p states. This is different from the ZnS compound, which has a weak p-d hybridization, in that the Zn 3d orbitals are not only fully occupied but are also much lower in energy than the S 3p states. As a consequence, the VBM of CuGaS2 is pushed upward, which results in a much smaller band gap for CuGaS2 (2.43 eV), compared to ZnS (3.78 eV) (Figure 8).168 In most Cu chalcogenides, the VBM is an antibonding state of the anion p orbitals and the Cu d orbitals.181,392−395 In addition, the VBM of sulfides is lower than that of selenides and

Figure 7. WZ-derived structures of the Cu-II-IV-VI quaternary compositions. Adapted from ref 250. Copyright 2014 Royal Society of Chemistry.

was predicted by density functional theory calculations, combined with evolutionary methodology.248 Wurtzite-kesterite and wurtzite-stannite structures have also been theoretically reported, which have similar cation arrangements to the respective kesterite and stannite structures, and they indicate that (Cu+Zn) disorder is possible in these structures.249 Since the Cu-II-III-VI composition is an alloy of the II-VI and I-III-VI compounds, it has different cation distributions around the group VI anions, compared to the Cu-II-IV-VI compounds.239 The stannite phase is the ground state structure in CuZn2GaS4 and CuZn2GaSe4. However, the energy differences between the kesterite, stannite, and PMCA structures are small in the Cu-II-III-VI compositions, and this indicates a higher possibility of forming mixed superstructures.

3. FUNCTIONAL PROPERTIES Cu chalcogenides show a broad spectrum of functional properties, due to their structural and compositional versatility, which have been exploited in a large variety of applications. In particular, multinary compounds offer an extraordinary flexibility for material design, and their functionality can be optimized by using only nontoxic and abundant elements. For example, the ternary CIS compound provides a proper direct band gap, which allows for the replacement of the toxic CdTe compound as an absorber in photovoltaic (PV) devices. Moreover, the partial substitution of In by Ga to form the quaternary CIGS compound provides a better adjustment of the band gap and the bands alignment with window materials. Beyond CIGS, quaternary CZTS or quinary CZTSSe compounds are particularly attractive materials for PV applications, as they reduce the dependence on scarce and expensive elements, such as In.251 Numerous Cu chalcogenide compounds have been synthesized and thoroughly studied, but several other compounds still remain unexplored.252 To provide an overview of the wide range of compositions attainable in Cu chalcogenides, we have summarized the measured properties and applications of Cu 5879


Crystal Phase

H, MC, OTR C, H, OTR, TR

H, OTR, TG H, C, TG C TR, H, C TG H TG H, OTR, C, TG TR, H, C TG TG TG TG OTR OTR TR, H, C, TC MC, H, C C, OTR MC, H, OTR, C OTR, C, TG TG, OTR TR, OTR H, OTR, TG H, OTR, TG TG

TG H, OTR, TG H, OTR, TG C, MC, OTR OTR OTR C OTR C TG TG OTR C

Material

Cu2−xS Cu2−xSe

Cu2−xTe CuInS2 CuIn5S8 CuInSe2 CuIn3Se5 CuIn5Se8 CuInTe2 CuGaS2 CuGaSe2 CuGaTe2 CuAlS2 CuFeS2 CuTlSe2 CuBiS2 CuSbS2 Cu2SnS3 Cu2SnSe3 Cu2SnTe3 Cu2GeS3 Cu2GeSe3 Cu2GeTe3 Cu3SnS4 Cu3AsS4 Cu3AsSe4 Cu3SbS4

5880

Cu3SbSe4 Cu3PS4 Cu3PSe4 Cu3SbS3 Cu3SbSe3 Cu3BiS3 Cu12Sb4S13 Cu4Bi4S9 Cu2SnSe4 Cu2MoS4 Cu2WS4 Cu4SnS4 CuCr2S4

1.22−1.60, 3.67 1.19−1.24 0.68−0.88 0.47 0.8−1.2 0.11−0.4 2.00−2.38 1.30−1.40 1.15−1.84 0.95−1.70 1.2−1.84 1.6−1.72 1.34 0.88−1.14 0.35−0.43 ∼1.71−1.74 1.74−2.15 1.0−2.32

2.19−2.62 0.90−1.90 0.59−1.8 0.44−0.96 1.18 1.5−1.6 0.78

1.2−2.0 2.0−2.3 1.1−1.5 1.1−1.5 1.45−1.50 1.29−1.51 1.04 1.21 1.15 0.96−1.06 2.43 1.68 1.23 3.49 0.50−2.0

Eg (eV)

direct

direct direct indirect direct indirect direct indirect direct indirect direct direct direct indirect indirect direct direct indirect direct direct direct

direct indirect direct indirect direct direct direct indirect direct direct direct direct direct direct direct direct direct indirect direct direct indirect direct direct direct direct direct direct

Band alignment

p

p p

∼105

p p p p p p p

p p n p n n p p p p p n, p p p p p p p p p p p p p

p p

Conductivity type

102

>105 ∼105

∼5 × 104 ∼105

∼105

∼104

∼104

∼105 ∼104 ∼105 >104

∼105 105 1−2 × 105 ∼105 ∼104

∼105

∼105 ∼104

4

∼10 ∼105

Light absorption coefficient (cm−1)

Table 1. Measured Properties and Applications of Cu Chalcogenide Compounds

3.3−5.8

2.5−4.6

4 5.1

10.6

4.1

3−5

TE, S, ES CAT, LED, S, BIO LED CAT, LED, BIO

TE CAT PV, CAT PV, ES, TE PV, TE PV, TE, ES, BIO PV, TE PV TE ES, CAT CAT PV, TE TE, SpE

TE, PV, Dc

PV, CAT PV

PV PV, TE

PV, TE PV, CAT, L/D, S PV, CAT, L/D, S TE PV, BIO ES, TE TE PV, CAT PV, ES PV, ES, CAT PV, TE

PV, PV, PV, PV,

PV, BIO, TE, CAT, S, ES PV, CAT, TE, S, ES, BIO

Exciton Bohr radius (nm) Applications 6,37,42 128,253,254

√ √

225−227,294 295−298 298,299 218,271,300 230,301−303 304−313 218,221,262,271,292 314−317 318−321 322−324 325−327 209,318,328,329 330,331

x x x √ x √ √ √ √ √ √, x √ x √ x x √ √, x x √ √, √, √, √, √, √ √ √ √

x x x x x

x

255 9,256,257 191,195 12,181,256,258,259 259 259 12,260 12,235,261 12,262,263 12,181 12,181 .264−266 267 268 269−272 273−275 275−277 278 279−281 215,281−283 284 285,286 287−290 287 271,291−293

√ √ x √ x x x

Ref

NC form?

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265,344,374

266,375−381 267−269,381,382 270,271,376,381,383 381,384−388 x

x x x

344,346−350 351−354 355,356 347,357−360 347,361 362,363 364−369 347,370−372 260,264,373

CAT, PV TE, PV TE TE, TP, PV CAT, S, PV, TE CAT, FE, PV, TE CAT, ES, PV, TE PV, CAT, TE 104−105 >104

∼104

∼103

3 × 102

direct indirect direct direct direct direct direct direct direct 2.05−2.25 1.21−1.63 0.17 0.54, 1.6 1.58 1.1−1.69 1.2−1.58 1.28−1.54

3 × 102

p p p p p p p p

∼6 p p p p p ∼105 ∼104 >104

MC, H, TG, C TG TG, H TG TG, OTR TG, OTR TG, OTR H, TG, OTR TG TG H, TG C, H, TG C, H, TG C, H, TG Cu2ZnSnSe4 Cu2ZnSnTe4 Cu2CdSnS4 Cu2CdSnSe4 Cu2CdGeSe4 Cu2FeGeS4 Cu2ZnGeS4 Cu2ZnGeSe4 Cu2HgSnSe4 Cu2HgGeSe4 Cu2CoSnS4 Cu2MnSnS4 Cu2NiSnS4 Cu2FeSnS4

direct indirect direct direct direct

∼105 direct C C H, OTR, TG, C CuCr2Se4 CuCr2Te4 Cu2ZnSnS4

1.45−1.8

Crystal Phase

1.0−1.5 0.58−1 1.4, 1.52 0.98 1.06−1.16

2.5−3.3

TE, SpE TE, SpE PV, CAT, LED, TE, ES, S, TP PV, CAT, TE, S, TP TE, PV PV TE, PV TE, PV p p p

tellurides, due to the p level of S being lower than that of Se and Te. Compared with other chalcogenide compounds (e.g. ZnX), the VBM variation is relatively small in Cu chalcogenides due to the p-d hybridization, which is stronger in the shorter Cu−S bond. For most Cu chalcogenides (with the exception of the Cu-Sb-VI compounds), a higher Cu concentration results in stronger p-d hybridization and, thus, a higher VBM position. Thus, binary Cu-VI compounds (i.e. Cu2S, Cu2Se, Cu2Te) have the highest VBMs, followed by the Cu2-IV-VI3 and Cu-III-VI2 compounds, with the Cu2-II-IV-VI4 compounds showing the lowest VBMs.396 In the ternary Cu-III-VI2 compounds, the conduction band minimum (CBM) is mainly composed of Cu 3s orbitals,which are hybridized with the p orbitals of the group III cation, and the VBM is dominated by the Cu 3d orbitals,which are hybridized with p orbitals of the group VI anion. 6,397,398 Thus, the band gaps in the Cu-III-VI 2 compounds can be adjusted by selecting different group III and VI elements, or by using combinations of different group III or VI elements with adjusted ratios, to form compounds such as CuInGaS2 (CIGS) and CuIn(S1−xSex)2. In this regard, the addition of a group III element leads to a decrease in the band gap in the order Al > Ga > In, and similarly in the order of S > Se > Te for the chalcogen.181 In the Cu-III-VI2 chalcopyrite crystal structure, the two cations are tetrahedrally coordinated by four anions, but the anion is coordinated by 2I + 2II cations, which have unequal bond lengths (RI-VI ≠ RIII-VI), and this results in a tetragonal unit cell. In addition, the Cu-chalcogenide bond influences the VBM (i.e. VBM is lower with a longer bond) and a shorter main metal−chalcogenide bond results in a higher CBM, due to a stronger main metal-s, chalcogen-p hybridization. Thus, the difference between the metal and chalcogenide bond lengths has a strong influence on the band gap of the ternary chalcogenide compounds.168,399 As noted above, the strong p-d hybridization is responsible for the considerable band-gap reduction in CuGaS2, due to an upward shift of the VBM. With respect to the quaternary compounds, the main differences are found in the conduction states. For example, by moving from CuGaS2 to Cu2ZnSnS4 (CZTS), it was noted that the Sn 5s orbital is deeper than the Ga 4s orbital (−10.8 eV, compared to −9.2 eV for Ga 4s), and this leads to a further reduction of the band gap.168 Furthermore, the CBM in CZTS is dominated by Sn and S, and so, by partially replacing Sn and S with alternative elements (e.g. Ge and Se, respectively), this allows the band gap to be easily tuned in quaternary compounds. In general, the CBM decreases from sulfides to tellurides and from group III, to group IV, and to group V elements.396 Almost 30 different Cu chalcogenide compounds have been theoretically identified, which have suitable band gaps for solar energy conversion between 1.0 and 1.5 eV.396 Density functional theory (DFT) calculations have become key to predicting the properties of new complex materials and to unravelling the often hidden composition−structure−property relationships of materials which have already been exploited. In this regard, several interesting works have been published by Zhang et al. using the modified Becke−Johnson potential, plus an on-site Coulomb U (mBJ + U) approach, to determine the fundamental electronic properties of Cu chalcogenide compounds.166,168,389,396 Other reports have focused on DFT calculations of the electronic structure of certain Cu chalcogenides, specifically CuS;32 Cu2A (where A = S, Se, Te); 9 4 , 1 3 8 , 1 4 0 , 3 8 9 , 4 0 0 − 4 0 3 Cu 3 BiS 3 ; 3 1 0 CuSbA 2 and

√ √, √ √ √ x √ √, √, √, √ √, √ √

x

Ref

331−334 335,336 337−345 √ √ √

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Table 1. continued

Eg (eV)

Band alignment

Light absorption coefficient (cm−1)

Conductivity type

Exciton Bohr radius (nm) Applications

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Table 2. Relevant Atomic Levels of s, p and d States389 Atom Cu S Se Te

Electronegativity 1.9 2.5 2.4 2.1

Covalent radius (pm) 132 105 120 138

Valence states 3d104s1 3s23p4 4s24p4 5s25p4

s-states −4.61 −17.26 −17.29 −15.11

p-states

d-states −5.04

−7.01 −6.51 −5.95

Figure 8. Density of states for (a) ZnS, (b) CuGaS2, and (c) kesterite Cu2ZnSnS4. For CuGaS2, the Cu 3d states are projected on t2 and e suborbitals; the X-ray photoemission spectroscopy (XPS) data is also shown.168 The magnitudes of conduction states are amplified for clarity. The VBMs are at zero, and the shaded areas show the band gaps. The results are obtained from the DFT+U+G0W0 approach. Reprinted with permission from ref 168. Copyright 2015 Elsevier B.V.

CuBiA2;272,404−406 Cu2-IV-VI3;407−409 CuScS2;410 Cu3PSe4298; Cu3-V-VI4;411 CuFeS2;412,413 CuYS2;414 Cu-III-VI2;181,415 CuInA2;248,416−418 CuGaA2;417−420 CuGa1−xAlxS2;421 CuInGaSe 2 ; 422 Cu 2 ZnTiA 4 ; 423 Cu 2 CdSnSe 4 ; 357 CZTS and CZTSe;167,424−427 and CuFe2GaSe4.428 The determination of the electronic structures of Cu chalcogenides is still a matter of intense research. However, there is some controversy over the calculated electronic structures and the apparent inconsistencies between theoretical and experimental works, obtained from X-ray photoelectron spectroscopy (XPS) and X-ray emission spectroscopy (XES) investigations.138,429−431 This is not only due to the complexity of the DFT calculations, but it is also attributed to the numerous phases, off-stoichiometric compositions, and low energy of formation of defects in the Cu chalcogenide compounds. This variability translates into substantial experimental uncertainties in the measured band gaps of these materials, with a large scattering of band-gap values evident in the reported results. In particular, the large density of defects in these materials may result in changes in the band gap, where band tailing reduces the band gap and limits the performance of kesterite solar cells432 and the Burstein−Moss effect leads to an increase in the effective optical band gap in the Cu2−xA (where A = S, Se, Te) compounds. In spite of their chemical simplicity, the binary Cu chalcogenide compounds represent a particularly dramatic case in this direction because they show complex phase diagrams, with several nearly degenerated metastable phases and stoichiometries possible, and this results in significant complexity in the electronic structure.389,389,433 3.1.2. Defects. An important characteristic of Cu chalcogenides is their high defect concentration, which plays a key role in determining most of their properties, including charge and heat transport, optoelectronic, optic, and plasmonic properties. Cu chalcogenide compounds tolerate a large range

and density of nonstoichiometry and antisite defects, which may result in spatial compositional inhomogeneities and ordered defect compounds.434−437 Several reports have focused on a calculation of defects in Cu chalcogenides, including CuInSe2,165,263,438−446 Cu2SnS3,447 CuGaS2,419 CuSbSe2,448 and CZTS.234,244,439,449−451 While both anion and cation deficiencies are tolerated, inducing n- or p-type conductivities, Cu chalcogenides with p-type conductivity are the most common, where their p-type conductivity is associated with Cu vacancies. As an example, it takes much less energy to form a Cu vacancy in CuInSe2 than to form cation vacancies in the II-VI compounds,165 and the low formation energy of the shallow defect Cu vacancies provides CuInSe2 with p-type self-doping. Nevertheless, controlled stoichiometry can also provide CuInSe2 with n-type self-doping, which has allowed for CuIn(S1−xSex)2 homojunction solar cells to be produced.1,452,453 In terms of CZTS, very low formation energies have been computed for several defects234,244,439,449−451 and some of these have been experimentally demonstrated.240,454−457 Cu vacancies and Cu−Zn antisites are considered to be the main defects even in stoichiometric samples, which is in good correlation with the p-type character of this material.458 In the Cu-poor and Zn-rich conditions used to process CZTS solar cells, Cu vacancies and anti-site ZnCu defects dominate, in which their formation energy is actually negative.167,244,451 Further details on CZTS defects and the influence of stoichiometry can be found in previous reviews.167,459−461 Binary compounds display even more extreme behavior. In binary compounds, the crystal structure behaves like a solid backbone of chalcogen anions, which are surrounded by a liquid of mobile Cu+ cations.50 In this structure, Cu ions show a superionic behavior and a related highly disordered distribution of Cu ions.98 The large stoichiometry deviations of Cu2−xA (0.0 < x < 0.3) have very high associated carrier concentrations, 5882


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up to 1021 cm−3 for Cu2−xS,6,37 which offers exciting technological opportunities and enables the design of new processes to control doping.462 This high carrier concentration translates into an increase of the effective optical band gap through the Burstein−Moss effect, which, in combination with structural and electronic variations associated with compositional changes, results in an increase of the band gap of Cu2−xS with the Cu deficiency, specifically from 1.1 to 1.4 eV for Cu2S, 1.5 eV for Cu1.8S, and 2.0 eV for CuS.6,33,42 The Cu deficiencies and related high carrier density also provide a plasmonic resonance band in the NIR region, which finds several important applications, as discussed later in this review.141 On the other hand, the high defect density introduces large amounts of nonradiative recombination sites in the bulk and surface trap states, which prevents Cu2−xA compounds from displaying photoluminescence (PL). 3.1.3. Surface Defects/Composition. In polycrystalline layers and grain boundaries, defects play a key role in controlling transport properties.463−465 In multicomponent devices, interface defects may also control charge transfer between distinct materials.466 Nanocrystalline materials are particularly susceptive to these surface effects,467 which offers extraordinary possibilities to control functional properties by engineering the surface composition.128,468 For example, Draguta et al. demonstrated that, by manipulating the surface chemistry, it was possible to control the majority charge carriers and produce p- and n-type CuIn(S1−xSex)2 QD-based films.469 In addition, charge carrier mobilities of 0.03 cm2 V−1 s−1 for ethylenediamine (EDA) treated films were also demonstrated. Lee et al. studied the electronic properties of Cu2−xSe NC-based thin films, after treating the NCs with several short ligands including S2−, SCN−, and Cl−, and they demonstrated charge carrier mobilities of 1 cm2 V−1 s−1 with hole densities of 1021 cm−3.470 Using field effect measurements, Brewer and Arnold studied the electronic properties of Cu2S QD-based films, treated with ethanedithiol or EDT, and demonstrated carrier mobilities of up to 10−3 cm2 V−1 s−1 and hole densities of 1020 cm−3 for EDT-treated films.471 Aigner et al. studied the effect of hydrazine and EDT treatments on the electronic properties of Cu2−xS NCs and demonstrated a several orders of magnitude increase in the electrical conductivity of up to 10−2 S cm−1.468 Based on DFT calculations, Zhang et al. proposed the use of potassium as a surfactant to alter the CZTS surface, in order to reduce deep electron traps generated by Zn at Cu antisite defects and the formation of ZnS secondary phases at the interfaces between crystals.472 In this scenario, the characterization of the surface composition of Cu chalcogenide QDs has become fundamental to optimize their functional properties. In this regard, Dierick et al. analyzed the surface chemistry of colloidal CuInS2 (CIS) NCs using nuclear magnetic resonance spectroscopy (NMR).473 The authors determined that the CIS NCs were stabilized by tightly bound L-type amines, making them neutral and stoichiometric, which is somewhat surprising for a ternary material that is prone to off-stoichiometric compositions. The tightly bound amines were shown to be relatively difficult to displace, requiring several steps at high temperature to be replaced by thiols. In addition, the use of technical oleylamine (OLA) during the synthesis was demonstrated to result in nonstoichiometric NCs in this report, where the NCs had L-type ligands and X-type impurities on the surface that can be exchanged for carboxylic acids.

technological application. Furthermore, the complexity of the different crystal structures, combined with a large range of nonstoichiometric compositions, has resulted in a range of different photoluminescence (PL) mechanisms being proposed. In binary Cu chalcogenides, the band gap is dependent on the composition and crystal structure of Cu2−xS and Cu2−xTe and can vary from 1.1 to 1.5 eV.1,42,75,144,255,474 In the case of Cu2−xSe and Cu2−xS, a direct (indirect) band gap has been reported and can be tuned from 2.1 to 2.3 eV (1.2−1.4 eV) and from 1.8 to 2.9 eV (1.05−1.8 eV), respectively.253,475−477 In addition, Luther et al.37 and Hsu et al.478,479 reported that the NC size and shape can also influence the band gap. The most common strategy to alter the composition of these binary materials is to make them Cu poor, which results in a widening of the band gap. This can be attributed to the increased number of free carriers in the NCs, due to Cu vacancies. The increase in the band gap with the availability of free carriers in the valence band leads to a low energy tail in the absorption spectrum, which makes the exact determination of the band gap somewhat difficult. Luther et al.37 and Xie et al.75 also reported that an increase in free carrier concentrations (with the Cu deficiency) gives rise to a localized surface plasmon resonance (LSPR), which has a broad absorption band in the NIR region. This was further explored for the dampening of the plasmon resonance by increasing the Cu concentration in binary Cu chalcogenide NCs, where the peak shifted to lower energies with observable radiative recombination. In addition, Dorfs et al.480 and Wang et al.481 reported that the incorporation of other metal cations (i.e. Ce4+, In3+) led to a dampening of the LSPR and a noticeable shift of the emission peak to lower energies. In the case of ternary and quaternary Cu chalcogenides, these materials exhibit a tunable direct band gap that can range from the visible to NIR spectral region, which makes these materials highly suitable for solar energy conversion482,483 and bioimaging applications.484,485 In the ternary Cu chalcogenide system, CIS and CISe NCs are the most investigated compounds, in which their optical properties can vary with a change in size (Figure 9a)9 and composition,174 with significant off-stoichiometry possible in the NCs (Figure 9b−c). In addition, the exciton Bohr radii are reported to be 4.1 and 10.6 nm for CIS and CISe, respectively.9,486 CIS and CISe NCs have several appealing features for technological applications and are of particular interest, due to their absence of heavy metals (i.e. less toxic), relatively high quantum yields,487 large Stokes shifts (>0.5 eV),12 larger PL lifetimes (up to hundreds of ns),9,488 and the ability to tune the absorption and emission of these NCs over the visible and NIR spectral regions.489 These ternary NCs (CIS, CISe) mainly exhibit structureless absorption spectra, without a sharp absorption transition or defined excitonic peak, due to shape and size inhomogeneities within the studied NC samples.9 The high occurrence of surface defects in smaller CIS NCs (with high surface to volume ratio) mainly leads to PL (radiative and nonradiative decay), as well as plasmonic behavior.171,481 These surface defects ultimately led to low quantum yields because of nonradiative PL decays. As CIS can occur in different crystal structures (i.e. chalcopyrite, zinc blende, and wurtzite), it can exhibit a slightly different emission behavior for each phase and, hence, a different emission mechanism. In relation to the chalcopyrite and zinc blende phases, CIS NCs can exhibit cation ordering and disordering behavior,184,490,491 respectively. In chalcopyrite NCs, the unequal bond lengths (i.e. RCu−S ≠ RIn−S) cause intrinsic

3.2. Photoluminescence

The correlation between the structural and optical properties of Cu chalcogenide NCs plays a significant role in their end 5883


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Figure 9. (a) Size tunable PL emission (650 and 800 nm) from CIS NCs with sizes varying between 2.5 and 4.0 nm. Adapted with permission from ref 488. Copyright 2011 American Chemical Society. (b) Compositionally tunable PL of Cu2ZnSn(S1−xSex)4 NCs with varying values of x. Reproduced with permission from ref 492. Copyright 2013 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim. (c) Absorption and luminescence spectra of CIS NCs with different [Cu]/[In] molar ratios. Reproduced with permission from ref 174. Copyright 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

significant change was noted in the size of the NCs; thus, the widening of the band gap does not result from the quantum size effect. Uehara et al.175 and Chen et al.174 also observed that the emission increased with the cation (Cu) deficiency, which resulted in a radiative defect, and that this decrease in the Cu content caused a reduction in cation−anion interband repulsions, due to a relative decrease in the Cu d-orbital character.165,499 Li et al.488 and others487,500,501 reported an increase in the PL quantum yield and PL lifetime for core/shell structures of CIS/ZnS NCs. The presence of a protective layer (e.g. ZnS, CdS) or extra cation (e.g. Zn2+, Ga3+) has been also observed to increase the PL quantum yield in multinary (ternary/ quaternary) NCs.488,502 Donega et al. have given the simple explanation of the passivation of surface traps and confinement of photoexcited carriers to the core, thus protecting the NC from oxidation and decreasing nonradiative relaxation pathways.503 Reiss and co-workers suggested that a change in the composition of multinary NCs (via the introduction of another cation) can vary the donor and acceptor state levels by disturbing internal electronic states, and hence, the resultant emission peaks can be shifted.1 A mechanism was suggested by Akkerman et al. that zinc cation alloys at the NC surface improve the surface states by either decreasing or removing the nonradiative recombination routes, thus increasing the emission intensity from the core−shell structure.501 In the case of chalcopyrite CIS NCs, Pan et al.482 and Nam et al.504 have suggested that by increasing the concentration of the Zn cation, the emission peak can be shifted to higher energies. This was attributed to an increase in the band gap, due to surface alloying with Zn and a decrease in the size of the NC core. Zhang et al. have suggested the passivation of intrinsic defects, which resulted in a decrease in donor−acceptor recombination and an increase in the emission intensities.505 The reports by De Trizio et al.502 and Torimoto et al.506 have also shown that the formation of a ZnS shell increased the quantum yield of the multinary NCs, up to ∼75%. Li et al. reported that the presence of a CdS shell can increase the quantum yield by up to 86% but the addition of Cd ions increases the toxicity of the NCs.488 In the case of the ZnS-CIS alloy NC system, Zaiats et al. reported that band-gap excitation gives rise to low and high energy emission states.507 In particular, a radiative and nonradiative hole recombination occurs when photogenerated electrons are nonradiatively transferred to intraband-gap sites, which gives rise to different decay times. These decay times are dependent

defects and, hence, broad PL emissions, when compared to the zinc blende analogues. In addition, it is feasible to tune the composition over the visible and NIR range in zinc blende NCs due to cation disorder, as cations can be easily exchanged with one another in the cation sublattice.184 There is considerable debate in the literature about the exact mechanism of PL emission in chalcopyrite CIS NCs. For example, Castro et al. mainly attributed the PL emission to the presence of surface defects.493 However, a number of reports in the literature have recently suggested that the emission in chalcopyrite CIS NCs is more likely to be due to intrinsic defects. Li et. al. suggested that the emission originated from electron transitions to acceptor defect states near the valence band.488 It was later suggested by Castro et.al.,493 Chen et al.,174 and Kraatz et al.494 that the presence of a donor−acceptor pair was a plausible reason, due to their faster relaxation pathways, as compared to band edge recombination. An alternative hypothesis by Li et al.488 and Knowles et al.495,496 suggested localized hole recombination with the conduction band electron states. However, Omata et al.497 and Kraatz et al.494 discounted the hole localization and recombination as a reason for emission and, instead, attributed it to radiative recombination between localized electrons and valence band holes. It has been argued that zinc blende CIS NCs should not exhibit emission, as defects cannot exist in a cation disordered structure.487 However, as colloidal NC solutions exhibit PL emission, it has been suggested that there is always a small proportion of chalcopyrite NCs, with ordered cation states, in any quantity of zinc blende NCs. In terms of wurtzite CIS NCs, they were determined to exhibit a similar PL emission behavior to chalcopyrite NCs, including broad emission peaks, large Stokes shift, and long PL lifetimes.498 In addition, the emission peaks of wurtzite NCs are slightly lower in energy, compared to chalcopyrite NCs. However, the PL emission mechanism has not been fully established for wurtzite CIS NCs, due to the relatively low number of literature reports. In general, the band gap can be adjusted in these systems by adjusting the cation ratios. For example, Uehara et al. reported that the composition of chalcopyrite CIS NCs could be tuned by adjusting the concentration of the Cu precursor in the NC synthesis.175 In this report, the authors observed that the emission peak was shifted to a higher energy (Figure 9c) by decreasing the cation ratio (Cu/In) in the synthesis. A corresponding increase in the band gap was also observed and no 5884


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intrinsic defects rather than surface defects. In addition, control of the emission intensities, peak position, and quantum yield was achieved by manipulating the Cu/Zn cation ratios in the NCs. In a separate report, Pan and co-workers reported the PL properties of Cu-Mn-In-S (CMIS) alloyed NCs with a ZnS shell.511 The authors observed that the quantum yield was as high as 20%, due to the reduction of nonradiative recombination centers, with a multiexponential emission lifetime decay evident with longer decay times. The longer lifetime decays were attributed to the initially populated core-state recombination, surface-related radiative recombination, and donor− acceptor transition recombination. Hua et al. studied core CIS and CMIS NCs (without any shell formation) and observed that the PL quantum yield decreased with Mn ion incorporation, as Mn2+ created nonradiative recombination centers and defect states in the NCs.512 In addition, a temperaturedependent PL study was conducted on CIS/ZnS NCs, and it was observed that the emission intensities decreased in the temperature range from 93 to 270 K, which was attributed to donor−acceptor recombinations of the carriers. From the calculation of the thermal activation energy for the system (200 meV), the authors concluded that a nonradiative recombination process (with the thermal activation energy) resulted in a decrease in the PL intensity with an increase in the temperature. In the CMIS/ZnS NC system, no significant effect of an increase in the temperature on the PL intensity feature was observed. It was concluded that the relatively large thermal activation energy for the nonradiative recombination process in this system provided thermally stable emission behavior in the temperature range. Zhou et.al studied a similar system with Mn-doped ZnCuInS NCs and investigated the energy level of different sized NCs and PL behaviors, as a function of temperature in the range 20−300 K.513 A decrease in the emission intensity with an increase in the temperature was also observed, and this was attributed to the low activation energy of the nonradiative relaxation process, which caused a decrease in the PL intensity, similar to the conclusion made by Hua et al.512 In the Cu2ZnSn(SxSe1−x)4 (CZTSSe) system, the possibility to tune the optical properties by varying the NC size is negligible, due to their very small exciton Bohr radius (∼2.5 nm).514 A more versatile route to tune the band gap in these materials is by controlling the stoichiometry. In this regard, most of the work has been done by simply tuning the anionic (i.e S/Se) ratio, while keeping the cation ratio constant, which leads to band-gap tunablities from 0.9 to 1.5 eV.492,515 Raadik and coworkers studied the low-temperature PL properties in CZTSSe monograin powders in solution form, in order to understand the radiative recombination process at 10 K.516 In this report, an asymmetric emission band was observed for the CZTSSe, which shifted toward lower wavelengths and became more asymmetric with increasing S content. This observation was associated with the presence of a high concentration of charge defects in this multinary system, where a high number of charge defects causes potential fluctuations and a local disturbance of the band structure, thus forming band tails and broadening the defect level distribution. Similar observations were also reported by Altosaar and co-workers for CZTSe monograin powders.517 Grossberg et al. studied the emission behavior of CZTS polycrystals at 10 K and proposed that the two PL emission bands at 1.27 and 1.35 eV arose from a bandto-impurity recombination process involving deep acceptor defects in the crystals, where the deep defects could be related to CuSn.518 Lin et al. studied the emission properties in

on the Zn and Cu ratios in the NC system, and at a Zn/Cu ratio of 1, a resolvable PL emission is observed that arises from two different emitting states. In addition, these studies showed that the high energy band from Zn does not change with varying cation ratios. This further supports the fact that the recombination of photogenerated electrons occurs near the conduction band, where emissive states reside. Furthermore, a number of emission peaks and broad absorption features have been observed in more recent reports, which are attributed to the presence of intraband-gap states in CIS NCs. For example, Manna and co-workers observed this phenomenon in their study of NCs (∼3 nm) with decreasing Cu ratios in the samples.502 The authors argued that the change in optical properties was not only due to quantum confinement effects, but also from the decrease in valence band edge energy, which resulted in an increase in the band gap. In addition, the large number of Cu vacancies is a probable cause for the strong p-type self-doping, which was evident in this system. Zhang et al.165 and Chattopadhyay et al.508 had earlier coined the idea of p-type, self-doping in CIS NCs, which explained the multipeak emission behavior in this NC system. The increase in the emission peak at longer wavelengths and the PL quantum yield have both been attributed to Cu vacancy recombination in the conduction band, as well as to a donor−acceptor recombination mechanism. Depending on the vacancy of Cu atoms in a NC, the quantum yield has been reported to increase by ∼7% in stoichiometric CIS NCs and up to ∼23% in nonstoichiometric CIS NCs. This observation was further supported from the emission lifetime of CIS NCs, where a slower emission decay contributed to both vacancy recombinations in the conduction band, as well as donor−acceptor recombination, whereas a fast emission decay could be from non-radiative recombinations. Over time, different reports have argued various reasons and mechanisms for the low PL emissive behavior of the CIS NCs. Reiss and co-workers studied the emission properties of the CIS/ZnS core/shell NCs to understand the evolution of emission in the NCs, by using time-resolved PL spectroscopy and different time delays.509 With different time delays, the authors observed two red shifted emission bands (at 650 nm and at 720 nm) with different decay times. The authors stated that the decay times for both spectral components (∼280 ns and 370 ns, respectively) were relatively long, which indicated that the emission bands originated from donor−acceptor pair recombinations rather than from excitonic-related emission. This behavior is a consequence of the contribution of the Coulomb energy, which decreases with increasing separation distance between the donors and acceptors. The lattice defects in the CIS NC system cause fluctuations to local carriers, and this ultimately results in long emission decays for the transitions. Maeda and co-workers also studied CIS NCs and the alloyed ZnCuInS (ZCIS) NC system to understand the emission.510 They reported that the PL quantum yield for the ZCIS NCs was ∼5%, which was higher than that reported for the CIS system. The authors attributed this finding to the alloying of ZnS with CIS, where the emission arose from donor−acceptor pair recombinations. By further overcoating the ZCIS NC with a ZnS shell, an enhancement of the PL emission intensity was achieved, due to the elimination the surface defects. However, no reduction in the full width half maximum (FWHM) of the PL band shape or in the large Stokes shifts occurred, which indicated that the PL emission originating from the donor−acceptor transition was due to 5885


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CZTSSe thin films and observed a single broad band emission, which was attributed to quasi donor−acceptor transitions and defect related recombinations.519 Singh et al. reported the first low temperature emission study on CZTSSe NCs, in which the emission could be tuned from 1.4 to 0.9 eV (Figure 9b) by systematically varying the S/Se ratio in the NCs.492 This resulted in a defined shift in the PL spectra, which correlated well with the NC stoichiometric composition. In addition, a broad emission band was observed, which was suggested to originate from near band-edge emission. Overall, multicomponent Cu chalcogenide NCs exhibit broadened PL bands with FWHMs in the range 80−120 nm, larger Stokes shifts (0.2−0.5 eV), and an extremely slow luminescence decay life (up to 500 ns), compared to the band edge emission of other metal chalcogenide NCs systems (i.e. II-VI, IV-VI).1,5,12,488,520,521 These results are attributed to various types of intrinsic defects (lattice vacancies or other point defects) that are present in Cu chalcogenide systems, which not only leads to possibilities of compositional inhomogeneities within NC samples but also consolidates many kinds of defect-based, nonradiative decay pathways in their luminescence.5,12,174,522,523

and Cu2−xTe141,255,543 and the related alloy NCs.546−548 Some recent tutorial reviews have already highlighted nonstoichiometric, Cu chalcogenide NCs as promising materials for tunable plasmonics.6,15,549−555 Importantly, the plasmonic behavior in Cu chalcogenides is the result of the collective oscillation of free holes, opposite to free electrons in metals. The fact that Cu chalcogenide NCs exhibit LSPR in the NIR region has sparked extensive research for applications where strong NIR absorbance is required, such as in photothermal therapy and photoacoustic tomography.477,556−558 Furthermore, it serves as a lower cost alternative to the conventional plasmonic noble metal NCs (i.e. Au) in biomedical applications.559 Cu chalcogenide NCs are regarded as “self-doped materials”, where the free carriers arise from a variation in the oxidation state of generally one of the elements in nonstoichiometric phases.15 This compensates for the substoichiometry of one of the components and results in holes in an otherwise filled valence band. In Cu chalcogenide NCs, the intrinsic formation of cation vacancies arises from a deficiency of Cu, which results in these NCs being p-doped. The concentration of free holes is set by the stoichiometry of the NCs. In particular, copper sulfide in its stoichiometric (Cu2S) form does not support LSPR due to the absence of free holes. However, upon increasing the Cu deficiency to form non-stoichiometric (Cu2−xS) compositions, the number of free carrier (holes) in the valence band increases, and this causes the plasmonic resonance peak to shift to higher frequencies (i.e. a blue shift).549 Figure 10 depicts the correlation between the charge carrier density and the LSPR frequency (either static or dynamic), which can be tuned by controlled doping of semiconductor NCs. This two-dimensional plot is based on the Drude model and allows the LSPR frequency to be extrapolated, as a function of the NC diameter (y-axis) and the number of free carriers (contour lines from x-axis), which highlights the wide range of LSPR frequencies that are assessable via control of the carrier density.37 While metal NCs (i.e. Au, Ag) have high carrier densities (∼1023 cm−3), it is difficult to change their carrier density and, thus, allow their LSPR frequency to be tuned outside the visible and NIR windows. However, in semiconductor NCs, carrier densities can be varied from 1016 to ∼1021 cm−3, and this allows their LSPR to be tuned from the terahertz (THz) regime up to NIR frequencies. Considerable efforts has been devoted to manipulating the LSPR frequency in Cu chalcogenide NCs by changing the NC size, shape, and composition. Despite many earlier papers on the synthesis of Cu2−xS NCs, it was Zhao et al.’s work in 2009 that started the field of plasmonic Cu chalcogenide NCs.42 In this work, the authors proposed that the NIR LSPR originated from free holes, which resulted from the presence of cation vacancies due to the Cu deficiency in the NCs. Three different chemical methods (sonoelectrochemical, hydrothermal, and solventless thermolysis methods) were used to synthesize Cu2−xS NCs in this report,42 with adjustment of the reduction potential, adjustment of the pH, and different precursor pretreatments, respectively, resulting in compositional variations from Cu1.97S to CuS (Figure 11a). The authors consistently found that Cu1.97S (djurleite) is more stable than Cu2S (chalcocite) under ambient conditions. In 2011, Luther et al. followed this up with a combined experimental and theoretical study, which conclusively demonstrated that the NIR absorbance is due to the plasmon resonance of free holes in Cu2−xS NCs.37 In this particular study, Cu1.94S NCs of 2.5-6 nm

3.3. Plasmonics

Plasmonic NCs are an interesting class of materials with potential uses in several fields, ranging from biological imaging and laser photothermal therapy,19,524−527 to sensors,528−533 to photonics,534 to improved photovoltaic devices.535−538 Their plasmonic properties originate from the collective oscillations of free charge carriers, which are in turn related to the surface plasmon resonances that are said to be localized in NCs.15,539 This gives rise to the so-called localized surface plasmon resonance (LSPR), where surface plasmons provide the opportunity to confine light to very small dimensions. LSPR is an optical phenomenon, which arises from the resonant interaction between electron-charged oscillations near the surface of metal NCs and the electromagnetic field of light.534,539 This strong light−matter interaction generates a locally enhanced electromagnetic field, where the LSPR either can radiate light (Mie scattering), for use in optical and imaging fields, or can be rapidly converted to heat (absorption) for therapeutic applications such as photothermal therapy.524 It is generally observed in NCs for sizes greater than a few nanometers (∼2−12 nm) and smaller than one-fifth of the LSPR wavelength.15 Over the past two decades, there has been a dramatic acceleration in progress in this field.540−542 In particular, noble metal NCs (i.e. Au, Ag) are the most systematically investigated plasmonic materials and support LSPRs in the visible spectrum.541,542 However, LSPRs are not only limited to metal NCs and can be achieved in semiconductor NCs with appreciable free carrier concentrations.37 The discovery of LSPR in Cu chalcogenide semiconductor NCs has created new opportunities for the manipulation of light, where the LSPR frequency of Cu chalcogenide NCs can be tuned by intrinsic doping via tailoring the stoichiometry of the NC. This provides Cu chalcogenide NCs with a particularly attractive property: a composition-dependent LSPR that can be tuned to support LSPR in the NIR region. This is arguably one of the strongest points in favor of Cu chalcogenide NCs over metal NCs, which are restricted to a fixed carrier density.15 In the past six years, LSPR has been observed in numerous nonstoichiometric Cu chalcogenide NCs such as Cu2−xS,37,42,543,544 Cu2−xSe,480,545 5886


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Figure 10. Dependence of the LSPR frequency on the density of free carriers. Reproduced with permission from ref 37. Copyright 2011 Macmillan Publishers Ltd: Nature Materials.

by charge transfer, leading to a red-shift and a weakening of the LSPR.562 Variation of the Cu precursor content is another way to vary the LSPR energy and the free carrier concentration. For example, by reducing the amount of the CuCl precursor in the one-pot synthesis from 0.5 to 0.25 mmol, the LSPR for Cu2−xS NCs red-shifted from 1290 to 1540 nm, respectively, and thus by a total shift of 250 nm.563 Dorfs et al. demonstrated that the exposure of Cu2−xSe NCs to either an oxygen or a cerium(IV) complex, (NH4)2Ce(NO3)6, resulted in the development of an NIR plasmon band due to the creation of Cu vacancies in the material.480 Upon controlled oxidation, the authors observed a blue shift in the optical absorbance, which was attributed to the increased concentration of Cu vacancies. This is in agreement with the findings from Luther et al.37 The as-synthesized Cu1.96Se NCs (no plasmon absorption, broad peak at 1700 nm) were gradually oxidized to Cu1.81Se NCs (narrow plasmon peak at 1100 nm) under ambient conditions (Figure 11d).480 Furthermore, the authors showed that the addition of an excess of Cu+ ions (Cu(I) salt) led to a reduction of the LSPR band in the NIR region and a well-defined red-shift of the plasmon absorption, with each subsequent addition of Cu(I) (Figure 11e). The same sample was oxidized again under ambient conditions for 19 h in total, thus proving that the process was reversible (Figure 11f). No changes in the crystal structure (cubic berzelianite) were observed, despite the creation and annihilation of numerous Cu vacancies. Kriegel et al. extended these studies to the whole family of binary Cu chalcogenide NCs (Cu2−xS, Cu2−xSe, and Cu2−xTe).543 Similarly to Dorfs et al.,480 oxygen exposure led to a gradual transformation from their stoichiometric forms (Cu2−xS, Cu2−xSe, x = 0) into their nonstoichiometric counterparts (Cu2−xS, Cu2−xSe, x > 0), which resulted in an intense LSPR in the NIR region after the transformation. However, changes in the crystal structure were observed after the oxidation process, in that Cu2S (chalcocite) transformed into Cu1.97S (djurliete) and Cu2Se transformed into Cu1.8Se. The charge carrier density extracted from the plasmon frequency gave values of 1.4 × 1021 cm−3 (for Cu1.97S), 3 × 1021 cm−3 (for Cu1.8Se), and 5 × 1021 cm−3 (for Cu2Te). In this particular report, the authors pointed out that the crystal phase of the Cu2−xTe NCs could not be unambiguously distinguished between rickardite Cu1.4Te and stoichiometric

in diameter were prepared, which had a free hole density of 1021 cm−3, and this corresponded to an LSPR energy in the NIR region at 0.7 eV. In comparison, the most common plasmonic metals (i.e Au, Ag, Cu) have free electron densities in the range 1022−1023 cm−3, with corresponding LSPRs in the visible region.560 Interestingly, the LSPR and excitonic transitions were observed in the same NCs, both of which were shifted to higher energies when the number of Cu vacancies was increased (Figure 11b).37 The authors attributed the shift in the LSPR energy to the increased number of free carriers. Moreover, the shift of the excitonic peak was caused by the Burstein−Moss effect, whereby the free carriers occupied the top of the Cu(I)S valence band and this led to an increase in the optical band gap. The LSPR energy in these materials can be tuned by several means, including the use of different surface ligands545,547 in the synthesis to vary the NC stoichiometry and, thus, the free carrier concentration; by variation of the aspect ratio in anisotropic structures;478,561 and by controlled chemical posttreatments such as oxidation and reduction.480,543 In particular, oxidation increases the number of free carriers and the degree of Cu deficiency and can be carried out under ambient conditions in air or through the use of an oxidizing agent such as cerium(IV) ammonium nitrate, (NH4)2Ce(NO3)6.480 As a consequence of oxidation, an NIR band emerges and is blueshifted in energy, depending on the extent of the Cu deficiency. In contrast, reduction decreases the number of free carriers and the Cu deficiency and can be carried out with reducing agents such as a Cu(I) salt, tetrakis(acetonitrile)copper(I) hexafluorophosphate (Cu(CH3CN)4PF6), or a Cu-free reducing agent such as diisobutylaluminium hydride (DIBAH).543 For example, Swihart and co-workers demonstrated that the stoichiometry and LSPR of Cu2−xS NCs could be tuned by varying the amount of oleic acid (OA) used in the synthesis.545 The authors explained that OA coordinates to the NC surface via its deprotonated carboxyl functional group and carries a negative charge. This coordination may trap holes and, thereby, reduce the effective free carrier concentration. The LSPR absorbance red-shifted by up to 270 nm in Cu2−xSe NCs and by up to 110 nm in Cu2−xS NCs (Figure 11c). In analogy with these findings, Jain et al. showed how common ligands bound to the surface of Cu2−xS NCs can electronically dope the NCs 5887


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Figure 11. (a) Extinction spectra of Cu2−xS NCs in the NIR region, with compositional variations from Cu1.97S to CuS. Reproduced with permission from ref 42. Copyright 2009 American Chemical Society. (b) Absorption spectra of Cu2−xS NCs, where both plasmonic and excitonic features are evident in the spectra. Upon oxygen exposure, the plasmon resonance becomes more pronounced due to the increased number of Cu vacancies. Reproduced with permission from ref 37. Copyright 2011 Macmillan Publishers Ltd: Nature Materials. (c) Absorption spectra of Cu2−xS NCs showing the dependence of LSPR frequency on the ligand combination used in the synthesis. Reproduced with permission from ref 545. Copyright 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim. (d) Time-dependent evolution of the LSPR in Cu1.96Se NCs upon exposure to air. A gradual blue-shift of the LSPR occurs over time due to their oxidation to Cu1.81Se. (e) Tuning the LSPR through stepwise reduction using a Cu(I) salt, (Cu(CH3CN)4PF6) which reduces the cation vacancy concentration. A red-shift in the LSPR is evident due to the reduction of Cu1.81Se to Cu1.96Se. (f) Reversible tuning of the LSPR by gradual oxidation in air where the oxidized particles again exhibited an LSPR. Panels d−f were reproduced with permission from ref 480. Copyright 2011 American Chemical Society. (g) Evolution of the NIR extinction spectra of Cu2−xSe NCs over time upon the addition of a copper-free reducing agent, DIBAH, where a gradual red-shift and decrease in the intensity of the NIR LSPR is observed. The inset shows the extinction spectrum of a nonoxidized sample. Reproduced with permission from ref 543. Copyright 2012 American Chemical Society. (h) Shape-dependent LSPRs for spherical NCs and nanodisks, with two LSPR modes observed for the nanodisks. Schematics of the LSPR polarizations for the nanodisks are included, where the out-of-plane mode for the nanodisks occurs at lower energy (1600−1900 nm) and the in-plane mode is observed only at higher energies (>3100 nm). Adapted with permission from ref 478. Copyright 2011 American Chemical Society. (i) UV-vis spectra of Cu2−xTe nanocubes, nanoplates, and nanorods. Reproduced with permission from ref 255. Copyright 2013 American Chemical Society.

Cu2Te. Another reducing agent, DIBAH, was used instead of Cu+, and similarly, this led to a gradual red shift and a decrease in the intensity of the LSPR band (Figure 11g) in the NIR region.142,543 The crystal structure changed during the reduction process, and this was attributed to a reduction of the Cu2+ species to Cu+ inside the NCs, which was triggered by chemically injected electrons.543 The authors concluded that Cu(II) did not leave the NCs upon oxidation and, instead, stayed on the surface of the oxidized Cu chalcogenide NCs, presumably in the form of CuO or as a monolayer of Cu(II) atoms bound to surface ligands. This result indicated that Cu ions could be re-inserted into the lattice from the surface layers, rather than from an external Cu-based reducing agent.

The extent to which changes in the NC composition affected the optical response was systematically studied by Xie et al., who developed an approach to access Cu2−xS NCs in several stoichiometries, starting from Cu1.1S up to Cu2S NCs.75 This approach involved a postsynthetic reaction of the as-synthesized Cu1.1S NCs with a Cu(I) complex, tetrakis(acetonitrile)copper(I) hexafluorophosphate, at room temperature to achieve a stepwise increase in the Cu stoichiometry. The Cu1.1S NCs exhibited a well-defined NIR absorption peak at 1090 nm, ascribable to an in-plane LSPR mode, with increasing amounts of the Cu(I) complex inducing a red-shift and a decrease in the intensity of the NIR absorption band, until a faint absorption band was observed at 1250 nm. 5888


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LSPR-expressing CIS NCs, which were suitable for photovoltaic and photocatalytic applications due to their modifiable surface chemistries and their LSPR peak stability. While synthetic methods to CIS NCs are prevalent in the literature, these methods generally result in NCs with charge neutral stoichiometries that resemble those of the precursors, which consistently lack plasmon modes.184,488,568−571 The plasmonic CIS NCs obtained by Niezgoda et al. were synthesized using a highly reactive sulfur source, bis(trimethylsilyl) sulfide (TMSS), and it was hypothesized that the combination of extremely labile S2− anions with the stable hexadecylamine (HDA)-capped cation species resulted in a hectic chemical environment, whereby sulfur planes were formed in the NC at a rate that restricted the complete filling of the cationic sites.567 This resulted in stable HDA-capped CIS NC with inherent cation vacancies, that had tunable band gaps from 1.5 to 2.1 eV and LSPR modes in the NIR region at 1200 nm. Based on their LSPR modes, the plasmonic CIS NCs were consequently tested in semiconductor sensitized solar cells (SSSCs), and they showed an 11.5% relative efficiency enhancement, compared to their nonplasmonic counterparts.572 This enhancement was attributed to augmented charge excitation, stemming from near-field antenna effects in the plasmonic-CIS-based SSSCs. A subsequent report by Wang et al. showed that plasmonic CIS NCs could also be formed using a more environmentally friendly sulfur source (i.e. sulfur powder/oleic acid),481 as opposed to the TMSS sulfur precursor used by Niezgoda et al. in ref 567. The plasmonic CIS NCs synthesized by Wang et al. were found to have LSPR responses that were strongly affected by the In content, where CIS NCs with the lowest detectable level of In incorporation exhibited an NIR LSPR peak at 1355 nm and NCs with the highest In cation fractions (47%) resulted in a red-shift of the absorbance and a LSPR peak at 1533 nm.481 This red-shift was accompanied by a significant broadening in the LSPR peak and an eventual disappearance of the LSPR, which was attributed to a decrease in the free carrier concentration (Cu vacancies) with an increased In content. A few reports of Cu-IV-VI NCs (i.e. CuxSnyS and CuxSnySe) with LSPR responses have been presented. For example, Liu et al. observed an LSPR response for Cu2SnS3 NCs in the NIR region and found that the LSPR peak was strongly influenced by the Cu:Sn elemental ratio.573 This demonstrated that a change in the NC composition produced a corresponding change in the carrier density, as reflected in the dampening of the LSPR at 1326 nm for low Sn content, followed by the eventual disappearance of the LSPR and a red-shift in the absorbance with an increased Sn content. LSPR modes have also been observed in Cu3SnS4 NCs, where an increase in the Cu:Sn ratio caused a red-shift and weakened LSPR bands, which was attributed to an increase in CuSn defects that were detrimental for LSPR.574 A similar trend was observed for Cu2SnSe3 NCs, where the LSPR decreased in intensity and redshifted with increasing Sn content.575 Cu-poor CZTS NCs were also studied, but no LSPR was observed, presumably due to the lower hole concentration within the range of ∼1015 to 1018cm−3, which may be too small to support an LSPR mode.574 Plasmonic properties have also been observed in the Cu-V-VI category of NCs, in particular for Cu3SbSe3 nanorods302 and Cu3BiS3 NCs.576,577 The ternary Cu3SbSe3 nanorods were found to display dual absorptions, one in the visible region from band-edge absorption and the second in the NIR region which was attributed to LSPR.302 It was noted that the LSPR peak red-shifted (from ∼1360−1490 nm) with an increase in

A similar in-plane LSPR mode was previously observed in CuS disks.561 On the contrary, Liu et al. synthesized digenite Cu1.8S NCs and discovered that the NCs underwent in situ phase transformation from Cu1.8S to CuS upon keeping the resulting colloidal solution in a vial for a few days.564 This transformation occurred as a result of an OLA-assisted oxidation process, and the LSPR related absorption spectra exhibited an obvious transition with the phase transformation. Shape can also affect the plasmonic properties. For example, Tao et al. reported the observation of shape dependent LSPRs for Cu2−xS nanodisks.478,479,565 The resonances were split into two modes: a transverse (out-of-plane, higher energy) mode and a longitudinal (in-plane, lower energy) mode, and these LSPR modes could be modulated by varying both the aspect ratio of the disk and the overall dopant concentration of Cu2−xS. Two asymmetric plasmonic bands peaking in the NIR (at 1600−1900 nm, out-of-plane mode) and the mid-IR (at >3100 nm, in-plane mode) were observed (Figure 11h).478 These out-of-plane and in-plane features were assigned by fitting the extinction spectra to a simple analytical expression based on an approximate NC shape and the Drude model. The authors reported that the LSPRs were blue-shifted as the aspect ratio of the nanodisks was increased and attributed this to the higher free carrier densities in NCs with higher aspect ratios. Some other reports on the Cu2−xS nanoplates have only been observed an in-plane LSPR mode.75,559,561 However, not all shapes have been shown to affect the plasmonic properties. For example, Li et al. synthesized Cu2−xTe nanocubes and nanoplates with an LSPR around 900 nm, but no LSPR was observed for Cu2−xTe nanorods (Figure 11i).255 The authors suggested that the small transversal dimension of the thin nanorods may not support a detectable plasmon. Another report investigated the LSPR response from more elaborate shapes such as tetrapods and found that they only exhibited a weak longitudinal LSPR.141 Most of the theoretical considerations are based on the application of the Drude model (i.e. assuming a free carrier, metal-like behavior for Cu chalcogenide NCs) to the particular morphology under investigation.37,42,478−480,566 However, it has been pointed out that this approach may only be reasonably adequate for metal NCs, and thus, it has to be taken with due reservation when applied to doped semiconductors, where carriers supporting LSPR may not behave as fully delocalized quantum oscillators.141,554 Analysis of the ultrafast LSPR dynamics has confirmed that a remarkably lower free carrier density can be accommodated in Cu chalcogenide NCs, compared to that found in noble metal counterparts.543,551,566 Following on from the binary Cu chalcogenides, ternary Cu2−xSySe1−y NCs have also been prepared, which had tunable plasmonic properties. For example, Dilena et al. developed a hot-injection method to prepare Cu2−xSySe1−y NCs and found that the LSPR energy depended mainly on the degree of the Cu deficiency (x) and only slightly on the crystalline phase or on the S/Se concentration (y).546 A more recent study on Cu2−xSySe1−y NCs found that the LSPR could be tuned between 1000 and 1600 nm by varying the amount of OA used in the synthesis.547 In addition to Cu2−xSySe1−y NCs, other ternary Cu chalcogenides have also shown plasmonic properties, which greatly expands the library of NC systems that support LSPRs. For example, Niezgoda et al. showed that quantum confined CuxInyS2 (CIS) NCs exhibited a strong LSPR in the NIR region.567 This work represented an important advance in the field of plasmonics, in that it was the first report of 5889


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field and an improved optical response.583 Another reason is the dependence of the plasmonic NCs on the dielectric of the material and its surroundings. This dielectric sensitivity can be used to manipulate the light interactions and modify the plasmonic response, by inducing a nonlinear change in the dielectric of the material with a controlled beam. The femtosecond response time in plasmonic excitation provides another advantage, as this paves the way to allow for ultrafast processing of optical signals. Typically, NCs embedded in inorganic glasses (as matrices) have been studied, due to several advantages over the organic film matrices, such as better optical quality and better energy damage threshold. 584−586 The favorable low structural dimensionality and the ability for electronic structure modulation in nanometer scale materials, such as QDs and wells, allow for an increase of light−matter interactions, which is important for the nonlinear optical properties. Metal chalcogenides exhibit promising nonlinear optical properties, due to their large energy band gaps and the occurrence of emission over a wide electromagnetic spectrum range, coupled with a strong quantum confinement of excitons in low dimensional structures. The partial orbital screening in semiconductor NCs, along with residual electron−hole interactions and interfacial charge trapping, highly impacts the nonlinear optical properties and electron dynamics. Nonlinear optical materials can produce coherent light by frequency conversions in the ranges where standard lasers are unavailable or perform poorly. In general, when two different frequency beams, f1 and f 2, are passed through a nonlinear material, the frequencies interact nonlinearly and this gives rise to four distinct output frequencies: 2f1 and 2f 2 by second harmonic generation, along with a sum and difference frequency generation (f1 ± f 2), respectively. This allows the range of laser applications to be expanded, by easily converting laser radiation from one frequency to another. Cu chalcogenides are finding specific interest in the field of nonlinear optics because different pathways exist to change the optical properties of the NCs, due to their multivalency, different stoichiometric forms, plasmonic response, and chargetransfer transitions. For example, in Cu2-xSe NCs, the surface layer can change at the interface due to the multivalency of Cu and its reaction with available oxygen, and this can significantly affect the optical response of the NCs. It has also been demonstrated that the NIR absorption band in Cu2-xSe NCs can be effectively bleached with laser pulses.584 Furthermore, the optical response of different NCs with similar band gaps, but with a different position of the absorption band onset, can be characterized by investigating their bleaching relaxation times and their ground state absorption cross sections. In the case of Cu2-xSe NCs, the bleaching and relaxation mechanisms have been explained by taking into consideration the inter-band energy levels within the unfavorable gap of the NCs. For example, Alexeenko and co-workers have shown that when silica glasses were doped with Cu2-xSe NCs, different bleaching effects and relaxation times (0.3−2.7 ns range) were observed for different NC stoichiometries.585,587 By increasing the Cu/Se ratios in the NCs, a decrease in the bleaching time and an increase in the peak absorption cross-section were observed. The authors explained this behavior by associating it with a change in the position of the inter-band energy levels, which gave rise to the absorption bands at 1 and 1.5 mm in these NCs. They also observed transient bleaching, with a time constant of absorption recovery of approximately 300 ps and 1.3 ns for the Cu2Se and Cu2−xSe doped glasses, respectively.

the refractive index of the solvent, and this allowed the authors to confirm that the absorption of the Cu3SbSe3 nanorods was due to LSPR. This was the first report on Cu-V-VI nanorods with LSPR activity,302 compared to previously reported Cu3SbSe3 NCs (ref 303) which did not facilitate LSPR. The authors assumed that the origin of the LSPR may be related to the exclusive Cu−Se bonding within the Cu3SbSe3 crystal structure and deduced that Cu3SbSe3 behaved similarly to Cu2Se, where the LSPR is known to originate from Cu vacancies. Cu3BiS3 NCs are another interesting material set that have been shown to exhibit strong LSPR in the NIR region.576 The Cu3BiS3 NCs exhibited an LSPR peak at 876 nm, which underwent a red-shift to 919 nm with the increasing refractive index of the solvent, as is expected for NIR absorption in Cu-deficient chalcogenide NCs. The LSPR peak was shown to be tunable in this report, where the NIR absorption was strongly affected by the Cu:Bi precursor ratio and blue-shifted with increasing Cu deficiencies. The intrinsic plasmonic properties in the Cu3BiS3 NCs were proven to be particularly useful for photothermal therapy (PTT) and computed tomography (CT) contrast enhancement, where the NCs were successfully demonstrated as a novel photothermal theragnosis agent for CT/IR thermal imaging and PTT. The CT imaging response of the Cu3BiS3 NCs was similar to that of Bi2S3 NCs, due to the large X-ray attenuation coefficient of Bi. While the LSPR peak was particularly strong in this report, a weaker and broader absorption peak was observed in a separate report on Cu3BiS3 NCs at 1150 nm, which was attributed to a LSPR mode that was induced by cation vacancies in the NCs.577 Nonetheless, it is evident that preparing Cu-III-VI, Cu-IV-VI, and Cu-V-VI NC compositions away from the stoichiometric compounds provides new opportunities to tune the charge carrier concentration in self-doped ternary NCs. Regardless of the synthesis route employed, the resultant LSPR modes are intrinsic to the NC stoichiometry. Therefore, it is crucial to ensure that control is instigated over the NC stoichiometry in the synthesis stage, as this regulates the concentration of free charge carriers and the resultant LSPR response. While Cu2−xS NCs are the most extensively employed Cu chalcogenide NC composition in bioapplications, particularly for imaging and photothermal therapy, ternary NCs may be even more stable to oxidative effects from air, than their binary NC counterparts. Undoubtedly, the added levels of customization, which can be achieved through the incorporation of different group III, IV, and V cations, make both ternary and quaternary Cu chalcogenide NC systems worthy of further investigation. 3.4. Nonlinear Optics

Nonlinear optical materials have attracted considerable research interest, due to their better response time and control over the frequency spectrum of laser light. Correlators, saturable absorbers, ultrashort-pulsed laser generators, and real-time holography are a few examples of optical device applications that have the potential for optical computing.543,578,579 In terms of Cu chalcogenide compounds, Cu2−xS and Cu2−xSe NCs have found application as nonlinear optical materials because the plasmonic effect enhances their inherent weak optical nonlinear response, which arises due to photon−photon interactions in the system.580−582 The nonlinear optical properties of Cu2−xS and Cu2−xSe NCs are enhanced due to several factors, namely due to coupling interactions between the surface plasmon features and light, which produce a strong local electromagnetic 5890


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carrier density in these NCs. This results in strong nonlinear effects at high fluences, both during the carrier temperature relaxation process and during the pump absorption process. Kriegel et al. demonstrated that the optical properties can be manipulated by tuning the Cu deficiency in Cu2−xS, Cu2−xSe, and Cu2−xTe NCs under oxidative/reductive conditions.543 Furthermore, Valle et al. also investigated the ultrafast nonlinear dynamics of Cu2−xSe NCs using pump−probe femtosecond resolution.582 Philip and co-workers investigated the nonlinear optical properties of CuS NCs, with femtosecond resolution, using the ultrafast open-aperture Z-scan technique and a regeneratively amplified Ti:sapphire laser for excitation.581 They discussed the possibility of free carrier absorption, multiphoton absorption, and saturable absorption, arising from the ultrafast laser pulses, as a possible pathway for the optical nonlinearity in the NCs. They also argued that the formation of surface traps and states in/around the band edge in the NCs can contribute toward the nonlinear absorption and that smaller NCs can have higher nonlinear absorption coefficients, which makes them better emitters. Recently, Qiu and co-workers have shown that, by using solution-processed Cu2−xS NCs, the ultrafast pulse can be generated and optically modulated.594 In this report, ultrafast pulse generation in the mid-IR and NIR region was due to the nonlinearity of the NC in that region, where superbroad band saturation absorption (>400 THz), ultrafast recovery (∼315 fs), and a large modulation depth (∼7 dB) were observed. From these reports, it is evident that the absorption and nonlinear optical response can be modulated, depending on the NC phase and surface properties. The nonlinearity in the optical features of NCs can be explained by a change in dielectric functions, which are induced by an increase in carriers and lattice temperatures. Cu chalcogenide NCs exhibit strong interband transitions in the visible region in their stoichiometric forms and a weaker absorption edge onset in the NIR and midIR region. The oxidation of these NCs leads to a transformation from their stoichiometric to nonstoichiometric forms, resulting in the evolution of a strong surface plasmon band in the NIR region. This optical nonlinearity has been observed for Cu2−xS (x > 0) and Cu2−xSe (x > 0) NCs in this spectral region, respectively. The enhanced nonlinearity of these plasmonic NCs contributes toward a utilization of their optical properties with reduced power in photonic applications, while the nanometer size range opens up new avenues to develop integrated photonic devices. In addition, the fast response time of plasmonic excitation, from the ultrafast nonlinearity of free electrons, improves the optical signal and permits workable femtosecond response times.

In an earlier work, Malyarevich et al. demonstrated that the short relaxation (in the picosecond and nanosecond range) of the induced absorption effect, under picosecond excitations, made the silica sol-gel glasses of CuS and CIS NC systems promising as fast laser darkening filters.586 In both systems, the authors observed that the induced absorption controlled the bleaching of the optical response. Klimov et al. studied the linear and nonlinear properties of CuxS NCs (∼8 nm) in glasses containing CdS QDs and reported a strong enhancement of 2-3 orders of magnitude in the third order of nonlinearity, when compared to glasses with CdS QDs alone.588 They also observed a high-energy shift in the lowest optical transition when the Cu ratios were changed in the NCs and attributed the increased nonlinearity to an increase in the linear absorption and the carrier lifetime.589 The femtosecond transient absorption studies of CuxS NCs were reported in a separate study, which indicated different nonlinear optical properties, depending upon whether the NCs formed direct or indirect band gaps.590 Depending on the Cu deficiency in the NCs, nonlinear transmission followed the mechanism of either state-filling-induced bleaching for CuxS (where x = 1.8, 1.9, and 1.96) or photoinduced absorption (where x = 2). Artemyev and co-workers studied the optical transient bleaching and absorption in CIS NCs in their oxidized form.591 They found that the additional absorption band, which appeared in the CIS NCs upon oxidation, bleached under picosecond excitation and that long bleaching (> 300 ps) and absorption features arose, due to the trapping of mid-gap surface states that were produced upon oxidation. Sumiyama and co-workers reported both the linear and nonlinear optical properties of chalcopyrite CIS NCs (∼2 nm).592 They concluded that surface defects played an important role in the nonlinear optical properties and that the Stark effect was the mechanism for the nonlinear optical properties, due to the presence of electrons and holes on the surface defects of the NCs. Klimov et al. suggested that the formation of strong local fields, from the separation of surface trapped electrons and holes, can result in bleaching at the original transition energy by shifting the transition energy toward lower energies.593 They also suggested the saturation of absorption as another mechanism for the nonlinearities in these NC systems, due to the state-filling effect, which is an exchange interaction between excitons and a screening of the Coulomb interaction between electrons and holes. In addition, the authors stated that this charge separation phenomenon has been observed in both glass and colloidal samples. In glass samples, charge separation occurs at high pump excitations and is further explained in terms of Auger related trapping at interface states, whereas, in colloidal systems, the trapping of holes in isoenergetic states can lead to charge separation. The nonlinear plasmon studies for Cu2−xS, Cu2−xSe, and Cu2−xTe (x > 0) compounds were conducted by pump−probe techniques in separate studies by Scotognella et al.566 and by Kriegel et al.543 This technique involves the investigation of the transient optical response of a system after exposing it to an intense pump pulse. Scotognella et al. studied Cu2-xSe NCs (∼13 nm) via ultrafast laser pulse in a pump−probe system, with femtosecond resolution, to understand the dynamics of localized plasmon resonance exhibited by Cu2-xSe NCs.566 The results revealed that the nonlinearities in the plasmonic absorption of the NCs were due to their lower carrier density. The authors argued that the smaller carrier heat capacity and the higher effective carrier temperature resulted from the low

3.5. Magnetism

In general, Cu chalcogenide NCs (e.g. CIS, CZTS) do not show any response to an applied magnetic field. However, it is noteworthy to highlight that the Cu(II) ion (3d9) is considered to have magnetic character, owing to the presence of an unpaired electron in its 3d shell.595,596 This results in Cu2+ exhibiting paramagnetic behavior, whereas Cu+ exhibits diamagnetic behavior because it has a completely filled 3d shell (i.e. 3d10). This differentiation has important consequences in terms of their magnetic properties because Cu+ has a negative magnetic susceptibility; thus, its internal induced magnetic field will oppose an externally applied magnetic field. Conversely, Cu2+ is paramagnetic and exhibits a small, positive magnetic susceptibility to an applied field. In terms of magnetic resonance imaging 5891


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80 K in CuFeSe2 was observed in a separate report, in which the Fe atoms on the 2e and 2a sites (of the tetragonal cell) have slightly different magnetic moments, thus resulting in a weak ferromagnetic behavior.618 The magnetic properties in chromium-based chalcogenide spinels (chalcospinels) have also been studied. In particular, CuCr2S4, CuCr2Se4, and CuCr2Te4 are ferromagnetic metals at temperatures close to room temperature, with Curie temperatures of 377, 430, and 360 K, respectively.619,620 The interest in this system originates from their sufficiently high Curie temperature (TC), which is needed for the room-temperature operation of spin-based electronic devices.621 While the CuCr2A4 (A = S, Se, Te) bulk materials show ferromagnetic character at room temperature, this character may be lost when the crystal size is reduced to the nanoscale due to size effects, but it could be recovered at lower temperatures in some cases. In this regard, Ramasamy et al. synthesized CuCr2S4 NCs and nanoclusters to investigate their magnetic properties.330 It was noted that both the NCs and nanoclusters showed superparamagnetic behavior at room temperature, whereas ferromagnetic character was observed at 5 K. In addition, relatively low coercivity values and magnetization values around 30 emu/g were observed,330 which were still lower than those of the bulk counterpart (86 emu/g).622 The magnetic properties of CoxCu1−xCr2S4 NCs were also investigated by Ramasamy et al. in a separate report, in which the effects of cobalt substitution on the saturation magnetization and coercivity values were evaluated.623 The CoxCu1−xCr2S4 NCs showed superparamagnetic behavior (for x values up to 0.6) at 200 K, whereas paramagnetic behavior was displayed for samples which had higher Co content under the same temperature. In addition, the NCs behaved as soft ferromagnets (x ≤ 0.6) at 5 K, and superparamagnetic behavior was observed for samples with higher Co content. It was noted that the NCs with higher Co content (i.e. Co0.8Cu0.2Cr2S4 and CoCr2S4) were much smaller than the other compositions;623 thus, the observed drop in magnetization could also be associated with a size effect.615 The selenium analog, CuCr2Se4, also displays a ferromagnetic behavior624 and has a pronounced magneto-optic Kerr effect625 at room temperature. In relation to NCs, slightly different magnetic activities were reported, depending on the NC synthetic method which was employed. For example, magnetic measuremernts of CuCr2Se4 NCs (synthesized from microwave radiation) showed a saturation magnetization of 15 emu g −1 and a coercive force of 80 Oe at room temperature.332−334 In another report, agglomerated CuCr2Se4 NCs (particle sizes between 25 and 200 nm) were formed in a solvothemal synthesis approach, in which a Curie temperature of 450 K presented saturation magnetization of 2.3 μB per Cr atom at 5 K.626 Similar results were also reported by Rao et al.627 A colloidal approach reported by Wang et al. allowed for much better size control of the resultant CuCr2Se4 NCs, in which hysteresis loops for NCs, with an average size of 15 and 25 nm, were measured at 300 and 10 K, respectively.628 In this report, both samples were observed to be superparamagnetic at room temperature but exhibited relatively low coercivity at 10 K, with saturation magnetization values of 37 emu/g and 43 emu/g being reported for the 15 and 25 nm sized NCs, respectively. In addition, ferromagnetic behavior was only evident below the blocking temperature of ∼250 K, as deduced from the lowfield (50 Oe) measurements. A colloidal approach was also used by Lin et al. to synthesize CuCr2Se4 hexagonal nanoplates

(MRI), the most commonly used contrast agents to enhance MRI are paramagnetic materials. A recent study by Xie et al. made an important advance in determining the oxidation state of Cu in Cu2−xS NCs and concluded that the sole oxidation state of (nonsurface) Cu in all Cu2−xS NCs is Cu+, regardless of the NC stoichiometry, which infers that Cu2−xS NCs should only display diamagnetic behavior.75 This data was also corroborated with superconducting quantum interference device (SQUID) and electron paramagnetic resonance (EPR) measurements, which can measure extremely subtle magnetic fields and, thus, can differentiate Cu+ (diamagnetic) from Cu2+ (paramagnetic). Contrary to this conclusion, Mou et al. demonstrated that Cu2−xS NCs provided a contrast enhancement for T1-weighted MRI and detected the presence of Cu2+ and Cu+ in the NCs via X-ray photoelectron spectroscopy (XPS) measurements.597 The success of their experiment was based on the longitudinal (r1 =) relaxation enhancement of water protons, induced by Cu2−xS NCs. In addition, a brightening of phantom images containing Cu2−xS NCs was observed. However, the authors noted that the relaxivity observed in these Cu2−xS NCs was quite modest, especially when compared to more common magnetic ions, such as Gd3+. Other possibilities to provide Cu chalcogenide NCs with magnetic character involve: (i) NC-labeling or their conjugation with magnetic molecules/ complexes;598−601 (ii) doping or alloying with magnetic ions;602−607 or (iii) forming heterostructures601,608−611 (dimers, core/shell structure), in which one component of the heterostructure is a magnetic NC. In addition, the compositional versatility in Cu chalcogenides also allows for the incorporation of magnetic elements (such as Fe, Ni, Co, Mn, and Cr) to form ternary or quaternary compounds that exhibit magnetic properties. The two most investigated Cu chalcogenide compounds with magnetic character are CuFeS2 and the chromium-based compounds CuCr2A4 (A = S, Se, Te). The mineral chalcopyrite (also known as CuFeS2) has been reported to exhibit antiferromagnetic behavior with a Neel temperature of 823 K, with the average magnetic moment of Fe atoms reported to be around 1.75 μB in the bulk.612 This is strikingly different from the magnetic moment of Fe3+ calculated in CuFeS2 NCs, which was determined to be around 0.12 μB at 4 K and was indicative of a very weak ferromagnetic behavior.613 This difference in the magnetic moment was associated with a size reduction to the nanoscale, surface effects on small NCs and interparticle interactions.614 In NCs, a reduction in size leads to an increase in the number of spins on the surface, which are higher than the number of spins forming the NC core. In addition, defects at cation and/or anion sites and non-fully-coordinated atoms at the NC surface, with resulting uncompensated spins, can also reduce the net magnetization.615 The particular phenomena behind the magnetic properties of NCs have been discussed elsewhere.616 A report by Wang et al. demonstrated that CuFeSe2 NCs displayed interesting magnetic properties at different temperatures.613 For example, magnetic measurements revealed that CuFeSe2 NCs exhibited ferromagnetic and paramagnetic behavior at 4 and 300 K, respectively. A sharp increase in the magnetic coercivity from 39 Oe at 300 K to 1780 Oe at 4 K was observed, which the authors suggested may be due to the reduced thermal fluctuation of magnetic dipoles. Previous reports have observed paramagnetic behavior from 300 K to ∼71 K in CuFeSe2, with weak ferromagnetic behavior observed below this temperature.617 A magnetic transition at 5892


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(15−24 nm) and flower-shaped nanostructures (19.8 nm).629 Magnetic measurements revealed that both the nanoplates and flower-shaped nanostructures exhibited ferromagnetic behavior below 350 K. A comparison between the two structures revealed that the flower-like nanostructures showed higher saturation magnetization, coercivity, and remanent magnetization higher than the hexagonal nanoplates, despite presenting lower Curie temperatures (380 K vs 420−430 K for the hexagonal nanoplates). This difference demonstrated that the magnetic properties of the CuCr2Se4 NCs were also strongly dependent on the NC morphology and/or surface anisotropy. Ramasamy et al. investigated the magnetic properties of CuCr2S4−xSex (0 ≤ x ≤ 4) NCs, which could be synthesized over the entire composition range.630 In this report, ferromagnetic ordering over the entire compositional range was revealed, with a systematic increase in the superparamagnetic blocking and Curie temperatures observed with increasing Se content in the NCs. In addition, the saturation magnetization and coercivity values of the CuCr2S4−xSex (0 ≤ x ≤ 4) NCs at 5 K were found to progressively rise up to x = 3.630 The magnetic properties of CuCr2Te4 NCs have also been studied, with magnetic measurements indicating ferromagnetic behavior at 5 K and near-superparamagnetic behavior at 300 K.335 A Curie temperature of 330 K was reported for the CuCr2Te4 NCs, and the saturated hysteresis loop obtained at 5 K showed a saturation magnetization value of 24 emu g−1. While exceptional spin-dependent properties and ferromagnetic ordering at room temperature are evident in the bulk CuCr2A4 (A = S, Se, Te) compounds, there is still a lack of suitable solution methods to synthesize these materials in NC form, which prevents further exploration of their magnetic properties. Undoubtedly, the development of synthetic routes to form size-, shape-, and phase-controlled CuCr2A4 (A = S, Se, Te) NCs, CuFeS2 NCs, and other Cu chalcogenide NCs (with Fe, Ni, Co, Mn, and Cr) would allow for a better understanding of the magnetic and electronic properties of these materials at the nanoscale.

for a higher degree of size and shape control because the respective NC nucleation and growth stages are separated in time. In the discussion of the colloidal synthesis method (section 4.1), further divisions are made within this section to separately discuss the mechanism behind NC nucleation and growth according to the La Mer classical theory (section 4.1.1), size focusing/defocusing concepts (4.1.2), and newer considerations in nucleation (4.1.3). A separate subsection on the importance of the hard-soft-acid−base (HSAB) theory is provided in section 4.1.4, which provides a qualitative understanding on how to regulate the reactivity of the different precursors in colloidal methods. Since the NC composition is intrinsically related to the functional properties of the material, the HSAB theory is particularly important for ternary and quaternary NC compositions to ensure that all the desired elements are present in the resultant NCs. 4.1. Colloidal Synthesis Method

The colloidal synthesis method is an oxygen-free, organic phase synthetic protocol that involves the thermal decomposition of organic precursors, typically inorganic salts or organometallic compounds, in a high boiling point organic solvent.631−635 The precursors are in turn chemically transformed into active atomic or molecular species (monomers) upon heating, thus providing the essential ingredients to form colloidal NCs.631,634,636,637 The colloidal synthesis method has been widely used to synthesize NCs because of its many advantages, such as the high crystallinity and monodispersity of the NCs and the high dispersion ability in organic solvents. This method generally involves several consecutive stages: (i) nucleation from an initially homogeneous solution; (ii) growth of the preformed nuclei; (iii) isolation of particles from the reaction mixture; and (iv) postsynthetic washing treatments to remove unreacted precursors and ligands from the NC solution. The colloidal setup involves a three-neck, round-bottom flask that is connected to a condenser and a Schlenk line to provide an oxygen-free environment for the synthesis. Typically, the reaction flask is evacuated at temperatures around 100 °C for 30−60 min, before the reaction temperature is further elevated, with the purpose of eliminating adventitious water and dissolved oxygen, which would otherwise affect air-sensitive or pyrophoric precursors in the reaction. There are two techniques in the colloidal synthesis method, which use homogeneous nucleation to synthesize NCs; the “heat-up” and “hot-injection” techniques, both of which have already been reviewed.8,638,639 The “heat-up” technique (Figure 12a) is a batch process, in which the precursors, ligands, and solvents are present in the flask initially and are mixed at low temperature, followed by heating up to the desired reaction temperature to initiate NC nucleation and growth. This technique is attractive and advantageous for largescale production because of its simplicity and it allows for gram scale quantities of NCs to be achieved.635,640,641 The nucleation period in the heat-up method is much longer in the hotinjection method, as pointed out by van Embden et al., due to the progressive generation of monomers with an increased supply of thermal energy.8 Thus, even after the nucleation process is initiated, a high supersaturation is maintained for an extended period, but the undefined separation in the nucleation and growth stages ultimately leads to a broadened size distribution in the heat-up technique. The “hot-injection” technique (Figure 12b) was pioneered by the Bawendi group in 1993 to synthesize monodisperse Cd

4. SOLUTION SYNTHESIS APPROACHES Numerous solution synthesis approaches have been employed for the synthesis of Cu chalcogenide NCs. In this section, we will focus on the five most prevalent solution synthesis routes in the literature: (i) the colloidal synthesis method (section 4.1); (ii) solvothermal and hydrothermal methods (section 4.2); (iii) template-directed synthesis methods (section 4.3); (iv) the nanoscale Kirkendall effect-induced method (section 4.4); and (v) the cation exchange method (section 4.5). While the cation exchange method is not a direct synthesis approach, it is discussed within this solution synthesis category because it allows for novel NC shapes and compositions to be achieved via the reaction of as-synthesized NCs with a desired cation solution. This permits the exchange of cations within the NC lattice to form novel NCs shapes and compositions, some of which have proven difficult to achieve through direct synthesis approaches. Of all the solution synthesis approaches, the colloidal method is undoubtedly the most extensively used method to synthesize Cu chalcogenide NCs, as it allows for exquisite control over the size, shape, and crystal phase of the resultant NCs. While the colloidal method encompasses both the heat-up and hotinjection protocols, the heat-up approach is arguably a much more attractive route in terms of up-scaling the yield and production of NCs. However, the hot-injection approach allows 5893


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preparation of uniform colloidal particles was initially published in 1950 by LaMer and Dinegar, in which the concept of “burst nucleation” was proposed.652 This pioneering concept was developed from their research on the preparation of monodisperse sulfur hydrosols, which were achieved from the decomposition of sodium thiosulfate in hydrochloric acid. In this particular work, the authors explained that the sulfur sol evolved in the solution in a two-step process, where free sulfur was formed from the thiosulfate in the first step and sulfur sols formed in solution in the second step. In this process, many nuclei were generated at the same time and these nuclei continued to grow without additional nucleation events occurring. This is the essence of the “burst nucleation” process, which makes it possible to control the size of an ensemble of particles during the growth stage. Burst nucleation has been adopted as an important concept in the synthesis of monodisperse NCs, whereby the formation of monodisperse NCs requires a temporally discrete nucleation event, followed by slower controlled growth on the existing nuclei.635,636 Thus, the requirements for monodispersity can be summarized as follows: (1) a high nucleation rate leading to the burst of nuclei formation in a short period of time; (2) an initial fast rate of growth to rapidly reduce the concentration below the critical nucleation threshold; followed by (3) an eventual slow growth rate leading to a long growth period. The mechanisms of nucleation and growth of NCs in solution have been discussed in separate review articles.653 The LaMer plot is particularly useful to visualize how the energy barrier works to induce burst nucleation.635,636 For NCs to nucleate and grow, free monomers are required in the solution. However, at low temperatures, the monomers are essentially “locked” into the precursors, in that they either are bound to ligands, that prohibit their reactivity, or are part of a complex, that needs to be thermally decomposed to allow the monomers to be released or free to participate in the reaction.8 These monomers eventually react to form the atomic units that comprise the final NCs, provided there is sufficient thermal driving force and a high free monomer concentration in the reaction, so much so that the solution is said to be supersaturated with free monomers. This supersaturated state can be achieved either by directly dissolving the solute at higher temperatures, followed by rapid cooling, or by adding the reactants in such a way as to produce a supersaturated solution during the reaction. In the hot-injection approach, the rapid addition of precursors, by means of an injection in a hot coordinating solvent, raises the precursor concentration above the nucleation threshold for a brief period of time, and it is the subsequent temperature drop (after the injection) that results in a supersaturation of particles in the solution.634,636,638,654 The nucleation and growth process through the LaMer mechanism can be essentially divided into three phases, marked as I-III in the plot in Figure 13.653,655 In phase I, a rapid increase in the concentration of free monomers (or atomic concentration) occurs in solution with time, as the precursor is decomposed by heating.655 Since the energy barrier for the initiation of a nucleation event is considerably high, the monomers cannot spontaneously condense into nuclei at the saturation concentration (CS); thus, the monomer concentration continues to increase until it reaches the minimum nucleation concentration (denoted as Cnu,min in Figure 13). In phase II, the monomer concentration reaches its minimum nucleation concentration (Cnu,min). At this point, the system is said to be supersaturated. It contains enough energy to overcome the energy barrier656

Figure 12. Schematic illustration of the colloidal heat-up and hotinjection techniques. (a) Heating-up technique. All the precursors are mixed together in the initial stages and are steadily heated up to the desired reaction temperature. Reproduced with permission from ref 641. Copyright 2007 American Chemical Society. (b) Hot-injection technique, where a rapid injection of a precursor solution into a hot solvent leads to a temporal separation of the nucleation and growth stages. Adapted with permission from ref 635. Copyright 2007 WileyVCH Verlag GmbH & Co. KGaA., Weinheim.

chalcogenide NC, whereby a cold trioctylphosphine (TOP) solution, containing the Cd and chalcogen precursor, was injected into a hot trioctylphosphine oxide (TOPO) solution.634 In this technique, a high degree of supersaturation is induced upon injection, which leads to instantaneous burst nucleation by relieving the excess free energy of the supersaturation.635,638 Although the heat-up and hot-injection techniques both involve an initial burst of nucleation followed by growth on the existing nuclei, better size uniformity is achieved in the hotinjection approach because the injection results in a single nucleation event, leading to uniform NC growth.634,636,637,642−646 An important characteristic of the hot-injection approach is its flexibility and its general applicability to produce NCs of various materials.639 Since the first report by the Bawendi group,634 the hot-injection approach has not only been extended to the synthesis of other Cd chalcogenides NCs but also metals,647 metal oxides,648 and metal chalcogenide seminconductor NCs649 by using appropriate combinations of reactive precursors, ligands, and solvents. These synthetic protocols all share a common point, in that the induction of rapid crystal formation by hot-injection yields uniform NCs. In terms of the reaction temperature, the conventional synthesis of Cd chalcogenide NCs is typically carried out at 250−300 °C to form a cationic complex between the Cd precursors and phosphine/phosphonate ligands. This differs from the Cu chalcogenide NC synthesis, where it can be performed at lower temperatures of 180−250 °C due to easier formation of a cationic complex. Furthermore, Cu chalcogenide NCs can be synthesized in the absence of phosphine ligands by using alkylamines or alkylthiols, and these “phosphine-free” synthesis methods253,273,545,650,651 are attractive, from both a safety and a cost point of view, because phosphine/phosphonate ligands are generally expensive and pyrophoric. 4.1.1. LaMer Classical Nucleation Theory. A clear understanding of the NC nucleation and growth mechanism is crucial to develop colloidal synthetic methods that are applicable to a range of semiconductor NCs. Research on the 5894


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Once the monomer concentration is sufficiently depleted, growth proceeds by Ostwald ripening. In this process, a “defocusing” of the size distribution occurs, whereby smaller particles dissolve due to their higher solubility and surface energy and, consequently, larger particles grow by receiving monomers from the dissolving particles.635−637,642 As a result, the supersaturation is kept constant at a low level and no additional homogeneous nucleation can occur during the growth stage. This is the classic concept of Ostwald ripening. Two size regimes exist at this point, with the critical nuclei size between the extremes of a smaller and larger particle. The defocused size distribution is difficult to revert back to a focused one, unless the growth time is extended to such a point that the smaller nuclei are completely depleted in solution. It is still possible to recover a monodisperse sample after the synthesis by using “size-selective precipitation”, a separation technique which involves the stepwise addition of a hydrophilic solvent to a stable NC dispersion in a nonpolar solvent to gradually reduce the solvation power and allow for precipitation.636 The larger NCs precipitate out first because of their greater van der Waals attraction, and NC fractions with narrow size distributions can be obtained, but it can be quite tedious and time-consuming and yield small quantities of material.636 4.1.3. Newer Considerations in Nucleation. Recently, a deeper understanding of nucleation and growth in NCs has been developed through a focus on precursor conversion reactivity rates. In particular, the research groups of Owen,660 Hens,661 and Bawendi662 have proposed kinetic models to describe the correlation between the precursor conversion rate and the NC size of II-VI and IV-VI compounds. For example, Owen and co-workers demonstrated that the conversion kinetics of thiourea governed the extent of crystal nucleation, where an increased thiourea conversion reactivity produced a higher NC concentration and a smaller final NC diameter at full conversion.660 The general applicability of this approach was impressive, in that control over the monomer supply kinetics allowed for nucleation to be adjusted in a range of metal chalcogenide NCs, such as PbS, CdS, ZnS, and Cu2−xS. However, the authors observed that this model did not fit exactly with the La Mer classical nucleation theory, in that a subcritical dependence was observed between the precursor conversion rate and the NC concentration, which still requires more accurate nucleation models to explain. Importantly, this work highlighted that control over the precursor conversion reaction kinetics played a central role in tuning the NC size, while maximizing reaction yield and minimizing the size distribution of NCs. It also greatly simplifies the synthetic considerations because tailoring the precursor reactivity allows for greater predictability of NC nucleation and growth, rather than using the conventional trial and error experimentation approach, where reaction temperature, time, solvent, and ligand concentration are typically adjusted to regulate NC nucleation. 4.1.4. Hard-Soft-Acid−Base (HSAB) Theory. The key to forming single-phase, ternary, and quaternary NCs is ensuring that the reactivity of the constituent precursors is balanced otherwise, the formation of two separate material compositions is encountered. While the use of the hard-soft-acid−base (HSAB) theory663 was originally developed for molecular coordination compounds, it can be extended to NCs and can be used to identify the ideal precursors to regulate the reaction rate.664 This theory remained relatively dormant for precursor selection in binary II-VI NCs, but it gained huge recognition in the synthesis of multicomponent Cu-III-VI2 and Cu2-II-IV-VI4

Figure 13. LaMer model describing the three different phases of the nucleation and growth process for monodisperse NCs, as a function of the reaction time and monomer concentration. The saturation concentration (CS) and minimum nucleation concentration (Cnu,min) are marked in the plot. Adapted with permission from ref 655, Copyright 2009 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim, and with permission from ref 652, Copyright 1950 American Chemical Society.

and undergoes “burst-nucleation”, where the atoms start to aggregate into small clusters (i.e. nuclei) via self- (or homogeneous) nucleation. These nuclei grow at an accelerated rate, which partially relieves the supersaturation, and this leads to a significant reduction in the concentration of free monomers in solution. No additional nucleation events can occur in solution once the monomer concentration drops below the level of the minimum nucleation concentration. In phase III, the nuclei grow into NCs of increasingly larger size, where growth is controlled by the diffusion of monomers via ongoing precursor decomposition, and this growth process occurs slightly over the saturation point (S) because this is a less energy consuming process. 4.1.2. Size Focusing/Defocusing Concepts. A limitation of the LaMer theory is that the evolution of the size distribution during the growth stage is not predicted. While the LaMer model describes the process of nucleation followed by the growth of stable nuclei, the characteristics of the growth stage remain more or less unspecified. Soon after LaMer’s work, Reiss et al. developed a growth by diffusion model and the concept of focusing the size distribution, whereby smaller crystals will grow more rapidly than larger crystals if the monomer concentration is high.657 The nucleation threshold defines the outcome of the nuclei. In essence, nuclei larger than the critical size will further decrease their free energy for growth and form stable nuclei that grow, while nuclei smaller than the critical size threshold will shrink.636,637 The uniformity of the size distribution is largely determined by the time period over which the nuclei are formed and begin to grow. If the nucleation period is short, the smaller particles grow more rapidly than larger particles, as their free energy driving force is larger, resulting in a “focusing” of the size distribution.642 However, competing processes such as aggregation or Ostwald ripening that result in broadened size distributions were not considered in these theoretical studies, but an accepted model was later developed by Lifshitz−Slyozov−Wagner.658,659 5895


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NCs, due to the stringent demands to carefully select precursors in which their relative reactivity can be precisely tailored.489,665−667 It provides a qualitative understanding through the basis that the affinity between metal centers (Lewis acids) and coordinating ligands (Lewis bases) will be higher if they share similar “hard” or “soft” properties.8 According to this theory, proposed by Pearson in 1963,663 the tendency of an elemental ion to either donate or accept an electron leads to them being categorized as behaving as an acid or a base, respectively. They are further classed as possessing hard, borderline, or soft character, based on their atomic/ionic radius, oxidation state, polarizability, and electronegativity. Hard acids prefer to bind to hard bases and have an affinity for each other which is ionic in nature.668 Soft acids and soft bases similarly form tight bounds and have an affinity for each other that is mainly covalent in nature. On the other hand, hard-soft or soft-hard combinations have weaker binding/ interactions. In general, precursors with a weak metal−ligand bond are ideal for nucleation, as they increase the availability of free monomers in the solution, and this facilitates rapid nucleation. However, ligands need to strongly coordinate to the NC surface during the growth phase, in order to prevent uncontrolled growth and form a stabilizing monolayer, but they need to allow monomers in solution to diffuse to the NC surface at the same time. Thus, careful selection of the metal precursors and ligands is required to allow this delicate balance to be achieved. Hard acids tend to have small atomic/ionic radii and high oxidation states.669−671 Some examples of hard acids, relevant to the materials covered in this review, include group III ions (In3+ , Ga3+, Al3+), group IV ions (Sn4+, Si4+), as well as the first row transition metal ions (Mn2+, Fe3+, and Co3+). On the other hand, soft acids tend to have large atomic/ionic radii and low oxidation states, such as Cu+, Ag+, Au+, Hg+, Cd2+. There is also a “borderline” character category, where their binding interaction lies in between the hard and soft character, and Cu2+, Zn2+, Sn2+, Sb3+, Co2+, and Ni2+ feature in this category. In terms of bases, the most well-known hard bases include Cl−, OAc−, acac−, and NO3− (all of which are notable counterions on metal salt precursors), but also OH−, H2O, NH3, and R-NH2. The category of soft bases includes I−, S(Se,Te)2−, thiols (R-SH, the most well-known thiol being dodecanethiol, DDT), and phosphine ligands (R3P), such as trioctylphosphine (TOP) and tributylphosphine (TBP), to name but a few. The hard and soft classifications of various metals and ligands used in the synthesis of NCs are illustrated in Figure 14. A classic example of where the HSAB theory proves extremely useful is in the synthesis of ternary CuInS2 (CIS) NCs,665 where copper(I) iodide is selected for its propensity to decrease the reactivity of the Cu+ ion with the thiol (sulfur precursor). Copper(I) iodide (CuI) is the most tightly bonded Cu precursor, as opposed to copper chloride (CuCl) or copper acetate (CuAc). Specifically, Cu+ is a soft acid and I− is a soft base, and so, tight binding is expected to result in its decreased overall precursor reactivity. Thus, the Cu+ ion reacts slowly with the thiol sulfur source (soft base, with I− > thiol), suppressing Cu2−xS growth, which permits the formation of ternary CIS NCs. This leads to the attainment of new crystal phases that are generally not stable in the bulk form, such as the wurtzite and zinc blende modifications. While the HSAB-based generalization of ligand exchange is very useful for predicting the binding affinity between a given ligand and NC composition, one should understand that it has some degree of limited

Figure 14. Hard and soft classifications of various metals and ligands used in the synthesis of NCs.

predictive power and does not necessarily dictate all the possible ligand exchange pathways.664 4.2. Solvothermal/Hydrothermal Method

The solvothermal synthesis method has aroused considerable attention among research groups, for its ability to access increased precursor reactivity and solubility under elevated temperatures and pressures.22,672 It is regarded as a modification to the conventional organic-solvent-based synthetic methods, which have been widely employed for the synthesis of monodisperse NCs since the early 1990s.634 The difference is that the solvothermal process is carried out in a specialized sealed reaction vessel called an autoclave, which consists of an inner Teflon liner, an outer stainless steel shell, and a stainless steel cap. All these components in the autoclave are capable of withstanding high temperature and pressure environments, which are created in the vessel over prolonged periods of time.672,673 This setup is illustrated in Figure 15a−b, in which an inert Teflon liner is used to protect the outer stainless steel shell from corrosive reagents employed in the synthesis, as well as avoiding any contamination from the steel shell. Autoclaves can be used up to temperatures of 270 °C, and depending on the engineering specification of the steel walls, pressures of 150 bar can be withheld.672 When the reactants and solvents are heated under high pressure in the autoclave, the reactants experience considerable increases in their solubility and reactivity, ultimately speeding up the reaction, which cannot normally occur at standard atmospheric pressure conditions. This results in the subsequent crystallization of the dissolved material from the solvent. The pressure in the reaction is self-generated and not only depends on the temperature but also relies on other experimental factors, such as the percentage fill of the vessel and the quantity of dissolved salts.672,674 While traditional solution-based routes are limited by the boiling point of the respective solvent, the solvothermal method possesses the added advantage that it can withstand temperatures up to the critical point of the solvent, provided that they are safe for the vessel. This merit makes the solvothermal method very flexible and attractive in designing high-quality inorganic NCs in this reaction system. 5896


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hydrophilic surface functionalization. Metal chloride salts have proven to be the precursors of choice in both solvothermal and hydrothermal methods, with a range of chalcogen precursors being used, such as thiourea, sodium sulfide, carbon disulfide, thioacetamide, and elemental sulfur or selenium powder. Different shapes, such as flower-like nanostructures,677−684 2D nanosheets,391,685 and 3D hollow spheres,678,686−691 can be produced by this method, typically with sub-micron dimensions. However, this method faces some drawbacks, such as the use of long reaction times, broad NC size distributions, NC aggregation, and the impossibility of studying in situ reactions, due to its closed reaction system. 4.3. Template-Directed Synthesis Method

The template-directed synthesis method has proven to be an efficient method for the preparation of well-distributed, onedimensional nanomaterials, such as sub-micron nanotubes and nanowires, or for the preparation of hollow structures with defined shapes. The template provides the necessary structural framework, within or around which the nanostructure forms and grows into a shape that is complementary to the template.686,692,693 Generally, templates employed in the synthesis can be divided into two categories: (1) hard templates, which are either used as physical scaffolds for deposition of the desired material; and (2) soft templates, which contain ligands, surfactants, and polymers.22 An important feature of the template method is its large versatility with respect to the diameter and length of the pores, in that the outer diameter of nanotubes, for example, is determined by the diameter of the template pores and their length is limited by the thickness of the template.694 The principle of the synthetic strategy is straightforward and consists of three main steps: (i) preparing the original template; (ii) depositing the target shell material onto the surface of the template; and (iii) selective removal of the original template to obtain hollow structures. Among the various types of templates, porous anodic aluminum oxide (AAO) is by far the most used and is an effective hard template, due to its narrow size distribution, high pore density, nearly parallel porous structures, and easily controlled pore diameter.695 AAO templates are typically prepared by the anodization of aluminum metal in an acidic solution, and its membranes contain cylindrical pores of uniform diameter arranged in a hexagonal array. Additionally, the templates are thermally and mechanically stable and can be employed under more rigorous reaction conditions. The dimensional flexibility and ordering in the template approach allow for the preparation of high aspect ratio nanowires and nanotubes, as well as ordered arrays, both of which have proven difficult to achieve in Cu chalcogenide NC systems to-date. However, there are some weaknesses to this approach such as the template removal problem, where the hard template needs to be selectively etched with aqueous acids and bases to release the fabricated nanotubes. The inherent difficulty of achieving high product yields and the lack of structural robustness and integrity of the shells upon template removal are also considered to be intrinsic disadvantages to the use of hard templates.696 In turn, the multistep procedures involved in template-based methods and the high production cost of template materials (e.g. AAO) impede the scale-up of such methods for practical applications. A schematic illustrating the formation of CZTS nanowires and nanotubes with AAO templates, in a modified sol-gel solution approach, is shown in Figure 16.697 In a sol-gel approach, a soluble precursor molecule is hydrolyzed to form a dispersion of

Figure 15. (a) Schematic of a Teflon line stainless steel autoclave, typically used for solvothermal/hydrothermal synthesis of NCs. Reprinted with permission from ref 672. Copyright 2002 Royal Society of Chemistry. (b) Photograph of the setup employed for the synthesis. Reprinted with permission from ref 674. Copyright 2016 Elsevier B.V.

When water is used as solvent, the process is more appropriately called hydrothermal synthesis and, in turn, possesses desirable merits of using a clean, green, low cost, and highly abundant solvent. The synthesis under hydrothermal conditions is usually performed above the supercritical temperature and pressure of water, which lie at 374 °C and 218 atm, respectively.673 At this critical point, water is said to be supercritical, in that it can exhibit characteristics of both a liquid and a gas. It boasts exceptionally high viscosities and can easily dissolve reactants, which would otherwise exhibit appreciably low solubilities under ambient conditions.673 Parameters such as water pressure, temperature, and reaction time can be tuned to maintain a high nucleation rate and a good resultant size distribution. Typically, both the solvothermal and hydrothermal processes for preparing Cu chalcogenide NCs consist of precursors, solvents, and organic additives. The precursors are usually organometallic compounds, metal complexes, or inorganic species, which are subjected to high pressure environments for extensive periods, some for up to 40 h, in the autoclave. Many factors, such as the chemical properties of the precursors, solvents, and additives and temperature and time, govern the solvothermal process. The most commonly used solvents are ethylenediamine (EDA), ethylene glycol, ethanol, and water. In particular, EDA is an excellent solvent, due to its strong polarity, relatively low critical pressure, and its strong chelating ability, and many inorganic species have reasonable solubility in EDA. Solvent choice affects the size and morphology control because the physicochemical properties of the solvent regulate the precursors reactivity, solubility, and diffusion behavior.22,675,676 If long alkyl chain fatty acids or amines are chosen as organic surfactants in the system, the resultant NCs will usually possess hydrophobic surface properties.675 This inhibits their application in the fields of bioimaging or biomedicine. Ethylene glycol alleviates this problem by producing NCs with 5897


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Figure 16. Schematic illustration for the formation of nanotubes and nanowires using AAO templates in a modified sol-gel solution approach. Reprinted with permission from ref 697. Copyright 2011 Royal Society of Chemistry.

phases sharing an interface. Moreover, it has the merit of avoiding the removal of templates because the template is depleted spontaneously during the hollowing process.711,712 In the simplest scenario, the synthesis of hollow nanostructures by the Kirkendall effect is a one-pot, two-step process, as illustrated in Figure 17.

colloidal particles (the sol), which acts as the precursor for an infinite network of particles (the gel).698,699 The gel is then typically heated to yield the desired material. The modified sol-gel approach was accomplished by conducting the sol-gel synthesis within the pores, in which the AAO template was immersed in the CZTS precursor solution for a desired amount of time and annealed in a sulfur atmosphere at 550 °C for 1 h, to directly omit the typical pre-drying step in air.697 The template was then etched in a 1 M NaOH aqueous solution to obtain the nanowires and nanotubes. Another approach involved immersing the AAO template in a CZTS precursor solution, followed by heating at 230 °C for 70 h in an autoclave and then etching off the template with NaOH, to obtain highly ordered nanowire arrays.700 Apart from hard templates such as AAO, sacrificial template methods have also been investigated for the fabrication of hollow nanostructures. In this method, a resultant shell forms around the surface of the sacrificial template, taking the shape of the template while being gradually consumed to form a hollow interior. The sacrificial template method (denoted as crystal templating in this review) offers many advantages, in that it avoids the template removal problem, which can prevent possible damage to the produced shell, and it also allows for the formation of a diverse range of shapes, which is simpler and more practical than directly synthesizing the hollow structures. It is also considered to be a more efficient process because it requires no additional surface functionalization and the formation of a shell is guaranteed by the chemical reaction. For example, Cu nanowires have been used as sacrificial templates for the formation of hollow CuS nanotubes,701 while Cu2O NCs (with defined morphologies such as cubic, octahedral, and star-like shapes) were used to form Cu2S mesocages.692 Wu et al. presented a universal sacrificial template method for the preparation of CIS, CIGS, and CZTS nanoplates using presynthesized CuS nanoplates as the starting template.702 A variety of other Cu chalcogenide NC compositions, such as CuSe nanotubes,703 CIS nanotubes,695 ultrathin CISe nanoplates,704 and CZTSe nanowire bundles,705 have also been reported via the use of sacrificial templates.

Figure 17. Schematic illustration of the nanoscale Kirkendall effect for the formation of hollow NCs. Reprinted with permission from ref 712. Copyright 2013 American Chemical Society.

The first step involves the reaction of solid particles of at least one element in the final shell (A). In the next step, a second element (B) in the solution or gas phase is reacted with A to form AB nanostructures.712 The consequent outward diffusion of the core element A through the growing AB composite generates a cavity inside the particle and removes the interior core.706,712−714 The nanoscale Kirkendall effect has also been employed to form hollow nanostructures, where differences in the diffusion rates between two components cause a supersaturation of lattice vacancies, which coalesce and create a void within the nanostructure.706,707,712−714 Functional nanomaterials with controlled hollow interiors and shell thickness have received considerable attention, due to their particular functional and mechanical properties, large surface area, low material density, and the ability to encapsulate multifunctional active materials within their interior cavity.696,715−717 This allows them to serve as ideal building blocks for the fabrication of lightweight structural materials, with potential applications in catalysis, sensing, drug delivery, and optical devices. The nanoscale Kirkendall effect is the most established approach to form hollow nanostructures in Cu chalcogenide compositions to-date. For example, Cu7S4 nanocages,718 Cu2Se nanoboxes,719,720 Cu2Te polyhedrons,142,143 and CIS nanoplates721 with a central hollow structure have been synthesized by employing the Kirkendall effect.

4.4. Kirkendall Effect-Induced Method

Since the initial report in 2004,706 the formation of hollow nanostructures through the nanoscale Kirkendall effect has been applied to the synthesis of a plethora of materials, including metal oxides and chalcogenides.707−710 The Kirkendall effect is a diffusion-related phenomenon occurring between reactive

4.5. Cation Exchange Method

Cation exchange has become a powerful method for the postsynthetic chemical modification of NCs. It gives access 5898


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growth mechanism, which allowed for the attainment of octapodshaped NCs, that had a sphalerite CdSe core and wurtzite CdE (E = S, Se, Te) arms. CIS NCs have also been used as parent NCs to achieve CIZS NCs by partial cation exchange.502,746 Even more complex CTSSe and CZTSSe NCs were obtained by controlling the amount of Zn and Sn precursors to be incorporated into the parent Cu2SeS NCs.729 Furthermore, Cu2Se NCs have been demonstrated to serve as important intermediates in the conversion of CdSe NCs into ZnSe NCs (Figure 18a−c), in that the Cu2Se NCs were able to

to the formation of NCs with compositions, morphologies, and phases that are not readily accessible by conventional synthesis methods.141,501,722−735 Cation exchange takes place instead of anion exchange because the cations have a higher diffusion rate than that of the anions. For example, Cd chalcogenide NCs can be transformed into their Cu/Ag analogues through a singlestep, room temperature process, involving the replacement of cations within the crystal lattice (i.e. Cd2+) with a different metal ion (i.e. Cu+/Ag+).114,722,723,736−739 In bulk materials, chemical transformation reactions are slow, due to the high activation energies required for the insertion and removal of cations.736 However, in NCs, the large surface to volume ratio, coupled with lower activation barriers to diffusing ions, permits rapid cation exchange reactions to occur within seconds at room temperature. This demonstration of dramatically accelerated kinetics in NCs is an attribute which is unique to the nanoscale.723 In recent years, many review articles on cation exchange have been published by various leading groups in the field.723,730,736,740−742 In a seminal paper in 2004, Alivisatos and co-workers reported the complete and reversible cation exchange of Cdbased CdSe NCs into Ag-based Ag2Se NCs under ambient conditions in a short period.736 The cation exchange reaction in the NCs was driven at room temperature by using an excess of the incoming cation (e.g. Ag2+ or Cu+) and/or preferential solvation of the parent cation (e.g. Cd2+).736,743 In the case of CdS to Cu2S NCs, the transformation occured within one second at room temperature and was apparent by a color change from yellow to brown, due to the different band-gap energies in the NCs.114 The driving force behind the cation exchange is solvation, meaning that the parent cation is more soluble in solution rather than in solid form, and this results in the net dissolution of the parent cation. Pearson’s HSAB theory663 plays an important role in cation exchange, in terms of predicting the affinity of metal ions to solvents; that is, Cd2+ is regarded as a hard acid, and so, a hard base such as methanol can preferentially solvate Cd2+ to allow for the formation of Cu2S NCs from the parent CdS NCs. The reverse cation exchange reaction (from Cu2S to CdS) can also be induced through a high excess of Cd2+ ions, coupled with the preferential solvation of Cu+ (a soft Lewis acid) with tributylphosphine, TBP (a soft Lewis base).738 It is the strong affinity of a soft Lewis acid (Cu+) with a soft Lewis base (TBP) that makes alkylphosphines like TBP or trioctylphosphine (TOP) suitable as effective agents to extract Cu+ ions from the NCs and replace them with other cations, without disrupting the anion sublattice. Partial and complete conversion to Cu2S is also possible by adjusting the concentration of substitutional cations in the solution and, thus, can be used to control the relative volume fraction of two crystals within a binary heterostructure.60,114,726,737,744 In particular, Cu chalcogenide NCs are often used as parent NCs or intermediates in sequential cation exchange reactions because Cu+ is easily exchanged by other cations. For example, nonluminescent Cu2−xS NCs were converted into luminescent CIS NCs731 and CIZS NCs501 by partial cation exchange rections. Another report used Cu2−xS NCs to form CIS nanoplates, which had a hollow or intact morphology depending on the indium precursor used and the reaction temperature.721 Cu2−xSe NCs have been used as parent NCs to achieve octapod-shaped CdSe/CdE NCs (E = S, Se, Te), where the Cu2−xSe NCs underwent a cation exchange with Cd2+ ions to form sphalerite CdSe NC.745 This was followed by a seeded

Figure 18. (a−c) Cation exchange of CdSe NCs into ZnSe NCs, where the Cu2Se NCs served as important intermediates and preserved the size, shape, and crystal phase of the starting CdSe NCs. Reprinted with permission from ref 724. Copyright 2011 American Chemical Society. (d) Cation exchange of dot-in-rod nanoheterostructures, starting from CdSe/CdS nanorods, transitioning through Cu2Se/Cu2S nanorods in the intermediate stage, and finally, yielding ZnSe/ZnS nanorods. Reprinted with permission from ref 735. Copyright 2012 American Chemical Society.

transfer all of their morphological and structural information to form ZnSe NCs and preserve the characteristics of the starting particles (i.e. size, shape, and crystal phase).724 Jain et al. investigated dot-in-rod nanoheterostrucutres and demonstrated that a CdSe NC dot embedded in a CdS rod (denoted as CdSe/CdS) could be exchanged to a PbSe/PbS rod via a Cu2Se/Cu2S rod structure.739 Building on this work, Manna and co-workers developed sequential cation exchange procedures to form Cu2Se/Cu2S and ZnSe/ZnS dot-in-rod nanoheterostructures by starting from CdSe/CdS dot-core/rod-shell nanorods (Figure 18d).735 The same group also demonstrated that it was possible to selectively exchange Cu+ cations in the core region of Cu2Se/Cu2S core/shell nanorods, with Ag+ and Hg2+ and form cores comprised of either Ag2Se or HgSe, respectively, without disrupting the Cu2S nanorod shell.727 Luminescent CISe/CIS dot-core/rod-shell heteronanorods were achieved in a separate work by van der Stam et al. by implementing a sequential topotactic cation exchange reaction, using CdSe/CdS dot-in-rod samples.733 In this particular report, Cu2Se/Cu2S nanorods were 5899


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formed by the cation exchange of Cu+ for Cd2+, before being transformed into CISe/CIS dot-in-rod nanoheterostructures by the self-limited partial exchange of In3+ for Cu+.

contain important synthesis parameters (i.e. precursors, ligands, solvents, growth temperature, and time), which allow the size, shape, composition, and crystal phase of the NCs from each literature report to be deduced. The investigated synthetic methods are abbreviated to the following in the “Method” column in the synthesis table: Heating-Up (HU); Hot-Injection (HI); Hydrothermal (HT); Solvothermal (ST); Crystal Templating Method (CTM); and the Kirkendall-Effect (KE). In many review articles, the injection temperature and the exact precursor (which is injected) are not specified for hot-injection procedures within the synthetic tables. Based on their importance in the synthesis, these specific details are denoted in square brackets “[ ]” within each table in the synthesis section. Note: In the “Precursor, Ligands, Solvents” column in each table, the injected precursor is outlined within the “[ ]” brackets, and its corresponding injection temperature is specif ied in the “Temperature/Duration” column. This note applies to all synthetic tables in section 5.

5. SYNTHESIS OF COPPER CHALCOGENIDE NANOCRYSTALS For many applications, it is desirable to fabricate NCs with controlled size, shape, crystal phase, and stoichiometry that are readily dispersible in solvents. Since these characteristics are known to directly influence the functional properties, it is of great interest to explore the synthetic routes to achieve control and tunability of these vital parameters. To provide a complete overview of the synthesis of Cu chalcogenide NCs, we have divided the synthesis section into four main categories: (i) binary semiconductor NCs; (ii) ternary semiconductor NCs; (iii) quaternary semiconductor NCs; and (iv) semiconductor NCs with two chalocogen components (i.e. sulfur and selenium). In the first category, a discussion of the binary Cu2−xS, Cu2−xSe, and Cu2−xTe materials is covered. In the second category, the wide variety of elemental compositions in ternary semiconductor NCs are discussed, starting with the most established ternary materials in the literature, which are the Cu-III-VI2 group of materials: CuInS2 (CIS) and CuInSe2 (CISe). The replacement of In with other elements (such as Ga, Ge, Sn, Sb, Bi, Fe, and Cr) is also explored in this section. The attainment of a diverse range of compositions has not only demonstrated the diversity and versatility of ternary Cu chalcogenide NCs, but it has also allowed for novel optoelectronic and photoluminescence (PL) properties to be achieved. In addition, these properties can be specifically tuned for the desired end application by varying the composition and stoichiometry of the synthesized NCs. Less extensively researched NCs are discussed in the “other ternary compositions” subsection, which covers NC compositions such as CuFeS2, Cu3BiS3, CuInTe2, and CuCr2Se4. The third category covers quaternary semiconductor NCs, which are typically comprised of In and Ga to form the Cu-III-VI2 compounds, CuInGaS2 (CIGS) and CuInGaSe2 (CIGSe), or consist of earthabundant elements such as Zn and Sn to form the Cu2-II-IV-VI4 compounds, Cu2ZnSnS4 (CZTS) and Cu2ZnSnSe4 (CZTSe). The ability to exquisitely tune the band gap in these quaternary materials, by manipulating the composition, has been a key motivator in the development of synthetic routes to these materials, with the view of providing optimum functional material properties for the desired end application. Other quaternary semiconductor NCs can also be formed by the substitution of an element with Ge, Cd, or other transition metal elements such as Co, Mn, Ni, or Fe. The fourth and final part of the synthesis section focuses on the synthesis of NCs with two chalcogens, which provide an additional degree of composition control through variation of the chalcogen (S/Se) ratio. The main objective of this synthesis section is to review the numerous solution-based synthetic protocols, which have been used to synthesize a wide variety of Cu chalcogenide NCs, with special emphasis on how the synthesis has progressed from the pioneer reports onto the many avenues that are now available to instigate control over the NC size, shape, phase, and composition. Furthermore, each Cu chalcogenide NC composition contains a synthesis table, which provides an overview of the experimental procedures to specific NC compositions. These tables can be used to pinpoint crucial work in the area, as well as expedite the discovery of new routes to other NC shapes, compositions, and phases. In particular, the synthetic tables

5.1. Binary Semiconductor Nanocrystals

In the following subsections, the breadth of solution synthesis methods to form Cu2−xS, Cu2−xSe, and Cu2−xTe NCs are covered, starting with the most explored NC composition, Cu2−xS, followed by the replacement of sulfur with a different chalcogen (selenium) to form Cu2−xSe and, finally, the effect of the heaviest chalcogen (tellurium) to form the Cu2−xTe composition. Each material is discussed in a focused subcategory to emphasize the synthetic procedures that are prevalent in the literature. 5.1.1. Copper Sulfide (Cu2−xS). 5.1.1.1. Introduction. Copper sulfide (Cu2−xS) is the most extensively studied binary Cu chalcogenide NC composition and is a p-type semiconductor with a direct band gap, ranging from 1.2 to 2.0 eV. Compared with its Cd- or Pb-based counterparts, it offers a benign environmental profile by using nontoxic and earthabundant materials, which makes it a promising material candidate for future technological applications. Cu2−xS exhibits a rich phase diagram and is known for its diversity of crystal structures and phases. In particular, the stoichiometric factor (2−x) in Cu2−xS varies in the range between 1 and 2, which gives rise to stoichiometry-dependent band gaps of 1.2 eV for chalcocite (Cu2S), 1.5 eV for digenite (Cu1.8S) ,and 2.0 eV for covellite (CuS).6,42 The stoichiometry of Cu2−xS has been observed to have a huge impact on the carrier density and the plasmonic properties, with an increase in the Cu deficiency causing an increase in the carrier concentration. Specifically, stoichiometric Cu2S has no free carriers, Cu1.93−Cu1.97S has a low carrier density, Cu7.2S4 has a moderate carrier density, and CuS has the highest carrier density.565 Thus, control over the stoichiometry provides a useful tool for achieving dynamic, even reversible tunability by varying the free carrier density. This feature is prohibited in noble metal NCs, as their plasmonic properties are permanently locked-in, once the NC size and shape have been engineered.561 Over the past decade, significant progress has been made in identifying routes to fine-tuning the crystal phase and shape of Cu2−xS NCs in solventless synthesis approaches747−750 and hydrothermal,751−757 solvothermal,64,676,758,759 and colloidal hot-injection and heating-up approaches.43,44,61,62,545,548,760−766 Across the solution synthesis methods, DDT is often used in the experimental protocol, acting as the sulfur precursor, reducing agent, and surface capping agent. The success of DDT in the reaction stems from the strong interaction of Cu+ (a soft 5900


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Lewis acid) with the thiolate moiety (a soft Lewis base).8 This leads to a weakened C−S bond in DDT, which lowers the temperature required to cleave it during metal sulfide formation. Cu 2−x S NCs can take on many different morphologies, including spherical nanoparticles,44,474,545,548,767,768 nanoplates,61,64,747,748,750,762,763,765,766,769,770 nanorods,771 nanowires, 749,772 nanosheets, 70 nanowhiskers, 773 nanoflowers,676,758,774−776 nanoflakes,753,777 and hollow nanospheres.778,779 In this regard, zero-dimensional (0D) NCs such as QDs or spherical particles have high surface areas and an active surface plasmon resonance, which are suitable characteristics for optoelectronic and biological applications. On the other hand, one-dimensional (1D) anisotropic structures, such as nanorods and nanowires, are more suitable in electronic devices and photovoltaic cells because of their directional, electronic transport and structural integrity. Moreover, the interest in studying the formation of Cu2−xS NCs is increasing, as Cu2−xS was revealed to be the parent compound to the much larger family of ternary and quaternary Cu chalcogenide semiconductor NCs. Of all the morphologies, Cu2S nanoplates are the most common shape reported across the solution routes and are inclined to self-assemble into face-to-face stacks, due to their van der Waals attraction.474,747,748,750,761−763,765,769,780 In particular, the attainment of nonspherical shapes such as nanoplates reflects the internal hexagonal crystal structure. As both end members, Cu2S and CuS, possess hexagonal structured crystal systems; these compositions are particularly liable to anisotropic growth. An important point to note in the characterization of nanoplates is that the acquired TEM image is a two-dimensional (2D) projection of a three-dimensional (3D) object, which can often lead to misinterpretation of nanoplates as nanorods, especially when the nanoplates are stacked side-by-side.747,748 Thus, tilting the stacked nanoplates should be implemented during the TEM characterization stage to give conclusive proof of nanoplate formation. An overview of the experimental procedures to Cu2−xS is outlined in Table 3. According to the literature, the most common strategy to tune the composition of Cu2−xS is by adjusting the molar ratio of the Cu and S precursors during the synthesis. Control over the size and shape can be achieved by changing reaction parameters, such as growth temperature and time, and through careful selection of the Cu and S precursors. In the case where two separate phases of Cu2−xS are obtained within a particular report, both phases are specified in the “Material” section of the synthesis table. 5.1.1.2. Initial Work on Cu2−xS Using Single-Source Precursors. Early studies on the synthesis of Cu2S NCs focused on the thermal decomposition of Cu thiolates, which behaved as single-source molecular precursors because they contained both the Cu and S elements.61,747−750,768−770,772,780,785,788−791 In general, single-source molecular precursors are of interest because the potential of any side reactions, occurring with separate metal and S sources, is eliminated. The Korgel group reported the first solventless synthesis method in 2003, in which Cu 2 S nanoplates were formed by the thermal decomposition of the Cu-dodecylthiolate (CuSC12H25) precursor.747,748,770 A series of steps were required prior to the synthesis to form this molecular precursor. Briefly, an aqueous copper nitrate solution was combined with chloroform and sodium octanoate to solubilize the Cu2+ ions in the organic phase. The resulting copper octanoate complex was added to

DDT, after which the solvent was evaporated to obtain a waxy Cu-thiolate precursor, that was dried at 148 °C for 140 min. The synthesis formed nanoplates, which self-assembled into ribbons of stacked platelets, exhibiting parallel alignment (Figure 19a).748 A typical preparation yielded 10−20 mg of purified nanoplates (10−20% yield). By increasing the reaction temperature to 200 °C, nanoplates were formed in the shorter growth time of 15 min, which was attributed to the enhanced rate of thermolysis of the C−S bond in DDT and the consequent increased availability of the sulfur monomer.748 Several groups adapted Korgel’s original procedure and investigated the effect of different synthetic modifications on the resultant shape of the NCs.749,790,792 For example, control over the stirring speed and time during the precursor preparation led to the formation of Cu2S nanowires, up to several micrometers in length, after the precursor was heated to 155 °C for 120 min.749 In this case, nanowire formation was attributed to the attainment of a highly polymerized thiolate precursor, which was generated at gentle stirring speeds. On the contrary, a lower polymerized thiolate precursor was formed by using high stirring speeds, which led to the formation of Cu2S nanorods under identical reaction conditions. Varying the alkyl chain length of CuS(C12H25) has also been investigated, with short chain precursors producing 2D nanosheets and longer chained variants producing Cu2S nanoplates.790 A recent in situ synchrotron XRD study proved that Cu2−xS NCs were formed from the thermolysis of CuS(C12H25) building blocks, where the C−S bond cleavage in DDT is the limiting step in their formation and growth.792 Other thiolate-derived, single-source precursors have also been investigated, particularly Cu-dithiocarbamate, which has been used to generate Cu2S nanobarrels,789 Cu1.8S quantum dots (QDs),785 and Cu2S nanowires.772 Quantum confinement effects were observed in the Cu1.8S QDs, and while their exact size was not quoted, the QDs had a band-gap energy of 2.35 eV and were considerably blue-shifted, compared to the bulk band gap (1.21 eV) of Cu2S.785 Even stronger quantum size effects were observed in Cu2S nanowires (Figure 19b), where the nanowires had diameters of 2.5 and 1.7 nm, which corresponded to band-gap energies of 3.47 and 3.69 eV, respectively.772 While single-source precursors are generally prepared in an organic− aqueous phase transfer reaction in which the aqueous phase is discarded, Zhuang et al. demonstrated that it is possible to directly heat the prepared organic−aqueous solution in an autoclave at 200 °C.474 In this two-phase reaction system, the upper oil/organic phase was composed of DDT or its toluene solution, and the lower aqueous phase contained the Cu ions and added anions (Ac− or Cl−) for controlling the growth. Spherical Cu2S NCs or Cu2S nanoplates were formed at the interface when Ac− or Cl− ions were used, respectively. 5.1.1.3. Heat-Up Approaches. In more recent years, a plethora of synthetic approaches have moved toward a simpler strategy, that of dissolving and heating metal salt precursors in an organic solvent, such as oleylamine (OLA) or DDT under high temperatures. This precluded the requirement to prepare precursors prior to the synthesis and alleviated the necessary evaporation/drying steps required in their preparation. The most adapted Cu precursors are metal salts such as copper(II) acetylacetonate (Cu(acac)2) or copper chloride (CuCl), while sulfur powder or DDT are the sulfur precursors of choice. For example, a heat-up approach using Cu(acac)2, sulfur powder, and OLA produced Cu2S nanoplates, when the reaction was kept at 230 °C for 2 h.762 Prolonging the reaction time to 4 h 5901


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Table 3. Overview of the Experimental Procedures for the Synthesis of Cu2−xS NCs Size (D × L) / Morphology

Ref

4 × 12 nm/plates 1.7−2.5 nm × 50 μm/wires 7−20 nm/spherical, plates 7−12 nm/spherical, 11 × 47 nm/plates 6−18 nm/spherical 6 nm/spherical, 4 × 9 nm/plates 3−10 nm/spherical 5 × 13 nm/disks 5 ×x 13 nm/plates 3 × 21 nm/plates 4−7 nm/spherical 3−13 nm/spherical

747,748,770 772 764 765 781 762 548 766 75 563 782 545

2−20 nm/spherical 9 × 17 nm/plates

767 763

9−23 nm × 10−100 nm/rods

771

HI HI

140−200 °C/15−140 min 160 °C/12 h 150−230 °C/15−140 min 200−250 °C/10−120 min 200 °C/8−210 min 200−230 °C/1−2 h 220 °C/30 min 182 °C/1 h 200 °C/30 min 85 °C/1 h 120−180 °C/40 min [115−140 °C], 115−140 °C/ 1.5−3 min [200 °C], 190 °C/15 min [160 °C], 160−220 °C/ 5−210 min [120−180 °C], 180 °C/ 5−10 min [180 °C], 180°C/1 h [25 °C], 25 °C/8 min

72−155 nm/dodecahedrons 4−34 nm/spherical, 4 × 7.5 nm/plates

783 559

HI HI

[180 °C ], 180 °C/10 min (180 °C), 150 °C/30 min

5 × 28 nm/disks 6 × 17 nm/disks

561 62

HI

[95 °C], 95 °C/18 h

3 × 453 nm/sheets

70

HI

[110 °C], 180 °C/3−20 min

2.5−6 nm/quantum dots

37,44

HI HI HI

13 × 70 nm/plates 6−20 nm/spherical 11−19 nm × 23−85 nm/plates, 5−9 nm/spherical

784 564 61

HI

[300 °C], 300 °C/10 min [180 °C], 180 °C/5−15 min [135−210 °C], 135−210 °C/ 15 min [250 °C], 250 °C/10−40 min

Unspecified/quantum dots

785

HT

90−110 °C/40 h

752

HT

130 °C/15 h

50−200 nm × 2−4 μm/wire-, tube-, vesicle-like aggregates of plates 20 × 400 nm/flakes, 6 × 20 nm/disks

HT ST ST ST HU

180 °C/12−48 h 120 °C/24 h 60 °C/3 h 140 °C/24 h 50 °C/2 h

50 × 500−800 nm/flower-like structures 8 × 26 nm/plates 27 nm/sheets 200 nm × 1−1.5 μm/cuboctahedrons 3 nm/spherical

677 64 391 786 787

HU ST

90 °C/15 min 60 °C/4 h

11 nm/spherical 200 nm/hollow spheres

556 686

CTM CTM

80 °C/12 h 0 °C/10−30 min

120−150 nm × 40 μm/hollow tubes 0.3−1 μm/hollow cages

701 692

KE

90 °C/6 h

320 nm/hollow cages

718

Material

Precursors, Ligands, Solvents

Method Temperature/Duration

Cu2S Cu2S Cu2S Cu2S Cu2S Cu2S Cu2S Cu1.8S, CuS CuS CuS Cu2S Cu1.94S, CuS Cu2S Cu2S

Cu-thiolate Cu(DDTC)2, DDT, OA Cu-oleate, DDT, OLA Cu(acac)2, DDT, OLA Cu(acac)2, DDT Cu(acac)2, S, OLA Cu(acac)2, S, DDT, OA Cu(acac)2, S, DCB, OLA, OA CuCl, S, OLA, ODE CuCl, S, OLA, TOPO CuCl, (NH4)2S, OLA CuCl, OLA, [S-OLA]

HU HU HU HU HU HU HU HU HU HU HU HI

Cu(St)2, ODE, [DDT] CuAc, ODE, TOPO [DDT]

HI HI

Cu1.94S

CuAc, ODE, TOPO, [t-DDT]

HI

Cu1.96S CuS

CuCl2, OLA, [TBDS] CuCl2·2H2O, OLA, toluene, [(NH4)2S] CuCl, ODE, OLA, OA, [S/OLA] CuCl2, OLA, [S/OLA] CuCl, OLA, OTA, [S/OLA/ OTA] [Cu(acac)2/OA], NH3-DDTC, DDT, OA [Cu(DDTC)2/OLA], OLA [CuCl/OLA/OA], S/ODE [CuTB/TOP, TBPT], DDT

CuS Cu1.75S, CuS CuS Cu2S, Cu1.93S Cu9S5 Cu1.8S Cu1.75S, Cu2S Cu1.8S Cu1.8S, Cu1.97S CuS, Cu2S CuS CuS Cu9S8 CuS CuS CuS CuS CuS Cu7S4 Cu7S4

[Cu(DDTC)2/TOP], [TOP/S], TOPO CuCl, thiourea, TEDA, TMEA, DBA Cu(NO3)2, CTAB, pentanol, hexane, DDT CuCl2, Na2S, PVP CuAc2·H2O, CS2, HDA, toluene CuCl2, thiourea, ammonia Cu(NO3)2·5H2O, S, EG CuCl2·2H2O, TAA, TGA, H2O, NaOH CuCl2, Na2S, sodium citrate CuSO4, CS2, Triton X-100, cyclohexane Cu NWs, thiourea, EG Cu2O NCs, Na2S, NaOH, ammonia CuSO4, thiourea, PVP, NaOH, ascorbic acid

753

were tuned in size by varying the reaction time or the amount of DDT.765 Specifically, the diameter increased from 7.3 nm (10 min) to 11.8 nm (60 min) with prolonged reaction times, while diameter increases of 6.1 nm (5 mL DDT) to 16 nm (15 mL DDT) were observed with increased volumes of DDT, for reactions kept at 200 °C for 60 min. Another heat-up approach demonstrated that Cu1.94S spherical particles (Figure 19c) could be formed, without the use of OLA, by solely heating Cu(acac)2 and DDT (30 mL) at 200 °C, in which the particle diameter was tuned by variation of the reaction time.781 These particles were further employed as

led to wider size and shape distributions, indicative of an Ostwald ripening process. By lowering the temperature to 200 °C, monodisperse spherical particles (5.9 ± 0.7 nm in diameter) were obtained, which demonstrated that the NC shape could be controlled by modulation of the reaction temperature. The method was further adapted to grow Cu2S nanoparticles on multiwalled carbon nanotubes and form a hybrid nanostructure, with potential application as an amperometric sensor for glucose detection.794 A similar approach using Cu(acac)2, OLA, and DDT as the sulfur precursor (instead of sulfur powder) also produced spherical particles at 200 °C, that 5902


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Figure 19. (a) SEM image of Cu2S nanoplates deposited onto a substrate, forming 3D colloidal crystals. Adapted with permission from ref 748. Copyright 2003 American Chemical Society. (b) TEM image of ultrathin Cu2S nanowires. Adapted with permission from ref 772. Copyright 2005 American Chemical Society. (c) TEM image of Cu2S spherical particles. Adapted with permission from ref 781. Copyright 2008 American Chemical Society. (d) TEM image of Cu2S nanoplates with face-to-face stacks of nanoplates evident, which could be misinterpreted as nanorods. Adapted with permission from ref 763. Copyright 2010 American Chemical Society. (e) SEM image of Cu1.94S nanorods. Adapted with permission from ref 771. Copyright 2012 American Chemical Society. (f) TEM image of Cu1.96S dodecahedrons formed at high precursor concentrations. Adapted with permission from ref 783. Copyright 2011 Royal Society of Chemistry. (g) SEM image of ultrathin CuS nanosheets (scale bar = 200 nm) forming domino-like superstructures. (h) SEM image of a monolayer of CuS nanosheets on a Si substrate (scale bar = 200 nm), with the inset showing a photograph of the CuS nanosheets in solution. Panels g−h were adapted with permission from ref 70. Copyright 2012 Macmillan Publishers Ltd: Nature Communications. (i) SEM of CuS cuboctahedrons. Adapted with permission from ref 786. Copyright 2006 American Chemical Society. (j) SEM image of hierarchical flower-like CuS nanospheres. Adapted with permission from ref 793. Copyright 2013 Royal Society of Chemistry. (k) SEM image of CuxS nanocubes prepared from Cu2O crystal templates. Adapted with permission from ref 692. Copyright 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim. (l) SEM image of hollow CuS spheres with a thin-shell. Adapted with permission from ref 711. Copyright 2008 American Chemical Society.

methods use temperatures >200 °C to produce uniform Cu2−xS NCs, but low temperatures ( 104 cm−1) and a band gap in the range 1.4−1.6 eV, which makes it an interesting candidate as a p-type absorber for thin film PV. However, there are only a handful of papers reporting the synthesis of this material in the NC form, to-date. For example, Yan et al. developed a hot-injection approach to form Cu3BiS3 NCs, by using Cu(acac)2 and Bi(NO3)3·5H2O as the cation precursors and OLA as the solvent.577 In this particular reaction, a solution of sulfur in oleylamine (S/OLA) was injected into the reaction flask at 220 °C and growth was allowed to proceed at this temperature for 30 min. The authors noted that there was a high potential for the formation of binary CuS and Bi2S3 phases in this system, with single-phase Cu3BiS3 NCs (10 nm in diameter) only occurring within a narrow synthetic window, relating to the choice of precursors and the growth temperature. A simpler but effective thermal decomposition approach was developed by Deng et al., in which a range of Cu3BiS3 NC morphologies were obtained, including spheres, nanosheets, nanowires, and nanoribbons using metal-diethyl dithiocarbamate precursors, with the NC shapes being largely affected by the ratio of precursors and the growth temperature.967 The most interesting application of this material set is their potential for use in photothermal (PT) theragnosis synergistic therapy, where a combined NIR and MRI response allows for detailed and exact information on tumors to be extracted at a very early stage of development.576 In this particular work, large hexagonal nanoplates (150 nm) were produced in a hydrothermal synthesis approach, by using poly(vinylpyrrolidone) (PVP) and poly(ethylene glycol) (PEG) as ligands to render

5.3. Quaternary Semiconductor Nanocrystals

Quaternary Cu chalcogenide compounds are formed by the additional occupation of a cation into the ternary Cu-III-VI2 5933


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measurements (Figure 28b) quantitatively defining red-shifts in the optical band gap, from 2.4 eV for the Cu1.0Ga2.0S3.5 NCs to 1.43 eV for the Cu1.0In2.0S3.5 NCs. A year later, Sun et al. formed chalcopyrite Cu1.0InxGa1−xS2 (0 ≤ x ≤ 1) NCs by the thermal decomposition of a mixture of two single-source precursors, (Ph3P)2Cu(μ-Set)2In(SEt)2 and (Ph3P)2Cu(μ-Set)2Ga(SEt)2, in the presence of benzyl acetate and 1,2-ethanedithiol (EDT) via microwave irradiation at 160 °C for 1 h.971 A progression to darker-colored solutions was observed, as a function of the increasing indium content, and this represented changes in both the band gap (2.3−1.59 eV) and NC size (2.7−3.3 nm) with the increasing indium content. In addition, the authors investigated the effect of using a higher reaction temperature of 240 °C and revealed that the NC size increased from 3.7 to 5.9 nm, and that the band gap could be tuned within a narrower range from 1.9 to 1.4 eV. A heat-up approach for the formation of color tunable, chalcopyrite CIGS QDs was reported by Song et al., in which the NCs had an intentional Cu-deficiency and varied In/Ga ratios.972 This approach was based on the reaction of CuI, InAc3, Ga(acac)3, and DDT at 230 °C for 5 min, and formed CIGS NCs with diameters in the range 1.7−2.2 nm that emitted red to deep-red colors (633−670 nm). Low PL quantum yields (14−16%) were observed for the CIS and In-rich CIGS samples, and negligible PL quantum yields ( Sn > Zn during their flow reactor synthesis of CZTS NCs, and they proposed a similar formation pathway for tetragonal CZTS NCs, in that Sn and Zn atoms subsequently get incorporated into the nucleated Cu2S NCs.989 While this pathway was proposed for tetragonal CZTS NCs, the same formation pathway still holds for wurtzite CZTS NCs formed in hot-injection approaches, as confirmed by Coughlan et al. by performing TEM and EDX line scan analysis of aliquots taken at different stages as the reaction proceeded.998 A separate report by Tan et al. demonstrated that surfaceenhanced Raman scattering (SERS) spectroscopy is a powerful characterization tool to fully identify and differentiate the mixed compositional phases present in CZTS, which also allowed for the growth mechanism to be elucidated.999 It was found that the formation of CZTS proceeded via the nucleation of Cu2−xS particles, followed by diffusion of Sn4+ into Cu2−xS to form kinetically driven Cu−Sn−S (CTS) particles and, lastly, the diffusion of Zn2+ into CTS to form CZTS NCs. Thompson et al. investigated the effects of the initial precursor concentration (loading) on the aspect ratio and composition of CZTS nanorods.1008 By decreasing the loading of the Cu precursor, the nanorods’ aspect ratio was increased because more precursors were left behind in solution for growth, as the rate of Cu sulfide nucleation was lowered. In contrast, increasing the loading of the Cu precursor or changing the loading of Zn or Sn precursor resulted in the formation of spherical NCs and/or phase segregated NCs. This demonstrated that the cationic precursor loading had dramatic effects on the morphology and composition of CZTS NCs. An overview of the experimental procedures to form CZTS NCs with various crystal phases and morphologies is outlined in Table 13. 5.3.3.5. Phase-Selective Synthesis of CZTS NCs. In the interest of controlling the phase in CZTS NCs, some reports have provided details on how to phase selectively synthesize CZTS NCs in colloidal approaches. Most routes employ 5941


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Table 13. Overview of the Experimental Procedures for the Synthesis of CZTS NCs Phase


Method

Temp/Duration

Size (D × L) / Morphology

Ref

KS KS KS KS

Cu(acac)2, Zn(acac)2, Sn(acac)2Br2, OLA, [S/OLA] [Cu(acac)2, ZnAc2, SnAc4, OLA], [S/OLA], TOPO Cu(acac)2, ZnAc2, SnCl2, S, OLA [Cu(acac)2, Zn(acac)2, Sn(acac)2Br2, OLA], [S/OLA], OLA CuCl2, ZnSO4, SnCl2, Na2S, PVP, EG CuCl2, ZnCl2, SnCl4, thiourea, PVP, EG CuCl2, ZnCl2, SnCl4, thiourea, PVP, EG Cu(NO3)2, Zn(NO3)2, SnCl4, TEG, S CuCl, ZnCl2, SnCl2, S, EDA, AAO template CuAc2, ZnAc2, SnCl2, ME, thiourea, AAO template

HI HI HU HI

[225 °C], 225 °C/30 min [300 °C], 300 °C/45−75 min 280 °C/60 min [250 °C], 250 °C/60 min

15−25 nm/irregular 12.8 nm/irregular 11 nm/irregular 12.5 nm/irregular

337,984 985 339,986 983

ST ST ST ST ST, CTM CTM

180 230 230 250 230 550

1002 684 690 691 700 697

HI

992

7−9 nm/irregular 15 nm/irregular 7.3 nm/pyramidal 5−8 nm/spherical 14.5 nm/irregular 20−50 nm/platelike, aggregated 3.3 nm/quasi-spherical 20 × 28 nm/nanoprisms 14 nm/nanoplates 11 × 35 nm/rods

1009 989 1010 1011 1012 1000 1001 994 994 246,995,998

HI

[150−175 °C], 150−175 °C/ 4 min [25 °C], 180 °C/1 h 300−320 °C/continuous flow 250 °C/30 min 250 °C/1 h [300 °C], 300 °C/30 min 200 °C/24 h 200 °C/10 min [240 °C], 240 °C/1 h [240 °C], 240 °C/1 h [150 °C], 240−260 °C/ 15−30 min [100 °C], 250 °C/1 h

10 nm/irregular 500 nm/flower-like 100−150 nm/spherelike 2 μm/microspheres 200 nm × several μm/nanowires 200 nm × 60 μm/nanowires, nanotubes 2−7 nm/quasi-spherical

12 nm/quasi-spherical

340,341,1006

HU

250 °C/30 min

8 ×15 nm/rods

993

HI

[120 °C], 210 °C/30 min

7.3 × 23.7 nm/rods

1008

KS KS KS KS KS KS KS KS KS KS KS KS OTR ZB WZ WZ WZ WZ WZ WZ

Cu(DDTC)2, Zn(DDTC)2, Sn(DDTC)4, ODE, OA, [OLA] CuCl2, ZnAc2, S, OLA, [SnCl4] CuCl2, ZnO, SnCl4, THF, S, OLA, ODE CuI, ZnCl2, SnCl4, CS2, DDT, OLA CuI, Zn(EtXn)2, SnCl4, CS2, DDT, OLA [Cu(acac)2, ZnAc2, SnAc4, S, OLA], TPPA CuCl2, ZnCl2, SnCl2, thiocarbamide, EDA, H2O CuCl2, ZnCl2, SnCl2, thiocarbamide, ethanol, OLA [CuCl2, ZnCl2, SnCl4, DDT], DDT, OLA [CuCl2, ZnCl2, SnCl4, DDT], DDT, OA Cu(acac)2, ZnAc, SnAc4, TOPO, ODE, [t-DDT/ DDT] CuCl2, ZnO, SnCl4, THF, OLA, ODE, [t-DDT/ DDT] Cu(DDTC)2, Zn(DDTC)2, Sn(DDTC)4, HDT, TOA Cu(acac)2, ZnAc2, SnAc4, TOPO, ODE, [t-DDT/ DDT]

HI HU HU HU HI HT ST HI HI HI

°C/12 h °C/24 h °C/24 h °C/48 h °C/70 h °C/1 h

which allowed for gram-scale quantities of kesterite CZTS NCs to be achieved at 300−330 °C.989 In this approach, a mixed solution of precursors was pumped though a 1-meter long bronze tube at a flow rate of 1−5 mL/min. This method formed particles, which were 15 nm in diameter and had a narrow size distribution (Figure 30h). Temperature was identified as an important parameter in defining the composition distribution, where higher reaction temperatures (i.e. 315 °C) led to narrower particle-to-particle compositional differences. It was also noted that compositions far from the stoichiometric composition proved to be unstable after annealing, and this led to phase segregation to form secondary phases such as SnS. Another up-scaled route to CZTS NCs was published around the same time by heating the precursors up to 280 °C for 1 h, and formed kesterite phase CZTS NCs, but the exact yield in grams was not stated.339 Chesman et al. developed a multigram heating-up method using two sulfide precursors of disparate reactivity, carbon disulfide and DDT, and formed 2.8 g (70% yield) of pyramidal kesterite CZTS NCs by heating the precursors up to 250 °C for 30 min.1010 By using a similar route but replacing the carbon disulfide with zinc ethyl xanthate, the authors obtained 3.3 g (80% yield) of kesterite CZTS NCs.1011 However, despite the size distribution being narrow, full shape control was not achieved in any of the up-scaled routes. A highly concentrated synthesis of CZTS NCs was demonstrated by injecting the metal precursor solution into a hot solution of triphenylphosphate (TPPA) at 300 °C.1012 The synthesis yielded 4.5 g of kesterite phased NCs with slightly irregular shapes, which was an impressive yield, but the hot-injection requirement was not a feasible approach for scale-up. A scalable heating-up

contained primarily Cu and Zn, and small particles contained Cu and Sn.1018 They suggested that longer synthesis times produced more consistent CZTS NC inks with higher yields, and with compositions closer to that of the starting precursors. Despite the presence of intraparticle compositional inhomogeneities, the selenization process led to an effective redistribution of the species and an overall enhanced homogeneity in the final selenized layer, thus facilitating the use of heterogeneous particles in optoelectronic devices. This observation suggested that the compositional nonuniformities inside particles were the result of kinetic effects associated with monomer incorporation during growth, and that the composition could be driven to a thermodynamic equilibrium with a sufficient thermal driving force. 5.3.3.7. Upscaled Synthesis Approaches of CZTS NCs. The majority of CZTS solution-based methods to-date have employed the hot-injection technique, using typically elemental sulfur, thiols, or dithiocarbamates as the sulfur source. However, the uncontrolled nucleation rates associated with the injection prohibit the scale-up of hot-injection-based techniques to commercial levels. Moreover, solvothermal methods are not very favorable for safe, large-scale production, as the process is conducted in a stainless steel autoclave with an extremely high vapor pressure. The heating-up method provides an excellent means to generate multigram quantities of NCs and avoids the limitation of precursor injection. That said, specific control over both the precursor and ligand chemistries is crucial, to obtain the necessary balance between the nucleation and growth rates, which are needed to afford high-quality NCs in large quantities. In 2012, Shavel et al. presented a continuous-flow process 5942


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Figure 31. (a) TEM image of CZTSe NCs with a polyhedral geometry obtained in a hot-injection approach. Reproduced with permission from ref 1027. Copyright 2010 American Chemical Society. (b) TEM image of CZTSe NCs with a mixture of triangular, hexagonal, and plate-like NCs obtained in a hot-injection approach. Reproduced with permission from ref 348. Copyright 2010 Elsevier B.V. (c) TEM image of CZTSe NCs forming flower-like nanostructures through aggregation. Reproduced with permission from ref 1028. Copyright 2015 Elsevier B.V. (d) TEM image of quantum confined CZTSe NCs with diameters of 3.4 nm. Reproduced with permission from ref 650. Copyright 2012 American Chemical Society. (e) TEM image of CZTSe nanosheets synthesized in a solvothermal approach. Reproduced with permission from ref 685. Copyright 2011 Royal Society of Chemistry. (f) TEM image of wurtzite-phased CZTSe NCs. Reproduced with permission from ref 349. Copyright 2012 NPG Asia Materials. (g) XRD pattern of CZTSe NCs synthesized with different solvents, namely OLA (pattern 1, red line), increased OLA and low volume of hexadecane (pattern 2, green line), low volume of OLA and increased volume of hexadecane (pattern 3, blue line), OA and hexadecane (pattern 4, pink line). Reproduced with permission from ref 349. Copyright 2012 NPG Asia Materials. (h) TEM image of wurtzite CZTSe NCs with a disc-like morphology. Reproduced with permission from ref 492. Copyright 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim. (i) TEM image of CZTSe NCs formed in a large-scale approach. Reprinted with permission from ref 346. Copyright 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

synthesis approach to wurtzite phase CZTS NCs was recently published by Shavel et al., which allowed for the production of several grams of highly monodisperse, quasi-spherical CZTS NCs per batch.1019 5.3.4. Copper Zinc Tin Selenide (CZTSe). 5.3.4.1. Introduction. CZTSe is an alternative quaternary chalcogenide with high potential for PV applications. It combines many advantages, such as a suitable band gap (1.0−1.5 eV),241,517,1020 a high optical absorption coefficient of up to 105 cm−1,1021 low toxicity, and the relative abundance of its elements. CZTSe is considered to be one of the most promising substitutes for CIGSe in PV, in which the latter material has higher efficiencies but contains the expensive and rare metals of indium and gallium. CZTSe is also promising for thermoelectric applications, both in the bulk1022 and in NC form.346,1023,1024 There are some discrepancies between reports for the band gap of CZTSe, in that the existing experimental data for the bandgap values are not consistent. The majority of the experimentally determined band gaps are between 1.4 and 1.5 eV for films prepared by top-down approaches,433,1020,1021,1025,1026

whereas a report on a monograin powder of CZTSe was shown to have a band gap of 1.0 eV.517 In general, selenides normally have a larger lattice constant and smaller band gaps than sulfides. Theoretical calculations of the band gap have shown that the band gap of CZTSe should be around 1.0 eV,241 which is consistent for the experimental value for the CZTSe monograin powder.517 5.3.4.2. CZTSe NCs in the Tetragonal Stannite Phase. Shavel et al. developed the first solution synthesis of CZTSe NCs in 2010, which relied on the hot-injection of trioctylphosphineselenide (TOP-Se) into a flask containing metal chloride precursors in hexadecylamine (HDA) and ODE solvents at 295 °C.1027 The synthesis formed 20 nm particles in the stannite (tetragonal) phase after a 5 min growth time (Figure 31a). Quantitative analysis of the chemical composition showed the NCs to be consistently Zn- and Sn-poor. Soon after this report, a phosphine-free, hot-injection method was also developed, but the authors injected the metal precursor solution into a flask containing selenium powder in OLA at 150 °C.348 The CZTSe NCs prepared by this method lacked a tight degree of shape 5943


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phase can be obtained through exquisite control of the reaction chemistry. For example, Wang et al. synthesized the wurtzite phase of CZTSe NCs for the first time in 2012 using a hotinjection approach.349 The approach involved the use of Cu-oleate, Zn-oleate, Sn(II) 2-ethylhexanoate (Sn(2-Eth)2), or OLA as solvent and diphenyl diselenide (DPDSe) as selenium source and formed quasi-spherical NCs (Figure 31f), after an injection of the DPDSe precursor at 190 °C followed by heating at 255 °C for 40 min. The choice of solvent was shown to directly affect the crystal phase (Figure 31g), where OLA formed wurtzite phased CZTSe NCs but its replacement with a mixture of oleic acid (OA) and hexadecane formed the stannite phase. In addition, it was noted that the morphology of the NCs changed from spherical particles to nanowires, when the OLA solvent was substituted by OA. The selenium source was also shown to affect the crystal phase, where substitutions of the DPDSe precursor with Se powder or selenourea resulted in NCs crystallizing in the stannite phase, as opposed to the wurtzite phase. A subsequent report by Ennaoui and co-workers used the same selenium precursor, DPDSe, to synthesize CZTSe NCs, but they showed that the obtained NCs could crystallize in three structurally different phases: tetragonal kesterite, hexagonal wurtzite, and orthorhombic wurtz-stannite by controlling the reaction conditions.1030 In addition, the phase stability of wurtzite CZTSe NCs was investigated in this report by annealing at a high temperature, where a complete transformation of the wurtzite type CZTSe NCs into kesterite type CZTSe grains was observed after annealing at 540 °C for 30 min under an Se-vapor atmosphere. Singh et al. used a modified hot-injection approach to form CZTSe NCs with a disc morphology (Figure 31h) and found that OA was critical in the synthesis to retaining the wurtzite phase for this composition, with kesterite CZTSe NCs occurring in its absence.492 Li et al. investigated the use of thiourea and selenourea precursors in a hot-injection approach and found that wurtzite CZTS NCs were formed with thiourea, but zinc-blende CZTSe NCs were formed with the selenium analog, selenourea.1032 In an attempt to understand the structural selection, the authors conducted an investigation on the Cu binary chalcogenide NCs, which formed in the initial stages of the synthesis. They found that deliberate control over the composition in the binary systems enhanced the corresponding structure of CZTSe (or CZTS) NCs, in which Cu2Se (or Cu1.96S, Cu9S5) generally preferred to form the wurtzite phase, whereas stoichiometric CuSe (or CuS) resulted in the zinc-blende phase. An overview of the experimental approaches to form CZTSe NCs is outlined in Table 14. 5.3.4.4. Upscaled Approaches to CZTSe NCs. Very few reports have investigated the scale-up potential of CZTSe NCs, where the ability to scale-up is pivotal for their use in thermoelectric and PV applications. Fan et al. developed a large-scale colloidal approach and achieved more than 10 g of cubelike CZTSe NCs (Figure 31i) per reaction.346 However, the only downside to this approach was that it involved the injection of a selenium precursor into the reaction flask, which is not an ideal requirement for up-scaled reactions. That said, this was the first report to generate gram-scale quantities of CZTSe NCs, as well as the first to investigate the thermoelectric properties of the material. The authors stated that the selenium precursor, selenium dioxide (SeO2) in ODE, proved to be easier to dissolve than Se powder and could be reduced to elemental Se by OLA or ODE at elevated temperatures. The average chemical composition of the NCs was determined to be

control, forming a mixture of triangular, hexagonal, and plate-like NCs (Figure 31b) after growth to 240 °C for 2 h. The chemical composition of the NCs was found to be Cu2Zn0.84Sn1.24Se4.08, showing slightly Zn-poor and Sn-rich content and deviating from the stoichiometric CZTSe NC composition. Moreover, a band gap of 1.52 eV was reported, which is higher than expected for CZTSe NCs. A heat-up approach to CZTSe NCs using metal chlorides, OLA, TOPO, and selenium powder also led to the formation of NCs with a stoichiometric ratio of Cu2Zn1Sn1.3Se3.75, but the particles suffered from high aggregation and joined together to form anisotropic 3D nanostructures (Figure 31c).1028 Contrary to the previous report, the CZTSe NCs were determined to have a band gap of 1.13 eV and displayed promising optoelectronic and photocatalytic properties, thus making the CZTSe NCs a potential candidate for PV, as well as for photocatalytic applications. Haas et al. reported an interesting investigation on the composition of CZTSe NCs, which were prepared by the injection of an Se/OLA solution into a mixture of the metal precursors in OLA at room temperature.1029 After heating the reaction mixture up to 230 °C, the synthesis formed 18 nmsized CZTSe NCs after a 90 min growth time. This report addressed a key issue often left unanswered in previous studies, that of determining the chemical homogeneity of the NCs. By using the excellent spatial resolution of EDX and EELS, the authors found that individual NCs showed a large variation in the content of Cu, Zn, and Sn, and they concluded that a broad range of chemical composition heterogeneity existed in the CZTSe NCs, due to differences in the reactivity of the precursors. However, by changing the types of precursors used, a stoichiometry close to that of theoretical CZTSe was achieved.988,1030 Liu et al. took advantage of the strongly coordinating and reducing character of DDT, by dissolving selenium powder in a mixture of DDT and OLA, which bypassed the conventional requirement of using extended heating to dissolve the selenium powder.650 This generated a soluble alkylammonium selenide precursor solution, which when injected into a solution of metal chlorides in OLA/DDT, produced quantum-confined CZTSe NCs (Figure 31d) for the first time, having an average diameter of 3.4 nm. EDX measurements indicted that the NCs had a low zinc content, and even doubling the amount of zinc precursor only slightly increased the zinc content in the NCs, again emphasizing the lower reactivity of the zinc precursor relative to the other precursors. The band gap was estimated to be 1.7 eV, different from the previous reports of 1.0−1.5 eV, but the authors mainly attributed the difference in the band gap to the small NC size. In terms of CZTSe morphologies, elongated nanostructures such as CZTSe nanowires have only been attained in solvothermal approaches, by using CuSe nanowire bundles as self-sacrificial templates followed by their reaction with zinc and tin precursors,1031 or by using anodic aluminum oxide (AAO) as a hard template.700 Considerably long reaction times were involved in these solvothermal approaches, where the reactions were carried out for 40−70 h at either 190 or 230 °C. There is only one report of CZTSe nanosheets (Figure 31e) formed in a solvothermal approach, where the thickness of the nanosheets was found to be tunable from 20 to 100 nm by decreasing the amount of ethylenediamine (EDA) solvent.685 Reports of other elongated 1D or 2D nanostructures of CZTSe have remained elusive in the literature to-date. 5.3.4.3. Phase-Selective Synthesis of CZTSe NCs. Phase selectivity from the stannite/kesterite phases to the wurtzite 5944


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Table 14. Overview of the Experimental Procedures for the Synthesis of CZTSe NCs Phase


Method

Temp/Duration


Ref

STN

HI

[100 °C], [295 °C], 295 °C/ 5 min [100 °C], 240 °C/2 h 250 °C/2 h 200 °C/unspecified 230 °C/90 min [100 °C], 280 °C/30 min 280 °C/30 min [180 °C], 180 °C/30 min 190 °C/40 h

20 nm/polyhedral

1027

STN KS STN STN MC STN STN STN

CuCl, ZnCl2, HDA, ODE, [SnCl4/heptane], [TOP/ Se] [CuAc2, ZnCl2, SnCl2, OLA], Se, OLA CuCl2, ZnCl2, SnCl2, Se, TOPO, OLA CuCl2, ZnCl2, SnCl4, TEA, Se CuAc, ZnI2, SnAc4, OLA, Se/OLA CuCl, ZnAc2, SnCl2, OLA, [SeO2/ODE] CuAc2, ZnAc2, SnCl2, Se, OLA CuCl2, ZnCl2, SnCl4, DDT, OLA, [Se/DDT/OLA] CuSe NWs, ZnAc2, SnCl2, TEG

348 1028 890 1029 346,1023 1024 650 705

STN

CuCl, ZnCl2, SnCl2, Se, EDA, AAO template

230 °C/70 h

17 nm/triangular, plates 35−40 nm/flower-like aggregates 23.5 nm/aggregated 18 nm/irregular 25 nm/cubelike 20 nm/irregular 3.4 nm/spherical 200−400 nm × 1 μm/nanowire bundles 200 nm × several μm/nanowires

700

STN WZ

CuCl2, ZnCl2, SnCl4, Se, EDA Cu-oleate, Zn-oleate, Sn(2-Eth)2, OLA, [DPDSe/ OLA] Cu-oleate, Zn-oleate, DPDSe, OLA, [SnCl2/OLA/ OA] Cu-oleate, Zn-oleate, DPDSe, OLA, [SnCl2/OLA/ OA] Cu(acac)2, ZnAc2, DPDSe, OLA, [SnCl2/OLA/OA] Cu(acac)2, ZnAc2, SnAc4, TOPO, OA, OLA, [DPDSe/OLA] [CuCl2, ZnAc2, SnCl2, OLA], selenourea, OLA

210 °C/15 h [190 °C], 255 °C/40 min

20 nm × 1−3 μm/nanosheets 19.3 nm/quasi-spherical

685 349

HI

[250 °C], 230 °C/60 min

15−80 nm/irregular, hexagonal

1030

HI

[230 °C], 250 °C/60 min

10−45 nm/irregular, hexagonal

1030

HI HI

[230 °C], 250 °C/60 min [155 °C], 270 °C/20 min

unspecified/irregular, hexagonal unspecified/irregular, hexagonal

1030 492

HI

[180 °C], 250 °C/30 min

13.5 nm/irregular

1032

KS, WZSTN WZ-STN WZ WZ ZB

HI HU HU HU HI HU HI CTM, ST ST, CTM ST HI

band gap toward the optimum value of 1.5 eV. Moreover, a reduction in the concentration of Sn in CZTGeSSe has been postulated to modify the deep traps associated with this element, potentially leading to improved material quality.1036,1039 Agrawal’s group made significant developments in the synthesis and application of Ge-containing Cu chalcogenide NCs, by developing hot-injection approaches to partially or completely replace Sn in CZTS and form Cu2ZnSnGeS4 (CZTGeS) or Cu2ZnGeS4 (CZGeS) NCs, respectively.369,1034,1035 In addition, the Ge/Sn ratio could be tuned by varying the relative portion of the GeCl4 and Sn(acac)2(Cl/Br)2 precursors.369,1034 In the synthesis, the germanium tetrachloride (GeCl4) precursor was injected into the reaction flask at 128 °C, after which a sulfur/ OLA injection was performed at 160°C, and NC growth was allowed to proceed at 280 °C for 1 h.369 It was noted that the injection temperature of the sulfur precursor played a critical role in the crystal stoichiometry, with a sulfur injection at either 130 °C or 225 °C resulting in Zn-deficient CZGS NCs, even when a 20% excess of the Zn precursor was used. An alternative Ge precursor, Ge(gly)2(H2O)2, was investigated for the synthesis of CZGeS NCs in a separate report and was prepared by heating an aqueous solution of germanium oxide (GeO2) and glycolic acid (glyH2).364 This precursor had the added advantage that it was both air and moisture stable, as opposed to the GeCl4 precursor, which was highly moisture sensitive and prone to hydrolysis. However, it proved difficult to obtain single-phase CZGeS NCs in this heating-up approach, even after changing the reaction temperature, time, and precursor concentrations, with XRD indicating that both tetragonal and orthorhombic phases were present in the final product. This suggested that the formation pathway to CZGeS NCs, compared to CZTS, was more susceptible to the formation of mixed phase products. This observation was consistent with predictions obtained from theoretical calculations, which showed that the energy difference between the tetragonal and orthorhombic phases of CZGeS was smaller than that of CZTS.249

Cu2Zn0.03Sn1.10Se2.98, with the low zinc content indicating that the Zn precursor (ZnAc2) showed much lower reactivity than the Cu and Sn precursors (CuCl, SnCl2). Based on the low zinc content, the obtained nonstoichiometric NCs were expected to accommodate the crystal structure of monoclinic Cu2SnSe3, rather than that of the kesterite or stannite phase of CZTSe. 5.3.5. Other Quaternary Compositions. The chemical and structural flexibility of the quaternary compounds has enabled new materials to be developed in the form of NCs. In terms of the Cu2-II-VI-VI4 compounds, CZTS has been the most extensively studied material combination, but the replacement of Zn and Sn with other group II and group IV materials, respectively, has allowed for the evolution of new material combinations. These materials can collectively be labelled as Cu2MIIMIV(S/Se)4 compounds, where MII = Zn, Cd, Hg, Fe, Co, Ni and MIV = Ge, Sn. In recent years, reports to Cu2ZnGeS4,364,365,369,1033 Cu2ZnGeSe4,347,370 Cu2CdSnSe4,357,358 and Cu2HgSnSe4373 NCs have emerged and have opened up new ways for designing novel functional materials for PV and thermoelectric applications. The incorporation of high spin, first row transition metal ions, such as Co, Mn, Ni, and Fe, to form Cu 2 CoSnS 4 , 37 5 −3 7 8 Cu2MnSnS4,382 Cu2NiSnS4,376,381 and Cu2FeSnS4385,386 NCs, has also been investigated, which could lead to new magnetic and optically active materials in the field of information technology, specifically for spintronics and magnetic data storage applications. An overview of the experimental procedures for the synthesis of these new quaternary material combinations is provided in Table 15. 5.3.5.1. Quaternary NCs Containing Germanium. The introduction of Ge into the CZTS system is viewed as a potential method for improving PV device efficiency.369,1034−1036 It is an attractive route, as it replaces Sn, the rarest element in CZTS, and also allows the band gap to be tuned between 1.5 eV (for CZTS) and 2.05−2.25 eV (for CZGeS).932,1037,1038 As the band gap of CZTS is lowered to ∼1.1 eV during the selenization process, the controlled addition of Ge permits an increase in the 5945


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Table 15. Overview of Experimental Procedures for the Synthesis of Other Quaternary NC Combinations Size (D × L) / Morphology

Material


Method

Temp/Duration

Cu2ZnSnGeS4

Cu(acac)2, Zn(acac)2, Sn(acac)2Cl2, OLA, [GeCl4/ OLA], [S/OLA] Cu(acac)2, Zn(acac)2, OLA, [GeCl4/OLA], [S/ OLA] CuI, ZnCl2, Ge(gly)2(H2O)2, OLA, CS2, DDT [CuCl2, ZnAc2, GeCl4, OLA], thiourea, OLA Cu(DDTC)2, ZnAc2, GeCl2.dioxane, tDDT, OLA Cu(acac)2, Zn(acac)2, GeCl2·dioxane, S, OLA CuCl, ZnCl2, Ge(MeO)4, S, EDA, AAO

HI

[128 °C], [160 °C], 280 °C/1 h 5−30 nm/irregular

369

HI

[128 °C], [160 °C], 280 °C/1 h 5−30 nm/irregular

369

HU HI HU HU CTM, ST HI

250 °C/30 min [180 °C], 285 °C/30 min 300 °C/1 h 310 °C/2 h 220 °C/36 h

364 1032 365 1033 366

Cu2ZnGeS4

Cu2ZnGeSe4

CuCl, ZnO, HDA, TDPA, ODE, [GeCl4/ODE], [SeO2/ODE] CuAc2, ZnAc2, GeI4, OLA, [DPDSe/OLA] CuCl, ZnCl2, Ge(MeO)4, Se, EDA, AAO Cu2−xSe NCs, GeI4, [ZnCl2/OLA]

Cu2CdSnSe4 Cu2CdSnSe4 Cu2CdSnS4 Cu2HgGeSe4 Cu2CoSnS4

Cu2MnSnS4

Cu2FeSnS4

Cu2NiSnS4

Cu(acac)2, CdAc2·2H2O, Sn(acac)2Cl2, OLA, [SeO2/ ODE] CuCl, CdO, SnCl4, ODE, HDA, ODPA, [SeO2/ ODE] CuCl, CdAc2, SnCl2, HA, CS2, MPA CuCl, HgCl2, HDA, ODE, [GeCl4/ODE], [SeO2/ ODE] CuCl2, CoCl2, SnCl4, thiourea, TMAOH Cu(acac)2, Co(acac)3, SnCl2(C5H7O2)2, OLA, [t-DDT/DDT] CuAc2, MnAc2, SnCl4, OLA, [S or thiourea/OLA] CuCl, MnCl2, SnCl2, HA, CS2, MPA Cu(acac)2, Fe(acac)2, SnCl2, [S/OLA] Cu(acac)2, Fe(acac)3, SnCl4, OLA, [t-DDT/DDT] Cu(acac)2, Fe(acac)3, SnCl4, ODE, OA, [t-DDT/ DDT] CuCl, Ni(NO3)2, SnCl2, HA, CS2, MPA

HI

[150 °C], [295 °C], 260−295 °C/5 min [230 °C], 230 °C/60 min

6.5 nm/irregular 15.9 nm/irregular 13 × 50−80 nm/wormlike 13.6 nm/quasi-spherical 200 nm × several μm/ wires 10−25 nm/tetrahedral

Ref

347,370

CTM, ST CER

220 °C/36 h

HI

[100 °C], 300 °C/30 min

25−30 nm/irregular, cubelike 200 nm × several μm/ wires 0.5−1.5 μm/triangular nanosheets 20−30 nm/quasi-spherical

HI

[285 °C], 285 °C/5 min

10−20 nm/tetrahedral

347,357,359

ST HI

5.7 × 26 nm/nanorods 10-40 nm/tetrahedral

355 373

ST HI

140 °C, 180 °C/2 h [120 °C], [285 °C], 285°C/ 5 min 200 °C/16 h [150 °C], 190 °C/30 min

150 nm/aggregated 16 × 32 nm/nanorods

375 377

HI

[270 °C], 270 °C/30 min

382

ST HI HI HI

140 °C, 180 °C/2 h [280 °C], 280 °C/90 min [150 °C], 210 °C/30 min [150 °C], 210 °C/30 min

9.3 nm/triangular, 23.1 nm/plates 4.9 × 46 nm/rod-like 13 nm/triangular, spherical 20 nm/spheroid 15 × 38 nm/plates

376 385 386 386

ST

180 °C/2 h

unspecified/nail-like

376

210 °C/2 h

371 366 283 358

have not yet been obtained for this composition. To the best of our knowledge, there is only one report of CZGeS nanowires, having diameters of 200 nm and lengths of several micrometers, which were obtained in a solvothermal approach using anodic aluminum oxide (AAO) templates.366 Ibáñez et al. developed hot-injection methods to form CZTSe NCs (Figure 32a) and Ge-containing Cu2ZnGeSe4 (CZGeSe) NCs (Figure 32b), with an unprecedented level of control over the NC size, shape, and composition.347,370 The CZGeSe composition has a direct band gap between 1.21 and 1.63 eV, as determined from experimental data and theoretical measurements.239,1037 The synthesis involved the injection of the Ge precursor (GeCl4 in ODE) at 150 °C followed by the injection of selenium precursor (selenium dioxide (SeO2) in ODE) at 295 °C and formed 10-25 nm NCs in the stannite phase after a 5 min growth time.370 It was noted that the partial replacement of Zn ions by Cu ions led to a substantial increase in the material’s electrical conductivity. CZGeSe NCs in the wurtzite phase were synthesized in a hot-injection approach using a different selenium precursor, diphenyl diselenide (DPDSe) in OLA, which was injected at the lower temperature of 230 °C and formed 25−30 nm cubelike NCs in the wurtzite phase after 60 min.371 The reaction temperature was found to play a crucial role in the resultant crystal phase in this synthesis, in that an increased reaction temperature of 260 °C NCs promoted the formation of a mixture of stannite and wurtzite

CZGeS is known to exist as two types of cation-ordered crystal structures, in that the tetragonal structure is a lowtemperature superstructure of the cubic zinc blende unit cell, and the orthorhombic structure is a high-temperature structure derived from the hexagonal wurtzite unit cell.1040 While the orthorhombic structure is generally only observed in the bulk at temperatures greater than 790 °C,367 Fan et al. reported a heatup colloidal method to synthesize phase-pure, orthorhombic CZGeS NCs under less extensive reaction temperatures, and found that an appropriate combination of precursors was crucial in the sole formation of this metastable phase in solution.365 The synthesis involved the use of copper diethyldithiocarbamate, Cu(DDTC)2, ZnAc2, GeCl2·dioxane, OLA, and t-DDT, and formed irregular, wormlike NCs after the reaction was kept at 300 °C for 1 h. The growth mechanism to form orthorhombic CZGeS NCs was also proposed, in which monoclinic Cu1.75S NC seeds nucleated initially in the reaction followed by the fast diffusion of Zn2+ and gradual incorporation of Ge4+ into the existing Cu1.75S nuclei. An earlier report had observed the formation of wurtzite-phased CZGeS NCs using CuCl2, ZnAc2, GeCl4, thiourea, and OLA in a hot-injection approach, but the synthesis generated irregular-shaped NCs which had a broad size distribution and a dominant wurtzite phase and minor zinc blende phase.1032 Overall, there are no reports of CZGeS NCs with exquisite size and shape control, and in addition, 1D nanorods and 2D nanoplates or nanosheets 5946


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Figure 32. TEM images and inset XRD patterns of (a) Cu2ZnSnSe4 (CZTSe) NCs, (b) Cu2ZnGeSe4 (CZGeSe) NCs, (c) Cu2CdGeSe4 (CCGeSe) NCs, and (d) Cu2CdSnSe4 (CCTSe) NCs. A photograph of the CCTSe nanopowder obtained from a scaled-up synthesis procedure is provided in the inset in (d). Panels a−d were reprinted with permission from ref 347. Copyright 2012 American Chemical Society.

phonons and an enhanced thermoelectric performance. The convenient structure layering in the stannite phase also allows the electrical conductivity to be decoupled from both the thermal conductivity and the Seebeck coefficient.1020 Inspired by this, Fan et al. developed a hot-injection method to form CCTSe NCs for the first time.358 The synthesis was later extended to form other quaternary chalcogenide NC compositions, such as CCGeSe (Figure 32c) and CCTSe NCs (Figure 32d) with tetradhedral shapes in the stannite phase, and was up-scaled to produce gram-scale quantities of NCs with narrow size distributions and controlled morphologies.347 In particular, the authors found that the presence of alkylphosphonic acids significantly promoted the incorporation of the II ion (i.e. Cd) into the lattice.347 Polarity driven, branched CCTSe NCs were subsequently reported and formed under modified reaction conditions in the presence of phosphonic acids and excess Cd in the solution.359 The chemical composition and crystal structure changed from the seed to the branches in these CCTSe NCs, where the initial seeds had a zinc-blende-like stannite structure with penta-tetrahedral morphology and branched out to form polytypic polypods with wurtzite arms. Changes in the electronic band structures were also observed, in that the branches changed from semiconducting to quasi-metallic character; such a combination is interesting, as it could be used as perfect epitaxial Schottky contacts between the metallic pods and the semiconducting seeds. The sulfur-containing analog, Cu2CdSnS4 (CCTS) NCs has also been obtained in a solvothermal approach and formed nanorods in the wurtzite phase, which had a band gap of 1.4 eV and demonstrated good photoresponse properties for potential thin-film PV applications.355 5.3.5.3. Quaternary NCs with Transition Metal Ions: Co2+, Mn2+, Ni2+, Fe2+. There are a few reports on the synthesis of other semiconducting quaternary NCs with first row, transition

phases, or the formation of a pure stannite phase at 270°C. Thus, this indicated that the wurtzite phase is less thermodynamically stable than the commonly observed tetragonal stannite phase. A cation exchange approach for 2D colloidal CGeSe, CZGeSe, and CZGeSSe nanosheets using presynthesized Cu2−xSe nanosheets as templates was also developed.283 This approach involved the partial cation exchange of Ge4+ ions and Zn2+ ions with Cu+ ions in Cu2−xSe NCs, in which the Ge4+ ions were incorporated by heating the Cu2−xSe NCs with the GeI4 precursor to 210 °C and formed CGeSe nanosheets after 2 h. CZGeSe nanosheets were formed by adopting a similar method, along with the injection of a zinc precursor solution, ZnCl2 in OLA, at 210 °C followed by growth at 210 °C for 2 h. In an attempt to investigate different group II elements in place of Zn, Cu2HgGeSe4 (CHGeSe) NCs with a controlled composition and a polytetrahedral morphology were formed in a hot-injection method.373 Excess amounts of Hg were required to form stoichiometric, single-phase CHGeSe NCs in this synthesis, but they were also experimentally observed to prevent the formation of twin defects by accelerating the overall reaction. By adjusting the cation ratios, control of the electrical conductivity of the NCs could be achieved, in which Cu-rich CHGeSe NCs were observed to have a higher conductivity than stoichiometric CHGeSe NCs. This increase in electrical conductivity was associated with the higher hole concentration, obtained by the replacement of Hg2+ ions by Cu+ ions. 5.3.5.2. Quaternary NCs Containing Cadmium. Novel quaternary chalcogenide compositions containing Cd, such as Cu2CdSnSe4 (CCTSe) and Cu2CdGeSe4 (CCGeSe), have also been developed. In particular, CCTSe has the required attributes to be a potentially excellent thermoelectric material, because intrinsically low thermal conductivities are associated with the diverse bonding types inside this compound. This leads to a naturally distorted structure that scatters more 5947


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XRD patterns, a successive shift of the peaks to lower degrees should be observed with the increase of Se content in the NCs, due to the larger atomic size of Se than S and an increase in the size of the unit cell. Several reports have been developed to form Cu2−xSySe1−y NCs.546−548,1042−1045 The hexagonal and cubic modifications of Cu2−xSySe1−y alloyed NCs were first reported by Wang et al. in a heat-up approach.1042 In this particular report, NC phase selectively was determined through judicious section of the solvents, in that DDT formed the hexagonal phase and a mixture of OLA and DDT formed the cubic phase. The compositions of both the hexagonal and cubic NCs were tuned by adjusting the amounts of the selenium precursor, diphenyl diselenide (DPDSe), and the sulfur precursor, dodecanethiol (DDT). Optical measurements indicated that the band gap of the Cu2−xSySe1−y NCs could be tuned from 1.44 to 1.85 eV by modulation of the S content. Another interesting observation was that the band gaps could also be tuned by varying the crystal structure, where the hexagonal Cu2−xSySe1−y NCs had a band gap of 1.79 eV, but the band gap of cubic Cu2−xSySe1−y NCs (with a similar Cu/S/Se ratio) decreased to 1.53 eV. However, the authors stated that the reasons for such a difference in the absorption, between the hexagonal and cubic NCs with similar composition, were unknown. Dilena et al. subsequently developed a hotinjection route to form the hexagonal and cubic modifications of Cu2−xSySe1−y NCs, using elemental sulfur and selenium powders as the chalcogen precursors.546 The plasmonic response of the NCs was found to be dependent mainly on the Cu stoichiometry and, surprisingly, was similar to that of Cu2−xSe even for the sulfur-rich alloy NCs. Swihart and coworkers followed on from these reports by presenting a general method for broad tuning of the localized surface plasmon resonance (LSPR), by varying the S/Se ratio and oleic acid (OA) ligand concentration.547 This method demonstrated LSPR absorbance peak tunability from 975 nm up to 1650 nm (Figure 33 a,b) and provided flexibility for using plasmonic Cu2−xSySe1−y NCs in photothermal therapy, in photoacoustic imaging, and as plasmonic photodetectors. Specifically, the LSPR was observed to shift to longer wavelengths with increasing sulfur content and with an increasing concentration of OA in the synthesis. More recently, an DDT and OA solvent combination has been employed to synthesize Cu2−xS NCs, alloyed Cu2−xSySe1−y NCs, and the tellurium derivative, Cu2−xSyTe1−y NCs, in a heat-up approach involving Cu(acac)2, elemental S, Se, or Te powders.548 The synergetic effect of DDT and OA was responsible for activating the chalcogen and forming highly monodisperse NCs, with sizes tunable in the range from 2−3 to 10 nm. The samples showed similar optical properties to those reported for pure binary Cu chalcogenides; i.e. at approximately 1800 nm (1270 nm after prolonged oxidation) for Cu2−xS; at 1300 nm for Cu2−xSySe1−y; and at 800 nm for Cu2−xSyTe1−y NCs. In agreement with the observations of Wang et al.1042 and Dilena et al.546 on alloyed Cu2−xSySe1−y NCs, the plasmonic properties of the NCs strongly depended not only on their size, shape, and surrounding media, but also on their crystal structures. Therefore, the degree of variability of the x value (stoichiometry) was restricted by the crystal phase during the oxidation process. While the aforementioned methods all formed spherical Cu2−xSySe1−y NCs, alloyed nanowire bundles1045 and nanoplate morphologies1043,1044 of Cu2−xSySe1−y have also been reported. 5.4.2. Copper Indium Sulfur Selenide (CuIn(S1−xSex)2). The incorporation of indium to form CuIn(S1−xSex)2 NCs with

metal ions. For example, cobalt (Co) has been used in place of Zn to form Cu2CoSnS4 (CCoTS) NCs, which have a band gap of 1.58 eV and display good photoresponse behavior, thus suggesting their potential use as absorber layer for thin-film PV.377 CCoTS NCs with the stannite structure have been reported in numerous solvothermal approaches,375,376,378,1005 whereas CCoTS nanorods with the wurtzite structure have been developed in a hot-injection approach.377 An interesting finding was that increasing the reaction temperature from 190 to 380 °C, or annealing the NCs at 400 °C, caused a change in the crystal phase from wurtzite to pure stannite. This was accompanied by a change in the morphology, in that the CCoTS NCs evolved from uniform nanorods to large size particles with increasing temperature, and this morphological evolution was attributed to a decrease in the fraction of capping agent bound to the NCs at higher temperatures and to changes in the crystal phase. Cu2FeSnS4 (CFeTS) NCs were synthesized by a hot-injection method and formed tetragonalstructured NCs, which had a band gap of 1.28 eV and a stable photoelectrochemical response, thus indicating their potential for PV applications.385 A subsequent report presented a method to form CFeTS NCs with tunable crystal phase, where zinc blende NCs were favored at a higher reaction temperature (310 °C) and had a band gap of 1.46 eV, and wurtzite NCs were synthesized at the lower temperature of 210°C and had a band gap of 1.54 eV.386 The larger band gap for the wurtzite-derived structure is consistent with recent electronic structure calculations on quaternary chalcogenide semiconductors.249 The incorporation of Mn has also been investigated and was used in place of Zn to form Cu2MnSnS4 (CMnTS) NCs, which had a zinc blende phase if sulfur powder was used or crystallized in the wurtzite phase if thiourea was used.382 The CMnTS NCs had a band gap of 1.1 eV and exhibited weak ferromagnetic behavior at low temperature (2 K). The different magnetic properties of four types of NCs were assessed in an interesting paper by Pan and co-workers, which revealed that the CMnTS NCs and Cu2NiSnS4 (CNiTS) NCs exhibited superparamagnetic behavior at low temperature, whereas the CCoTS and CFeTS NCs showed ferromagnetic behavior.376 Partially doped CZTS NCs with Mn2+, Co2+ and Ni2+ ions and with different dopant concentrations were synthesized in a separate report, but no insight into the effect of each dopant on the magnetic properties was provided.1041 5.4. Semiconductor Nanocrystals with Two Chalcogens

In this subsection, Cu chalcogenide NCs that contain two anionic chalcogens are discussed, in which the two chalcogens are typically sulfur and selenium. This gives rise to NC compositions, which range from ternary Cu2−xSySe1−y, to quaternary CuIn(S1−xSex)2 and quinary Cu2ZnSn(SxSe1−x)4 (CZTSSe). As these are the most popular NC compostions in the literature, a subsection is dedicated to each of these materials in the following three subsections. Other quaterary NC combinations including Cu2Ge(S3−xSex), Cu2Sn(SxSe1−x)3, and CuSb(SxSe1−x)2 have also been investigated, although they have not been synthesized to the same extent as the Cu2−xSySe1−y, CuIn(S1−xSex)2, and Cu2ZnSn(SxSe1−x)4 (CZTSSe) NC compositions. An overview of the methods to form a variety of NC compositions with two chalcogens is outlined in Table 16. 5.4.1. Copper Sulfur Selenide (Cu2−xSySe1−y). Ternary alloyed copper sulfur selenide (Cu2−xSySe1−y) NCs provide an effective way to finely tune the optical band gap, by controlling the S/Se chalcogen ratio in the resultant NCs. In terms of the 5948


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Table 16. Overview of the Experimental Procedures for the Synthesis of NCs with Two Chalcogens Material


Method

Temp/Duration


Ref

Cu2−xSySe1−y Cu2−xSySe1−y Cu2−xSySe1−y Cu2−xSySe1−y Cu2−xSyTe1−y Cu2−xSySe1−y

CuCl2, DPDSe, DDT CuCl, OLA, [S, Se, ODE] CuCl, OLA, [S/OLA, Se/OA] Cu(acac)2, S, Se, DDT, OA Cu(acac)2, S, Te, DDT, OA Cu(NO3)2, S, Se, NaOH, H2O

HU HI HI HU HU HU

225 °C/10 min [300 °C], 300 °C/5 min [225 °C], 205 °C/2.5 min 180−200 °C/30 min 200−220 °C/30 min 170 °C/6−8 h

1042 546 547 548 548 1045

Cu2−xSySe1−y Cu2−xSySe1−y CuInSxSe1−x CuInSxSe1−x CuInSxSe1−x

CuCl, octadecanol, OLA, [S, SeO2, ODE] Cu(NO3)2, S, Se, NaOH, H2O CuCl, InCl3, S, Se, OLA Cu(acac)2, In(acac)3, ODE, DDT, [TBP/Se] CuI, InAc3,OLA, DDT, [TOP/Se]

HI HU HU HI HI

CuInSxSe1−x

CuI, InAc3,OLA, DDT, [Se/OLA/DDT]

HI

2.9−6.6 nm/pyramidal

1049

Cu2Ge(S3−xSex)

Cu(acac)2, OLA, [GeCl4], [S/OLA]

HI

11−19 nm/quasi-spherical

281

Cu2Sn(SxSe1−x)3 CuSb(SxSe1-x)2 CuSb(S2−xSex)

CuCl, SnBr2, S, Se, OLA, DDT Cu(NO3)2, SbCl3, OLA, DDT, SeO2 Cu(acac)2, SbCl3, OLA, [DPDSe/tDDT/ DDT] Cu(acac)2, SnBr2, OLA, [GeCl4/ODE], [S/OLA/Se/ODE] CuCl, InCl3, GaCl3, S, Se, OLA [Cu(acac)2, ZnAc2, SnAc4, OLA], [S/Se/ NaBH4/OLA], TOPO [CuSt, ZnSt, SnSt, OLA], thiourea, Se, OLA, ODE [Cu(acac)2, Zn(acac)2, SnCl2, OLA], S, Se, OLA CuI, ZnAc2, SnCl2, OLA, DDT, [DPDSe/ OLA] CuI, ZnAc2, SnCl2, OLA, DDT, [DPDSe/ OLA] Cu(acac)2, ZnAc2, SnAc4, TOPO, ODE, [DPDSe/DDT] CuCl, ZnAc2, SnAc4, TOPO, HPA, TDPA, OLA, [DPDSe/DDT]

HU HU HI

unspecified/irregular 400 nm × 2 μm/nanosheets 45−50 nm × 10−47 μm/ mesobelts 12−24 nm/quasi-spherical

949 958 220

HU HI

[130 °C], 250 °C/30 min 100 °C/10 h 265 °C/90 min [180 °C], 220 °C/60 min [220 °C], 230 °C/ 20−30 min [170−210 °C], 230 °C/ 10−60 min [125 °C], [160 °C], 280 °C/2 h 240 °C/1 h 200 °C/1 h [250−255 °C], 250−255 °C/10-30 min [120 °C], [160 °C], 280 °C/2 h 265 °C/90 min [325 °C], 285 °C/5 min

8.2−9.4 nm/spherical 10.2−11.6 nm/quasi-spherical 4.7−8.8 nm/spherical 3−8 nm/spherical 3−8.5 nm/spherical 200−250 nm × several μm/ nanowire bundles 15 nm/nanoplates 100−200 nm/nanoplates 15−17 nm/quasi-spherical 1−3 nm/quasi-spherical 4.2 nm/triangular

15−20 nm/irregular faceted 7.8−11 nm/irregular

1046 515

HI

[270 °C], 270 °C/1 h

11−13 nm/quasi-spherical

1053

HI

[130 °C], 240 °C/2 h

8−12 nm/irregular

1052

HI

[180 °C], 280 °C/30 min

17−26 nm/irregular

1055

HI

[180 °C], 240−320 °C/1 h

13 × 22 nm/rugby-ball like

185

HI

[155 °C], 270 °C/15 min

11 nm/quasi-spherical

492

HI

[155 °C], 270 °C/15 min

8 × 12 nm/ellipsoid

1056

Cu2Ge1−xSnx(S3−ySey) CuIn1−xGax(SySe1−y)2 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4 Cu2ZnSn(SxSe1−x)4

HI

1044 1043 859 1047 1048,1050

1051

a similar hot-injection method1048−1050 but investigated their use in quantum dot sensitized solar cells (QDSSCs).1050 Other indium-free combinations including Cu2Ge(S3−xSex),281 Cu 2 Sn(S x Se 1−x ) 3 , 94 9 CuSb(S x Se 1− x ) 2 NCs 22 0, 95 8 and Cu2Ge1−xSnx(S3−ySey)1051 NCs have also been investigated. 5.4.3. Copper Zinc Tin Sulfur Selenide (Cu2ZnSn(SxSe1−x)4). There are several reports on the synthesis and band-gap tuning of Cu2ZnSn(SxSe1−x)4 (CZTSSe) NCs todate, which are based on the adjustment of the S/Se chalcogen ratio. Tetragonal structured CZTSSe NCs were developed in three similar methods by tuning the relative amounts of the sulfur and selenium powders, but there are some discrepancies in terms of the reported band gaps of the NCs.515,1052,1053 The reliable determination of their band gap is also hindered by the fact that Cu chalcogenide NCs do not show a pronounced absorption edge, which presents an issue for Tauc plots in that defined inflections are required for precise determination of the band gap.492 For example, Riha et al. developed a hot-injection method to form compositionally tunable CZTSSe NCs, which involved the simultaneous injection of the cation and anion precursors into a flask containing TOPO at 325 °C, followed by growth at 285 °C for 5 min.515 The band gap of the CZTSSe NCs synthesized in this method was controlled in the range of 1.54 eV for pure CZTS, to 1.47 eV for pure CZTSe (Figure 33e). Ou et al. synthesized CZTSSe NCs in a modified

controllable S/Se ratio was investigated by Tuan and co-workers,859 where the S/Se ratio could be tuned across the entire composition range of x from 0 to 1. While the S and Se powder were not soluble in OLA at room temperature, these precursors immediately melted above their melting points (i.e. Se at 220 °C; S at 120 °C) in the developed heat-up method, which was conducted at 265 °C for 90 min. A series of CuIn(S1−xSex)2 NCs with tunable band-gap energies in the range between 0.98 and 1.46 eV (Figure 33c) were formed. Moreover, the method was up-scaled to form gram scale quantities of CuIn(S1−xSex)2 NCs, demonstrating its potential for providing ultralarge quantities (i.e. kilograms) of band-gap tunable NCs. The same group subsequently incorporated another element, gallium, and formed quinary CuIn1−xGax(SySe1−y)2 NCs in a gram scale, heat-up synthesis with band gaps tunable in the range of 0.98−2.40 eV (Figure 33d), wider than that of any of the quaternary NC compositions.1046 Panthani et al. subsequently reported on the synthesis of ZnS coated CuIn(S1−xSex)2 QDs with bright fluorescence in a hot-injection approach.1047 The ZnS shell increased the PL quantum yield from 10 to 40%, and their application as fluorescent contrast agents was also tested, where the ZnS coated CuIn(S1−xSex)2 NCs were used as an imaging agent to track a microparticle-based oral vaccine administered in mice. Klimov and co-workers synthesized CuIn(S1−xSex)2 QDs using 5949


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Figure 33. (a−b) TEM image of Cu2−xSySe1−y NCs with broadly tunable NIR localized surface plasmon resonance (LSPR). Reprinted with permission from ref 547. Copyright 2013 American Chemical Society. (c) Plot of the optical band gaps of CuIn(S1−xSex)2 NCs versus the composition x, with the inset showing a photograph of the gram-scale quantities achieved in the synthesis. Adapted with permission from ref 859. Copyright 2011 American Chemical Society. (d) UV−vis−NIR spectrum of CuIn1−xGax(SySe1−y)2 NCs with varying In/Ga(x) and S/Se (y) ratios. Reproduced with permission from ref 1046. Copyright 2011 Royal Society of Chemistry. (e) Plot of the band-gap energies of Cu2ZnSn(SxSe1−x)4 NCs where a linear decrease is observed with increasing Se content. Reproduced with permission from ref 515. Copyright 2011 American Chemical Society. (f) Low-temperature photoluminescence (PL) emission of Cu2ZnSn(SxSe1−x)4 NCs with varying values of x. Reproduced with permission from ref 492. Copyright 2013 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

follow up study, by using both solution-based and solid-statebased phase transformations, respectively.1057 In this report, the authors observed that the phase transformation pathways were selective, in which wurtzite CZTSSe NCs were converted into the zinc blende phase in a post-treatment solution-based approach, whereas the conversion of wurtzite NCs into kesterite grains was achieved by employing a thermal-annealing, solid-state approach. Singh et al. reported the first low temperature PL emission study (Figure 33f) of wurtzite CZTSSe NCs, with a compositionally tunable band gap ranging between 0.9 eV for Cu2ZnSn(S0.2Se0.8)4 and 1.4 eV for CZTS, that directly correlated to the sulfur-to-selenium precursor ratio.492

phosphine-free, hot-injection approach, but achieved band gaps in a much larger tunable range from 1.5 eV for pure CZTS and 1.0 eV for pure CZTSe.1053 Experimental and first-principle calculations have also been investigated on the band gaps of CZTSSe alloys,185,1054 where the band gaps increased from 1.0 to 1.5 eV with an increase in S content. This suggested that the band gaps reported by Ou et al. on tetragonal CZTSSe NCs were closer to those obtained by first-principle calculations. Compared with the tetragonal derived CZTSSe alloys, wurtzite-derived alloys are more difficult to synthesize due to their two-fold energetic metastability, in that the wurtzite phase is metastable relative to zinc blende-derived structures. Fan et al. published the first report of CZTSSe NCs with wurtzitederived structure by using DDT and diphenyl diselenide as the sulfur and selenium sources, instead of elemental S and Se powders.1055 Similarly, the authors found that the band gap of the wurtzite-derived CZTSSe NCs ranged from 1.0 to 1.5 eV, which showed good agreement with the predicted values from ab initio calculations. The same group also published a method to form linearly arranged polytypic NCs, where each NC consisted of two zinc blende-derived ends and a wurtzitederived core.185 Polytypic CZTSSe NCs with different wurtzite and zinc blende phase ratios were obtained by tuning the reaction temperature from 240 to 320 °C, with nucleation in the wurtzite-derived structure preferred at relatively low temperatures (240 °C) and the zincblende-derived structure favored at higher temperatures (320 °C). In a separate report by Singh et al., polytypic CZTSSe NCs with phase and shape control to form dots, ellipsoids, and arrow- and bullet-shaped NCs were obtained by the judicious selection of ligands and metal precursors in the synthesis.1056 In addition, single-phase wurtzite CZTSSe NCs were obtained in this report by tuning the reaction chemistry. A selective structural phase transition study of metastable wurtzite CZTSSe NCs into more stable phases, such as zinc blende and kesterite, was reported in a

6. NANOCRYSTAL INTERACTIONS AND ASSEMBLY STRATEGIES 6.1. Assembly Introduction

The complete level of size and shape control achievable in Cu chalcogenide NCs makes them optimal building blocks for their hierarchical assembly into useable architectures, where their functional properties can be collectively harnessed. This has been achieved via self-assembly, and prototypical examples of self-assembled NC structures include close-packed NC superlattices and supercrystals that exhibit either a short-range or long-range periodicity.1058−1060 Beyond that, more complex superstructures starting with NCs of different compositions (metals, metal oxides, and groups II−VI, III−V, and I−II− IV-VI semiconductors) and shapes (i.e. spheres, cubes, rods, wires, prisms, plates, tetrapods) into different arrangements, such as body-centered cubic (bcc), face-centered cubic (fcc), and hexagonal close packing (hcp), have also been demonstrated.1058,1061−1068 By manipulating NC building blocks into higher order assemblies, there exists the opportunity to engineer unique properties that are not only dependent on properties intrinsic to the NC but also on extrinsic attributes, 5950


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Figure 34. TEM images of self-assembled superstructures of Cu chalcogenides from (a−c) simple hcp packing to (d−f) long range order superlattices in corncob and columnar (ribbon) structures. Panel a is reprinted with permission from ref 253. Copyright 2010 American Chemical Society. Panel b is reprinted with permission from ref 666. Copyright 2011 Royal Society of Chemistry. Panel c is reprinted with permission from ref 474. Copyright 2008 American Chemical Society. Panel d is reprinted with permission from ref 767. Copyright 2008 Elsevier Inc. Panel e is reprinted with permission from ref 763. Copyright 2010 American Chemical Society. Panel f is reprinted with permission from ref 1078. Copyright 2012 Royal Society of Chemistry.

In this review, while the focus is on Cu chalcogenide NCs, the overall understanding of assembly is not confined to a particular composition but is defined by the different internanocrystal forces that drive the self-assembly process. The basic understanding of these forces and how they are influenced by the experimental methods used in assembly is a prerequisite for the predictive design of new nanoparticle-based assemblies. The progress of these assemblies in terms of the internal forces between the NCs, with particular emphasis on Cu chalcogenide NCs, is highlighted in this section.

such as NC arrangement, that influence their collective behavior.1069−1071 Several examples include advances in charge transport, surface-enhanced Raman scattering, photonics, photovoltaics, magnetic, biomedical, and catalysis.631,1069,1072−1074 One subset of NC assembly that has garnered increasing interest is the assembly of anisotropic NCs.1075,1076 In typical isotropic systems (i.e. spherical NCs), only positional ordering is usually considered. However, for anisotropic shapes and faceted NCs, orientational ordering presents an additional variable. Apart from their inherent shape anisotropy, which is a barrier to the manifestation of long range order, this anisotropic feature also introduces different surface chemistries along different facets (due to different surface energies) and stronger interparticle interactions, which can be utilized in the selfassembly process. Over the years, many strategies have been explored and developed to induce the self-assembly of anisotropic NCs into tunable superstructures over device scale areas. Particularly, orientational control over the assembly of anisotropic NCs has been achieved by manipulating the interparticle interactions and NC surface chemistry, using templated assembly and application of external stimuli.1064,1076,1077 The assembly possibilities for Cu chalcogenide NCs have been highly intuitive, due to flexibility in tuning their shape (spherical, platelet, cube, rod, wire, biyramids) and crystal structures (djurleite, hexagonal covellite, zinc blend, wurtzite, chalcopyrite). These factors influence the surface chemistries along different facets in both shape and crystal structure and, hence, can be exploited in the assembly process. This allows the possibility to assemble Cu chalcogenide NCs in various superstructures, ranging from traditional hexagonal close-packing or face-centered cubic to long range stacking of nanoplate superlattices (as ribbons or columnar assemblies), as shown in Figure 34.

6.2. Interactions between Nanocrystals

6.2.1. van der Waals Interactions. Self-assembly in NCs is mainly governed by the intrinsic forces present in the NC system.1079 Generally, these forces are van der Waals interactions and are necessary to drive the assemblies into superlattices. van der Waals interactions originate from the dipole−dipole interactions present in the materials and can be classified depending on whether the dipoles are permanent or induced (from the polarizability of the nanocrystals).1080−1083 The organic ligand capped NCs exhibit an attractive force component of the van der Waals interactions, and dispersion forces, which arise due to the induced dipoles in NCs when the spontaneous fluctuation of the electronic cloud in one atom causes fluctuation of electrons in its neighboring atom. Although Keesom and Debye forces are stronger forces that van der Waals, with the exception of highly polar materials, it is the London dispersion forces that account for most of the van der Waals attraction, which is responsible for assembly. Both Keesom and Debye forces involve at least one permanent dipole, which indicates the existence of electric dipole moments in NCs, due to the chemical environment of the surface ligands and the crystal structure of self-assemblies. The strength of 5951


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interactions decay rapidly with r−6, thus differentiating it from the typical colloidal behavior. While the dipole−dipole interaction energy between two nanocrystals under applied electric field can be given by eq 4:

dispersion forces in nanocrystals depends on the polarizability of the NCs and can be given by eq 1: disp Eab ≈

−3IaIb α aα b 2Ia + Ib r 6

(1) d−d = E12

where Ia and Ib stand for first ionization potentials of atoms, r is the intermolecular distance, and α is the dipole polarizability. As evident from the above equation, dispersion forces are inversely proportional, r6 , which makes the strength of the van der Waals force act as a small but significant perturbation to other long-range interactions, such as electrostatic interactions. The interaction energy of dispersion forces can be calculated according to Hamaker’s theory, which is based on the assumption that the dispersion potential between two colloidal NCs can be presented as a summation of the dispersion interactions between pairs of atoms located within the two NCs. The expression for dispersion interaction energy (E) between two spheres can be approximated by the additive Hamaker approach: E=−

(4)

where p1 and p2 are the permanent dipole moments of the two NCs, r1 and r2 are the radii, and n is the unit vector between the nanocrystal centers. In the case of colloidal NCs, the dipole moment can be experimentally measured for an anisotropic structure, such as a nanorod. By using transient electric birefringence methods, Li et al. showed that wurtzite CdSe nanorods possessed a large dipole moment on the order of 100 to 200 D and scaled with the volume.1081 This is due to small crystallographic deviations of wurtzite structure from the ideal structure. Further based on this observation, the dipole moment for other materials (CdS, ZnSe, PbSe, etc.) has been calculated/predicted experimentally and theoretically.1082,1086,1088 Recently, Zhuang et al. has used the anisotropic cell structure for hexagonal Cu2S NCs and found that it has a layered structure (alternation of Cu and S rich layers), in which a Cu5S layer is on one side and a CuS layer terminates the other side, which ultimately leads to inducing a dipole moment along the [001] direction (due to structural asymmetry) as shown in Figure 35.474

2r1r2 2r1r2 AH ⎡ ⎢ 2 + 2 6 ⎢⎣ l + 2r1l + 2r2l l + 2r1l + 2r2l + 4r1r2

⎞⎤ ⎛ l 2 + 2r1l + 2r2l ⎟⎥ + ln⎜ 2 ⎝ l + 2r1l + 2r2l + 4r1r2 ⎠⎥⎦

1 p1 ·p2 − 3(n·p1 )(n·p2 ) 4πε0 (r1 + r2 + l)3

(2)

where the negative sign represents the attractive nature of the interactions, AH is the Hamaker constant for the two interacting NCs, which depends upon the polarization properties of the two nanocrystals and the medium which separates them, with radii r1 and r2, and l is the center to center distance between the two NCs (eq 2).1084,1085 The Hamaker constants depend on the precise structure of NCs, which are not known exactly, but they are generally not the same as that of the bulk material. 6.2.2. Dipole−Dipole Interactions. Semiconducting NCs can have permanent dipole moments, originating from electronic or magnetic association, respectively, and these dipole−dipole interactions, if strong enough, can induce directional assembly or aggregation as the NCs approach each other in the solution.1080 In NCs, the electric dipole moment originates from the non-centrosymmetric nature of the atomic lattice, and these short-range interactions can be scaled with nanocrystal volume.1081,1082,1086 Depending on the arrangement of the facets in the NC solution, the NCs can have net zero dipole moment if it has center of symmetry and hence no assembly, while the lack of symmetry will induce a net dipole moment and can drive the assembly of the nanocrystals.1087 The dipolar moments affect the interactions between NCs through dipolar coupling, and the interaction potential between two dipoles can be given as μ Vdipole = [2 cos θ1 cos θ2 − sin θ1 sin θ2 cos ϕ] 4πεε0r 2

Figure 35. Scheme of the crystal structure of hexagonal Cu2S along the [001] direction. Reprinted with permission from ref 474. Copyright 2008 American Chemical Society.

6.2.3. Electrostatic Interactions. Electrostatic interactions, which can be either attractive or repulsive, strongly affect the self-assembly process of NCs due to their long-range nature. Depending upon the strength of the interactions and solution environment, the NCs can either end up as aggregates or mesoscopic self-assembled structures. When considering these electrostatic interactions, it is convenient to begin with the simplified framework proposed by Derjaguin, Landau, Verwey, and Overbeek. In this formalism, the strength of electrostatic interactions between colloidal particles can be adjusted by controlling the electrolyte dielectric constant, ionic strength, ionic valence and temperature. In high dielectric constant media, for instance, charges dissociate easily and the length of the electrostatic interaction between NCs can be calculated on the basis of the double-layer thickness, as shown in eq 5:1089

(3)

where the dipole moment μ separated by a distance r and θ1, θ2, and Φ are the orientations of the dipole (eq 3).1080 The orientation of the dipole in a NC and the distance between the dipole individually affect the interaction potential in a NC. In turn, the orientation and dipole distance are dependent on the NC size and morphology and can be tuned and varied in rods, cubes, and faceted nanocrystals.1081,1082,1086 The dipole−dipole interactions in semiconductor NCs behave like that of molecules as the interactions scale by r−3 while the van der Waals

k−1 =

5952

ε0εr kBT e 2 ∑i cizi2

(5) DOI: 10.1021/acs.chemrev.6b00376 Chem. Rev. 2017, 117, 5865−6109

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where ε0 is the vacuum permittivity, εr is the dielectric constant of the solvent, kB is the Boltzmann constant, T is the absolute temperature, e is the elementary charge, and ci and zi are the number densities and valences of the electrolyte ions. The value of the screening length (k−1) increases with decreasing electrolyte concentration, which leads to long-range electrostatic interactions between charged colloidal NCs. For considering the interaction between NCs, it is also convenient to introduce an alternate metric called the Debye length, traditionally defined as the distance between two unit charges, where their Coulomb energy equals the thermal energy kT and quantifies the length scale over which a counterion cloud is perturbed by a charged surface. Generally, Cu chalcogenide NC surfaces become charged, either from the dissociation of ionizable surface groups (e.g., acidic groups under basic conditions) or from the adsorption of charged (ionic) species onto a charge-neutral surface.1090 When these colloidal NCs with similar charges interact with each other, the electrostatic repulsions are generated and the electric double layers of NCs overlap with each other. These repulsive interactions prevent the NCs from aggregating and result in solution stability. While in the case of oppositely charged NCs, self-assembly of binary NCs could be obtained and precisely controlled through the concentration and size of the NCs,1090−1092 particularly in NCs where surface charge is distributed asymmetrically or that have permanent electric polarization, where the electrostatic interactions can be directional. The electrostatic interactions along with dipole−dipole attractions play a crucial role in determining the assembly process and can lead to 1D, 2D, or 3D superstructures.1090,1093,1094 The formation of an assembly process in NCs can be further controlled by tuning the electrostatic interaction by the geometric curvature of nanoparticles, as the density and arrangement of surface ligands vary with the morphology of the NCs and, thus, can result in different dimensional superlattices.1093,1095 6.2.4. Depletion Interactions. In colloidal solutions, the NC interactions can be manipulated using small nonadsorbing solute or solvent molecules, called depletants, to generate longrange attractive depletion interactions. The small nonadsorbing depletant molecules are dispersed in the continuous phase of the multicomponent system and can be excluded from the colloid solution. In a large colloidal NC system, small nonabsorbing polymer molecules are widely used as depletants to drive the self-assembly process.1096 The depletion attraction phenomenon in colloids has been explained in different ways. One approach can be explained considering the repulsive interactions of the colloidal and solute particles or solvent molecules that generate the depletion layer around the particles. When colloidal particles approach each other at a distance smaller than the diameter of the solute particles, the depletion layers interact and the total volume available to depletants increases as compared to noninteracting colloids and bulk. This interaction induces an entropically driven attractive depletion force and leads to assembly. The other approach considers that in a system that contains larger colloidal particles and small solute molecules, e.g. polymer chains, as the two larger colloidal particles collide and become close, the small solute molecules are excluded from the space in-between. This exclusion of solute molecule is enough to cause a net movement of solvent out of the depletion region to generate a net osmotic pressure on the colloidal particles, which drives the colloidal particles closer, and this helps in the assembly of the larger particles. If the attractive forces are large enough, the colloidal solution can result in phase separation. The increase in volume ΔV is accompanied by a

decrease in the exclusive region of the polymer, and is therefore available to polymers with increased entropy, as described in eq 6.

ΔFd = kBTnpΔV

(6)

where ΔFd denotes the reduction in free energy, kB is the Boltzmann constant, T is the temperature, and np stands for the number density of small molecules. The reduction of free energy is directly proportional to the increase in overlapping exclusion volumes (ΔV) due to the depletion interactions.1097 The depletion interactions between NCs are shape dependent and can be manipulated by adjusting the size and concentration of solute molecules and solution temperature, which facilitates the control of the long-range ordered NC structures, i.e. into self-assembly. It has been shown in the literature that the semiconducting CdSe/CdS nanorods yield close-packed hexagonally ordered arrays of 2D monolayers through depletion attractions.1096,1098 If the colloidal system contains adsorbing solute molecules, the colloidal particles will remain separated due to the absence of depletion interactions and loss of conformational entropy, and the system will be colloidally stabilized. 6.2.5. Capillary Interactions. Prior to discussing capillary interactions, it is important to first note that these forces are orders of magnitude larger than the aforementioned van der Waals, electrostatic, and dipole−dipole forces. This is evident in their original application to assemble large micrometer colloids where other forces do not provide the necessary attraction for assembly. For nanoscale systems, they remain relevant because they can be used to force the assembly of particles, even in cases where the net balance of the aforementioned forces in a system does not favor assembly. There are three predominant types of capillary forces that can affect the assembly of colloids: bridging, floatation, and immersion.1099 Briefly, the bridging force arises when a section of solvent that spans two particles exhibits a curvature in the direction of the bridge phase and this curvature induces a Laplace pressure that pulls the two particles closer. On the other hand, floatation and immersions forces arise from particle interaction with a liquid phase. Floatation forces arise from a balance of meniscus forces, buoyancy forces, and gravitational forces acting on a colloid on a liquid surface but are actually not pertinent for particles below the micrometer scale, where gravity forces can be largely neglected. Immersion forces arise from wetting of the particle surface by a liquid matrix, usually in the form of an evaporating liquid droplet containing particles. As the matrix evaporates, a meniscus is formed between particles, and wetting of the liquid on the particle surface determines the shape of the meniscus and therefore determines the force, be it attractive or repulsive, between particles.1100,1101 In all three cases, wetting of the liquid matrix on the particle surface determines the direction and extent of force between particles. On the microscopic scale, this translates to differences in particle shape, curvature, and surface properties (i.e. surface chemistry, crystallographic facets). In the ideal case of spherical particles with a smooth surface, the equilibrium contact angle is governed by the surface energies, as dictated by surface chemistry and crystallography at the three-phase interface. When particle surfaces are rough or when particles are nonspherical, however, the system may no longer easily adopt equilibrium contact angles due to pinning of the two-phase interface (air−solvent etc.) at tips and edges.1099 This leads to increased difficulty in capillary force driven assembly of multifaceted and anisotropic particles: the large density of corners and edges 5953


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process and also provides the freedom to manipulate the final superstructure from a simple one-component system to a multicomponent system. As discussed in the synthesis section, the primary importance of the capping ligand is to control the composition, shape, and crystal structure in the solution synthesis of multicomponent Cu chalcogenides. The capping ligand is also crucially important, as it mediates the interaction between the NC and solvent for the purpose of assembly. For example, Li and co-workers showed for the first time that the columnar selfassembly of Cu2S hexagonal nanoplates can be induced by in situ formation of a Sn-X complex (inorganic ligand) during the synthesis by the reaction between tin(IV) bis(acetylacetonate) dichloride (Sn(acac)2Cl2) and 1-dodecanethiol (1-DDT).808 This inorganic complex (Sn-X) not only changes the shape of the Cu2S NC from a spherical particle ((Figure 36a−b), formed when there is no Sn(acac)2Cl2 involved in the reaction) to plates (Figure 36 c) but also influenced their self-assembly. Most nanoplates tend to stack face-to-face to form a columnar superstructure. The distance between the nanoplates is close to 1 nm (Figure 36 d), which further confirmed the presence of a shorter ligand than the 1-DDT molecule (2 nm), as seen in the case of spherical NCs (Figure 36b). This approach has also been used for the columnar assembly of Cu2−xSe NCs.1078

complicates experimental control of any receding three-phase interface during drying, that is responsible for the forces that ultimately lead to assembly. Similarly, when particle surface chemistries are incompatible with the host solvent, the wetting behavior can create repulsive forces that disallow particle− particle interaction to form critical nuclei from which ordered assembly can arise. Apart from the above-mentioned interactions that play an important role in driving and governing the self-assembly process of Cu chalcogenide NCs into superlattices and supercrystals in 1D, 2D, or 3D, other different interactions and forces also contribute and promote the self-assembly process. The entropic effect arises from the steric repulsions of the surface ligands, ordering at high volume fractions, solvophobic interactions between the organic surface ligands and polar solvents, ligand−ligand interactions, and externally applied convection forces.1079,1080,1085 6.3. Role of Surface Ligands in Nanocrystal Assembly

The NC surface plays a significant role in the self-assembly process, owing to the high surface reactivity due to the high surface area and the resultant presence of a high number of dangling bonds, as compared to their bulk counterparts.1090,1102 The surface ligands separate the core of the NC from the bulk solution and direct control over the interaction between the NC and its environment.631,636,1090 The reactivity and composition of NC surface atoms decide their interaction with different ligands and, thus, can be used to manipulate the interparticle interactions. Hence, the choice of ligand for surface passivation is key not only in synthesis of NCs but also for manipulating their properties and the self-assembly process. Ligands serve as a protection layer to avoid direct contact between the NC cores. The long chain organic ligands (hydrophobic in nature) passivate the trap states on the NC surface, by binding to the low-coordinating surface atoms, which helps to solubilize the NCs in nonpolar solvents and form stable colloids. When the NCs are in solution, the ligands keep on adsorbing and deadsorbing on NC surface. The ligand nature, and their interaction with other ligands or solvent molecules, determine the extent of colloid stability. In general, the long chain organic ligands behave like polymer chains in a good solvent. When two NCs approach at a distance smaller than twice the ligand length, the chains compress and yield a repulsive force between the NC to make the dispersion stable. The high density of ligands on the NC surface will make these repulsions stronger, and also the size/molecular weight of ligands, as bulky groups, will cause the steric hindrance. Also, the chain length and chemical composition of the capping ligands can influence the dipole moments of the NCs, their reactivity, and their stability in solution and direct the selfassembly via electrostatic interactions or bonding. It has become ever more important to understand this synergism between the NC surface and the capping ligands, which has significant consequences for the properties of assembled materials, as compared to individual NC building blocks. In the case of monodisperse spherical NCs with an organic ligand covered surface, it is expected that the assembled structure would arrange into a face-centered cubic (fcc) lattice to achieve the maximum packing efficiency.631,1060 Recently, the superlattice structure from a simple fcc lattice to a body centered cubic (bcc) one was obtained by just changing the NC carrier solvent and the ligand coverage.1103 Thus, this has opened a new avenue in understanding the self-assembly

6.4. Strategies of Nanocrystal Assembly

Over the years, different strategies have been employed for the assembly of a wide variety of NCs with four types of assembly processes dominant: (i) Drying mediated assembly; (ii) Assembly in solution; (iii) Assembly at the interface; and (iv) Directed assembly. These are outlined for Cu chalcogenides in Table 17. 6.4.1. Drying Mediated Assembly. Self-assembly of NCs by solvent evaporation/drying has been extensively used to form superstructures (from centimeter to micron size) on a range of substrates.1090,1110 While the whole process seems to be just evaporation of the solvent, the interactions that drive the organization are complex, as discussed (dipole−dipole interactions, electrostatic attraction and repulsion, depletion attraction, steric forces (associated with surface ligands), capillary forces, and the solvent evaporation dynamics during the controlled or uncontrolled solvent evaporation). The control factors for these interactions are solvent dielectric, volatility, polarity, NC charge, NC size and concentration, dispersibility, temperature, and surface ligands.631,1094,1110−1112 As this parameter set is too broad to allow for prediction of the preference for assembly, the approaches typically involve trial and error methods to elucidate the optimal conditions for assembly, and once this is achieved, more control can then be implemented to tune properties such as assembly size, location, and ordering of the assembly. In the process defined by solvent evaporation, the NCs are dispersed in an appropriate solvent (mostly organic, such as toluene, hexane) and are dropcast onto a clean and flat substrate, followed by solvent evaporation. During solvent drying, the local NC concentration gradient (underneath the liquid surface) increases, which reduces the distance between NCs. This initiates various attractive (dipole−dipole) and repulsion interactions (Columbic), which ultimately stabilize or force the NC into a close packed 2D structure at the surface.631,1094 These structures may be further assisted by surface tension, which helps in holding the assembly at the surface. The first attempt to assemble colloidal NCs into a highly ordered superstructure was obtained by simply drying the NC 5954


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Figure 36. (a) Schematic representation for the synthesis and assembly procedure for Cu2S NC (sphere and hexagonal plates). (b) TEM images of spherical NCs synthesized in the presence of 1-DDT, with the inset showing the selected-area electron diffraction (SAED) pattern. (c,d) TEM and HRTEM images showing the columnar self-assembly of Cu2S nanoplate arrays standing edge-on perpendicular to the substrate, respectively. Reprinted with permission from ref 808. Copyright 2010 American Chemical Society.

Table 17. Overview of the Assembly Procedures for the Cu Chalcogenide NCs Assembly Process

Assembly Range

Shape

Material

Ref

Drying Mediated

100 1000 nm

At the Interfaces In Solution

1μmcm2 100 nm2 μm

Directed

1μmcm2

Spherical Rod Platelet Pyramidal Rod Spherical Rod Platelet Rod

Cu2S, Cu2−xS, Cu2−xSe, CuInSe2, Cu2ZnSnS4, CuS, Cu2S, CuInS2, CuIn1−xGaxS2, Cu2ZnSnS4, CuS, Cu2S, Cu2Te, Cu2−xSe, CuInSe2, CuInSe2 Cu−In−Zn-S Cu2S, Cu2−xSe CuInS2, CuIn1−xGaxS2, Cu2CdxZn1−xSnS4 CuS, Cu2S, Cu2−xS, Cu2−xSe CuInS2, CuIn1−xGaxS2, Cu2ZnSnS4

474,765,1104,1105,253,907 771,474,747,666,973,975,995 64,763,765,770,762,770,728 836,907,203 1106 1107,835 975,1108 780,62,70,769,808,808,1109,835,1078 973,975,246

in a glass vial tilted at approximately 45−70°, and the solvent is then allowed to evaporate over the course of several hours at 45−60°C (depending on the boiling point of the NC) in a drying oven. Recently, Murray and co-workers have shown that the tetragonal bipyramid CISe NCs can be assembled into close-packed oriented films (monolayer to multilayer) over large-areas by using a similar setup (Figure 37b−e).203 Lu et. al also used this approach to achieve a vertical assembly of CIS nanorods.666 Many reports have demonstrated the drying mediated assembly for different shaped Cu chalcogenide NCs, which have a wide variety of packing/superstructures, as evident in Table 17. Apart from drying rate and solvent, other key parameters have been shown to influence NC assembly, such as concentration, surface charge, and NC−substrate interactions. These parameters have been extensively discussed in the literature for NC assemblies.631,1058,1064,1075 6.4.2. Self-Assembly at the Interface (Liquid−Liquid and Liquid−Air). Self-assembly of particles at interfaces such as liquid−liquid and liquid−air has been investigated for more than a century. This idea of trapping solid particles at the interface between two immiscible liquids to prevent gravity sedimentation is based upon the Pickering−Ramsden

solution on a substrate (i.e. a TEM grid, silicon nitride membrane, silicon wafer). The well-dispersed NCs experience a short-ranged steric repulsion while in the organic solvent, and there is no significant driving force for the formation of a superstructure. As the solvent is evaporating, the local NC concentration increases and the NCs become compressed together, due to reduction of the free volume available to the NC that leads to the formation of a superlattice structure driven by entropy of the system. The overall entropy of the system (NC assembly) can be further increased by controlling the temperature of the process (based on Helmholtz free energy).631,1113 A higher temperature will not only increase the entropy but also ensure a constant solvent evaporation rate, which is highly desirable for the high-ordered assembly. For example, Murray and co-workers have shown that slowing down the evaporation rate of a NC solution, in a controlled manner, is favorable to increase the overall packing order, uniformity, and size of the superstructure.631,636,1067 Figure 37a shows a schematic illustration of a low-pressure assembly setup that was used in the drying mediated assembly process for a wide variety (shape, size, type) of NCs. Typically, a substrate (i.e. TEM grid, Si wafer) is immersed in the NC dispersion, 5955


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Figure 37. (a) Schematic illustration for drying-mediated assembly setup. (b, c) STEM image and model of a self-assembled monolayer of oriented tetragonal bipyramid CISe NCs. (d, e) TEM and SEM images show the multilayered assembly of CISe NCs, with the inset in d showing the SAED pattern. Panels b−e were adapted with permission from ref 203. Copyright 2013 American Chemical Society.

emulsions phenomenon,1114 and it was later exploited to design new techniques to assemble colloidal NCs into simple (one particle type system) and complex superlattice and superstructures (2D and 3D from binary to quasicrystal type).1066,1115−1117 These techniques not only provide a rich tool for lab scale NC assembly, but their scalability to wafer level substrates suggests possible extension to real world applications. Interfacial assembly (ΔE) is driven by the total reduction of interfacial free energy between two immiscible phases due to particle adsorption on the interface. Depending upon the particle size, ΔE can be higher (micron size) or comparable (nanometer size) to thermal energy (kBT), and this will govern the adsorption of particles at the interface (higher ΔE cause the fluctuations). In general, smaller particles adsorb more weakly than larger ones at the interface, and hence, the assemblies from larger particles are more stable. Apart from particle size, particle shape and surface chemistry also play key roles in the overall assembly process and the stability of final superstructures, as they influence the wettability (change in surface energy) and the interparticle interactions. The important role of the liquid interface is to provide a mobile surface for the particle to move freely and rearrange into their closely packed configuration. The assembly of colloidally synthesized NCs, at interfaces with different degrees of short- and long-range ordering, has been studied for isotropic (i.e. spheres, cubes) and anisotropic NCs (i.e. nanorods, nanowires, nanoplates, nanoprisms) in both metal NCs and semiconductor NCs.1066,1077,1110,1115,1117−1120 In the case of spherical NCs, close-packed assembly is generally favored due to the interparticle interactions. In contrast, for anisotropic NCs, shape and spatially dependent (interface mediated) capillary forces govern the overall assembly, due to undulation of the contact line at the interface. Colloidally synthesized semiconductor nanorods were first shown to selfassemble at the liquid−liquid (oil/water) interface by Russell and

co-workers, who noted that anisotropic NCs behave differently than their spherical counterparts.1120 The assembled structure displayed morphologies ranging from low-density smectic packing at the edges, to columnar structures, and finally to a crystalline-like phase at the center. More importantly, the work hinted at the complex phase diagram (smectic to columnar to crystalline) that governs assembly at the liquid−liquid interface and that different phases can be accessed by controlling parameters such as NC concentration, aspect ratio, and the interfacial energy between liquids and the NC. Recently, Murray and co-workers reported a simple technique of dropcasting the NC dispersion on an immiscible subphase and then controlling the evaporation rate to drive the assembly for a wide range of NC morphologies (spheres, plates, rods, prisms) with complex structures (binary, ternary to quaternary).1066,1115−1118 In this evaporation-controlled assembly, the NC dispersion (in hexane, toluene solvents) was drop-casted in a Teflon well, containing the immiscible solvent surface, such as ethylene glycol, diethylene glycol, and acetonitrile, and the assembly was controlled by tuning the NC concentration and the solvent evaporation rate. The resulting centimeter scale assemblies were suspended on the subphase solvent surface, allowing for easy transfer onto arbitrary substrates such as TEM grids, glass, silicon, and indium tin oxide (ITO) coated glass. In the case of compound semiconductors, Singh et al. has similarly shown that by controlling the nanorod concentration at the liquidinterface, the assembly orientation can be achieved in either lateral or vertical alignments, as shown in Figure 38.1106 For example, Cu-In-Zn-S nanorods dispersed in hexane can be assembled in a Teflon well, on an acetonitrile or diethylene glycol surface, by drop-casting the nanorod solution and allowing the solution to evaporate slowly (Figure 38a). The floating nanorod superstructure film can be transferred to any substrate previously immersed in the subphase. Furthermore, the nanorod concentration also controls the extent of assembly. 5956


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Figure 38. (a) Scheme showing the nanorod assembly at the liquid−air interface setup. (b−d) HRSEM images showing the nanorod assembly obtained at the liquid−air interface for nanorod solutions with different initial rod concentrations. Reprinted with permission from ref 1106. Copyright 2015 American Chemical Society.

At low concentrations, no orientational long-range assembly occurs and the nanorods partially align side-by-side. However, as the concentration is increased, lateral assembly of nanorods becomes more apparent, and finally at significantly high concentrations (∼2 μM), vertical assembly of nanorods was achieved (Figure 38b−d). This progression in the overall assembly process can be rationalized by considering the influence of various forces on the rods, and also the subsequent changes in these forces during the evaporation process at the liquid−air interface. The long-range order in the assembly, leading to periodic nuclei formation (nucleation of the ordered phase) and the dynamic nature of nanorods (attachment and detachment), resulted from a balance between the interparticle repulsive interactions (from the overall charge on the rods) and net attractive interactions (from the combined affect of van der Waals and dipole interactions). At low concentrations, negligible or no nuclei will form, while, at medium concentration ranges, long-range dipole interactions will account for the arrangement of the nanorods in a side-by-side or end-toend manner to some extent. At high optimum concentration ranges, as stable nuclei will form, the nanorods will undergo preferential attachment/detachment and will be further adsorbed as a perpendicular closed-packed superlattice at the interface and decrease the overall interfacial surface tension. This observation suggests that the nanorod assembly can be governed by the concentration, while keeping the evaporation rate constant, which in turn determines the nuclei size and also controls the clustering of the nanorods. 6.4.3. Assembly in Solution. An interface is not necessary for assembly and particles can coaelesce in solution, due to interactions such as van der Waals, dipole−dipole, or electrostatic along with ligand−ligand, ligand−solvent present in NC solutions, which govern the assembly. Even though the technique cannot be completely controlled, as this relies on the strength of interactions, adequate control can be achieved with careful choice of solvent and capping ligands. The different

interactions lead to the formation of small bundles of NCs in solution, and these act as the nucleation sites for the incoming NCs under Brownian motion, that grow larger and more ordered along with a decrease in potential energy of the system.631,1064,1090,1121,1122 The addition of anti-solvent and control of temperature provide additional factors for manipulation of this process. By addition of anti-solvent into the NC dispersion, the solution can be destabilized. This destabilization leads to a significant increase in the attraction potential between NCs and results in their assembly. Murray and co-workers have shown that by using a combination of low boiling point organic solvent and high boiling point alcohol, controlled aggregation can be achieved.636,1060With an increase in the dispersion concentration and controlled transition of NCs between the solvents, equilibrium for the incoming NC (to find a site to attach on) can be achieved and stable superlattices can form in the solution. In the case of multinary Cu chalcogenide NCs, Ramasamy et al. showed that wurtzite phase Cu2CdxZn1−xSnS4 nanorods can be assembled into ordered structures by changing the dispersibility of the nanorod solution using an antisolvent (ethanol, Figure 39a), which will trigger the aggregation of nanorods in solution.1108 As the assembly is happening in very short time scales (few minutes), the size of these clusters is somewhat limited to a few hundred nanometers, as can be seen in Figure 39b−c. Sun et al. showed that a similar approach can be used to obtain superparticles of CIS/ZnS core/shell NCs by introducing the antisolvent slowly (< hour) using a peristaltic pump.1123 Singh et al. showed that if the assembly process is slowed down to occur over several hours, 3D assemblies of CIGS nanorods can be achieved from a highly concentrated nanorod dispersion, without the addition of any antisolvent.975 In the resultant superstructure, the nanorods are not just aligned perfectly side-by-side and end-to-end (Figure 39d−e) but also the sizes of the superstructures are similar, which strongly indicates that the assembly process followed the typical crystal growth phenomena. In this process, nucleation happens 5957


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Figure 39. (a) Schematic diagram for the NC assembly in solution. (b, c) SEM images show the nanorod assembly achieved by adding nonsolvent in the nanorod solution. Panels (b,c) were adapted with permission from ref 1108. Copyright 2013 Royal Society of Chemistry. (d, e) STEM and SEM images show the 3D assemblies of CIGS nanorods. Panels (d,e) were reprinted with permission from ref 975. Copyright 2012 American Chemical Society.

6.4.4. Directed Assembly Process (Electrophoretic Deposition). NC assembly methods, ranging from simple evaporation processes (solvent) to solution (in vial)-based methods, provide an avenue to glean important insights into the assembly process, with respect to their physical properties (surface charge and dipole moment), but they are ultimately limited in scalability. This limitation can be overcome by utilizing external applied forces (electric and magnetic field) to facilitate NC assembly. Applying an external field to direct the NC is remarkably versatile, not only to control assembly size but also to precisely tune the location where an assembly forms.1125−1128 Initial studies on the directed assembly of NCs (spherical) in the presence of an electric field by Herman et al. provided seminal guidelines about the process and the tunable factors (charge, solvent, drying etc.).1129 Over the years, further research has been done to extend control, beyond just the number of deposited layers and close packed order, but also to the localization of spherical NC assemblies to discrete lithographic trenches on a patterned substrate.1125 In the case of arranging anisotropic NCs, which requires additional control in both translation and orientational order, the application of external fields can provide the necessary control. In the earlier reports of nanorod assembly in the presence of an electric field, a dispersion of nanorods from toluene was drop-cast on a gold-coated Cu electrode and the solvent was allowed to evaporate in the presence of an applied electric field.1130 During this process, the nanorods experience torque that rotates and aligns them in the direction of the applied electric field (i.e. perpendicular to the faces of the electrodes) as the solvent evaporates. This method can yield vertical nanorod assemblies, with micrometer domain sizes or laterally aligned nanorods, when lithographically fabricated interdigitated electrodes are used, which have micrometer spacings between the electrodes.1131 Even though the external applied field approach gives assemblies of micron size order, the scalability of the process remains limited. An extension that

at appropriate concentrations and the nanorods arrange at the growing supercrystal to reduce the potential energy and maximize the packing efficiency. The balance between the attractive forces and the repulsive interactions between the NCs governs the change in free energy of the system, as speculated by the nucleation and crystal growth theory. The formation of a supercrystal is further supported by other entropic forces/ interactions operating between the NCs, where the excluded solvent volume between the NCs maximizes the entropy of the system and drives the assembly. These superstructures are not robust and can be easily disrupted by a mild ligand exchange treatment with amine, which ultimately modulates the charge on the nanorod surface and increases the electrostatic repulsion between nanorods. Manna and co-workers have shown that by simply adding various additives to a stable colloidal solution of complex nanostructures (such as octopods), higher order assembly can be induced by depletion forces.1124 These additives are generally long-alkyl-chain fatty acids and amines (oleic acid, hexadecylamine) or polymers (polystyrene, poly(ethylene glycol) methacrylate), which do not interact with the NC solution. The important parameter is the selection of additives, in such a way that the size of the additives should be larger than the separation between the nanostructure to induce the depletion interactions. When the additive is excluded from the space between the NCs, a positive osmotic pressure (attractive) occurs due to change in the local concentration gradient, which induces the nucleation of NC assembly in solution, followed by the growth of the superstructure. The occurrence of nucleation is dependent upon on the strength of the depletion interaction, which ultimately depends on the size and concentration of additive, respectively. The most important output of assembly in solution, compared to other approaches, is that the assembled superstructures recover minor structural differences, such as different inert-NC spacing, due to double-layer/electrostatic effects. 5958


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Figure 40. (a) Schematic representations of experimental setups used in the assembly of anisotropic NCs, under the influence of electric field. (b) SEM cross-section image showing a multilayer vertically aligned CIGS nanorod assembly over a large area, achieved by using electrophoretic deposition. (c) SEM cross-section images showing the randomly deposited CIGS nanorods. (d, e) SEM cross-section images showing the CZTS nanorod deposited using electrophoretic deposition and the subsequent thin film (with micron sized grains) formed after annealing. Panel b was reprinted with permission from ref 975. Copyright 2012 American Chemical Society. Panels d and e were reprinted with permission from ref 246. Copyright 2014 Macmillan Publishers Ltd: Nature Communications.

inherent property of the nanorods and can be altered only during the synthesis process (i.e. by changing the aspect ratio). The net charge on the nanorod can be tuned postsynthetically through ligand exchange procedures (with conductive ligands) or by performing excess washing with antisolvent to remove any free and loosely bound ligands (creates more dangling bonds). Any increase in the overall net charge on the nanorod surface will affect the mobility of the nanorods (fast for higher charge) under the applied electric field. This can inhibit the formation of close-packed assemblies (superlattice or superstructure), as there will not be sufficient time for nanorods to align in the direction of the electric field. This has been observed in the deposition of CIGS nanorods that have high net surface charge (35 ± 4 mV; after ligand exchange), which resulted in randomly distributed nanorods with their c-axis parallel to the plane of the substrate, as shown in Figure 40c. One highlight of the potential of ordered nanorod arrays is gleaned from work by Mainz et. al.246 In this work, the authors demonstrated how an ordered vertical array of wurtzite phase CZTS nanorods (Figure 40d), deposited using electrophoretic deposition, can be used to form a polycrystalline CZTS thin film with micrometer grain sizes (Figure 40e) by postannealing at much lower temperatures (400 °C) and in much shorter time (few seconds) than traditional processes. These are both attributed to the utilization of a metastable (wurtzite) to stable (kesterite) phase transition that drives grain growth, which is further amplified by the close proximity of nanorods in these superstructured films. In addition, the close structural relation between the kesterite and wurtzite phases with a very fast phase transition (∼9 s) enables the growth of a single-phase, thin film with minimum ion diffusion.

builds on this effort with device level scalability utilizes electrophoretic deposition.1132 In this approach, a direct current is applied between two copper electrodes having ∼2−3 cm spacing within a nanorod solution. A conductive substrate is clamped to one of the electrodes depending upon the net charge on the nanorods. The physical properties of the nanorods (dipole and charge) still dictate the resulting degree of the final superstructure. Figure 40a schematically shows examples of experimental setups, which have been used for the alignment of colloidal nanorods. In particular, two electrodes of exactly the same dimensions are placed 1 mm apart in a vial and a potential difference of 200−300 V (for 1−15 min) is applied, as shown in Figure 40a. The dimensions of electrodes should be exactly equal to prevent dielectrophoresis. The substrate for assembly can be attached to one of the electrodes, if nonconductive, or it can be used directly as the electrode, if suitably conductive. In this case, the net charge and aspect ratio of the nanorods determine the mobilities under the applied field, thus directly affecting the deposition time. The dipole moment of the nanorods assists reorientation of the rods parallel to the electric field, while the net charge of nanorods forces them to migrate toward the electrode and deposit on the substrate. As the assembly of nanorods is field driven in electrophoresis, the process eliminates the requirement of evaporation or external additives. To prevent the formation of cracks and obtain closepacked assemblies, the nanorods should be cleaned multiple times before deposition, as excess ligand coverage and the presence of surfactant will affect the uniformity of the deposited layers. This process has been used for a variety of NCs with different composition, surface chemistry, and aspect ratio, and the assembly is remarkably consistent. Ryan and co-workers have explored this assembly process for nanorods with different compositions from simple binary (CdSe, CdSe) to ternary (CdSeS) to the quaternary Cu chalcogenides (CIGS and CZTS).114,246,975,1064,1127 Figure 40b shows the perpendicular assembly of CIGS nanorods achieved using electrophoretic deposition, where nanorods are not only aligned complete in close packed (side to side) order, but sequential layers of nanorods can also be deposited serially with no voids (free space) in the deposited area.975 The dipole moment is an

7. PHOTOVOLTAIC APPLICATIONS Owing to their outstanding optoelectronic properties, Cu chalcogenides are excellent candidate materials as light absorbers in the field of photovoltaics (PV). As absorbers, PV devices require semiconductor materials with a direct band gap in the range of 1.0−1.5 eV, and a high absorption coefficient. These properties are matched by several Cu 5959


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processing techniques. This is a completely unique case, since electronic and optoelectronic devices commonly show superior performances when produced using vacuum-based technologies, in addition to high-purity precursors and very clean environments. The reason behind the exceptionality of CZTS solar cells is the possibility to produce and optimize homogeneous layers of complex materials in an easier and faster way using solution processing techniques, when compared with vacuum-based technologies demanding additional efforts and time to screen a number of compositions/ parameters/conditions. To produce high-efficiency CZTSSe solar cells, an additional disadvantage of vacuum-based technologies over solution-processing ones is the volatility of chalcogens and tin at the relatively large annealing temperatures used in a vacuum atmosphere. This volatility makes the control of the layer composition, a key parameter to produce high efficiency devices, extremely difficult.167 This anomaly was not observed during the development of CIGSSe solar cells probably because, at the time, solution-processing technologies were not as well established as in the present, which allowed single-crystal growth techniques and evaporation technologies to provide the best efficiencies from the very beginning. Besides the current eventual record efficiency for CZTS, inkbased solution processing technologies offer obvious advantages over vacuum-based technologies. The main advantage is a reduction of the energy payback time, which is mainly associated with the purification and processing costs of the elements and absorber layers in vacuum processed solar cells. Solution processing technologies have lower production costs, related to the need for lower precursor purities, lower equipment investment/depreciation, potential continuous production of large area devices in roll-to-roll processes, lower energy usage, high composition homogeneity, safer and more environmentally friendly processes, and higher material yield and productivity. Furthermore, the use of NC-based inks allows for simple adjustment of the composition to the atomic scale and enables high-throughput production of large-area devices with homogeneous composition, and with a potentially outsourced material production. The high-potential economic, social, and environmental impact of PVs and, particularly, of low-cost solution-processing PVs has motivated a very large interest in the synthesis of potential absorber materials (such as CIGSSe and CZTSSe) in NC form, the formulation of inks based on these NCs, the development of large-scale NC deposition technologies, and, in some cases, the annealing and characterization of the optoelectronic properties and testing of the absorber layers and PV devices produced from these NCs. Numerous works have also been devoted to the detailed study of the specific problem associated with this technology, where the main challenge is the maximization of the charge carrier life time before recombination and the enhancement of the charge transport within the absorber layer. The availability of suitable semiconductor materials in the form of NC-based inks offers obvious advantages to process solar cells. However, since efficient devices require charge to be extracted from relatively thick layers (when compared with size of NCs), the small size of NCs, the related high density of interfaces and surface defects, and the presence of organic ligands on the NC surface become important constraints. To overcome these limitations, the easiest approach is to sinter the NC-based layers to produce a highly crystalline absorber with

chalcogenides, such as the binary Cu2−xS, ternary CIS, and quaternary CZTS, among others.396 Currently, silicon technology dominates the PV market. However, silicon is an indirect band-gap semiconductor with a low light absorption coefficient. Thus, silicon solar cells require thick absorber layers (∼100 μm) that impose long diffusion lengths for the charge carriers to reach current collectors. To minimize losses of photogenerated charge carriers, crystalline silicon with a very high purity is necessary. The production of such high-purity silicon demands a high energy consumption and, thus, this results in elevated costs. More than half of the relatively large energy payback time of single-crystal Si solar cells (ca. 2 years) is associated with Si purification and processing.1133,1134 To overcome the intrinsic limitation of Si, a second generation of solar cells was developed on the basis of direct band-gap semiconductors, characterized by higher absorption coefficients. The improved light absorption in these materials allowed for a significant reduction of the absorption layer thickness (∼1 μm) and, consequently, for larger densities of defects, including grain boundaries and impurities. The first thin film solar cells, dating from 1954, were based on CuxS, particularly on a CuxS/CdS junction, and displayed 6% power conversion efficiencies (PCEs).1135,1136 However, Cu2S solar cells are limited by a low stability, which is associated with the degradation of Cu2S to Cu2−xS and to the diffusion of Cu to CdS.38,1136,1137 Cu2−xTe was also used in combination with CdTe to produce heterojunction solar cells.1138 Currently, CdTe is a mature technology, but it has the limitation of using toxic and scarce elements. In the last decade, some ternary and quaternary Cu chalcogenides have become key absorber materials in the field of thin film photovoltaics. In particular, Cu(In,Ga)(S,Se)2 (CIGSSe) is a well-established absorber material, which is at the basis of the fastest growing commercial thin film PV technology and is expected to dominate the thin film PV market over the next number of years. CIGSSe has an optimum direct band gap in the range of 1.04−1.50 eV, depending on composition, a very high absorption coefficient (3−6 × 105 cm−1), excellent stability under operation conditions,1136,1139 and appropriate electron affinity and lattice constants to allow suitable heterojunction formation with wurtzite CdS or ZnCdS. Owing to these advantages, PCEs up to 12% were already obtained in 1974 using a CISe/CdS heterojunction architecture, which was made by vacuum evaporation of CdS onto single-crystal, p-type CISe.1136,1140 Currently, CIGSe holds laboratory record efficiencies up to 22.3% obtained using vacuum-based technologies. However, the high cost of indium within CIGSe solar cells has motivated the search for lower cost PV materials. New alternative materials can be found by moving toward more complex compositions, which have similar or improved functional properties but employ more abundant, lower cost, and also less toxic elements. Aside from the boom of perovskite solar cells, which have already reached efficiencies above 20% but still present important stability issues and a limiting dependence on toxic Pb, another Cu chalcogenide, Cu2ZnSn(S,Se)4 (CZTSSe), has appeared as a main candidate absorber material for future solar cells. CZTSSe solar cells have already reached power conversion efficiencies (PCEs) up to 12.6% using a similar architecture as CIGS solar cells.245,1141 While they are still under development, a particularity of this absorber material is that record PCEs have been obtained, quite unexpectedly, from absorber layers produced using solution 5960


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improved charge transport properties. This is the approach followed to produce NC-based, solution-processed sintered solar cells using either a planar configuration or a nanocrystalline scaffold, as reviewed in sections 7.1 and 7.2, respectively. Alternatively, nanocrystalline absorber layers can provide high PCEs at lower processing costs than sintered layers if proper ligand exchange/displacement treatments are applied to facilitate charge transfer between NCs and to passivate surface traps. In this review, we will refer to the PV devices where charge transport takes place through the NC absorber layer, as NC solar cells (NCSC), which are reviewed in section 7.3. A third option to maximize charge carrier collection is to attach/combine/blend surface-treated NCs with other materials that have optimized transport properties and are able to collect one or both types of charge carriers. This is the strategy behind the so-called semiconductor sensitized solar cells (SSSCs) and hybrid organic-inorganic solar cells, reviewed in sections 7.4 and 7.6, respectively. Cu chalcogenide NCs can also be used as absorber materials on third generation solar cell concepts such as multiple exciton generation (MEG) solar cells and intermediate band solar cells, which are reviewed in section 7.7. Furthermore, beyond their use as absorber materials creating electron−hole pairs, Cu chalcogenides have also been used as counter electrodes in SSSCs (section 7.5) or as photoluminescent materials for luminescent solar concentrators, luminescent down-converters, or plasmonic enhancers of light absorption (section 7.8). In the following subsections, we will discuss the use of Cu chalcogenide NCs in all these PV-related applications (as illustrated in Figure 41). Further

details on the use of NCs for solar energy conversion through PV can be found in a significant number of reviews and books focused on this topic.1142−1147,20,167,1136,1148−1152,23 7.1. Thin Film Solar Cells Based on Sintered NCs

The solution processing of Cu-based sintered thin film solar cells from NC-based inks/baths involves (Figure 42) : (i) the NC preparation; (ii) in some cases, a ligand displacement/ exchange step; (iii) the formulation of an ink/bath; (iv) the deposition of the NC ink by techniques such as spin coating, dip coating, doctor blading, spray deposition, or electrophoretic deposition, generally over molybdenum coated (∼800 nm) soda-lime glass substrates; (v) the thermal treatment of the films to sinter the NCs into a highly crystalline absorber, typically in a Se atmosphere at temperatures around 500 °C for 20 min; (vi) a selective etching process with KCN generally applied to remove undesired binary CuSe phases, which would lead to shunted cell behavior; and finally, (vii) the completion of the PV device by the chemical bath deposition (CBD) of a thin (∼50 nm) buffer layer of CdS, radiofrequency (RF) sputtering of a thin (∼50 nm) i-ZnO film, and RF sputtering of a transparent conductive oxide (TCO) top layer (∼200 nm), typically Al-doped ZnO (AZO) or indium tin oxide (ITO), the thermal evaporation of a patterned electrode array, commonly a Ni/Al grid as a top contact, and, in some cases, the addition of a MgF2 antireflection coating (∼100 nm) deposited on top of the device by electron beam evaporation. Typically, laboratory prototype samples are mechanically scribed into individual solar cells with a size of around 0.5 cm2.337,1035,1156 The main advantage of PV devices produced from absorber layers obtained from NCs, as opposed to absorber layers obtained from conventional vacuum-based technologies, is cost reduction. In this scenario, NC-based strategies compete with other low-cost, solution-processing strategies, such as those based on inks obtained from molecular or ionic precursors or electrodeposition. When compared with competing solution processing strategies, the a priori main advantage of NCbased processes is the potential saving of the selenization/ sulfurization step, as the chalcogen can be already properly incorporated and in the correct amount within the NC. However, the annealing step in a chalcogen containing atmosphere is rarely saved. A second advantage of NC-based solution processing methods is the a priori higher purity and compositional homogeneity of the produced absorber layers, which should translate into devices with higher efficiency. However, this point remains to be demonstrated. Current state-of-the-art solution-processed devices have been obtained through the use of hydrazine and printingbased technologies. Hydrazine-based processes hold the absolute efficiency record for CZTSSe solar cells at 12.6% and the NC ink-based record for CIGSSe solar cells at 15.2%.245,1157 Hydrazine inks are prepared by dissolving powders of the corresponding binary chalcogenides in hydrazine, with an excess of chalcogen to readily form metal chalcogenide complexes coexisting with hydrazinium species. In the case of CZTSSe, zinc compounds such as ZnS and ZnSe show negligible solubility in hydrazine. Thus, zinc powder is added to in situ form ZnSe nanoparticles1158−1160,1156 or is incorporated in a soluble form with the help of hydrazinocarboxylic acid, for instance.1161 However, hydrazine is highly toxic and is a very unstable compound that requires extreme caution during handling and storage; thus, its industrialization has important limitations.

Figure 41. Cu chalcogenide NCs have been used as absorbers in different types of solar cells (SCs). In some architectures, NCs supported on planar substrates or on porous frameworks are sintered to form highly crystalline films. In others, they are used as light absorbers in NC form. Cu chalcogenide NCs can also play other roles in the PV field, including the enhancement of the light absorption in the SC and the electrocatalytic regeneration of the electrolyte at the counter electrode of semiconductor liquid−junction solar cells. The counter electrode panel is reproduced with permission from ref 1153. Copyright 2015 Royal Society of Chemistry. The optical enhancement panel is reproduced with permission from ref 1154. Copyright 2015 Macmillan Publishers Ltd: Nature Nanotechnology. The 3 rd generation panel is reproduced with permission from ref 1155. Copyright 2013 Macmillan Publishers Ltd: Scientific Reports. 5961


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Figure 42. Scheme of the typical process to produce sintered thin film solar cells from NCs.

selenized layers) present the best PCEs. Record CZTSSe solar cells are obtained with a mixture of S and Se, at a composition [S]/([S]+[Se]) = 0.23 for sputtered absorbers and 0.09 for hydrazine-based inks.167 This optimum chalcogen composition could be partly related to an enhancement of the nonradiative lifetime of minority carriers, due to the vanishing of the tetragonal distortion around these compositions. A similar argument was used to justify the effect of Ga in a concentration [Ga]/([Ga]+[In]) = 0.2 in CIGS solar cells.1175,167 Following a similar strategy, that the partial replacement of In by Ga helps to improve CIGSe solar cell performance, the partial replacement of Sn by Ge has recently been demonstrated to improve the Voc and the overall PCE in CZTS solar cells.369,1034,1035,337,1036,239 The incorporation of Ge increases the minority charge carrier lifetimes by reducing the deep level traps associated with multivalent Sn.1035,1036,1039 It also allows the band gap to be adjusted, which can be used to produce band-gap-graded solar cells and multijunction cells.1176 Undoubtedly, the realization that composition plays a key role on PV performance has had a strong impact in the field of NC synthesis, triggering the development of mechanisms and routes to control the composition of ternary, quaternary, and even quinary Cu chalcogenide NCs for PV applications. Multinary NC-based inks require NCs that overall are Cu-poor and have appropriate amounts of Ga and In or Ge and Zn,337 which may be an important challenge. Luckily, to produce sintered crystalline films, exceptionally narrow size distributions are not required, although wide size distributions may be indicative of large compositional inhomogenieties.1177 However, if no phase segregation takes place, excellent composition distributions within the NC ensemble are neither required and the composition distribution within each NC1178 may not even be a significant parameter. Instead of excellent particle and composition distributions, the technological application of NCs in large-area sintered devices (such as solar cells) requires lowcost, environmentally friendly safe, and large-scale synthesis protocols that are capable of producing highly concentrated NCs using minimum amounts of (recyclable) solvent, ligands, and additives and at a very high yield. In addition, ligands that are easily decomposable would help in reducing the number of technological steps.1012 At the same time, automated NC purification processes will be required. If NC-based technologies are to be relevant in such a high volume for PV applications, a complete redefinition of the synthesis procedures currently used to produce multinary NCs is probably a prerequisite. Luckily, a relatively broad range of processing parameters seems to provide adequate NCs for sintered thin film solar cells. For example, Ford et al. tested different CISSe NC precursors, such as acetates, acetylacetonates, iodides, chlorides, and nitrates, and concluded that the synthesis conditions can be optimized for each precursor to yield NCs with similar properties.1179 The NCs were consequently used to form absorber films, which had roughly the same charge carrier concentrations of 1016-1017 cm−3, and devices with similar

Strategies to replace hydrazine by compounds such as butyldithiocarbamic acid in a low-toxicity ethanol-based molecular ink have been developd.1162 However, devices produced by these methods did not reach the high efficiencies attained with hydrazine. Beyond hydrazine-based processes, most works use either salts, usually in combination with thiourea in solvents such as DMSO, alcohols, or water,1163,1164 or NCs in a variety of solvents. In the latter direction, which is the only one considered in this review, sintered solar cell devices based on Cu chalcogenide NCs have reached efficiencies up to 9.8% for CZTSSe1035,983 and 15% for CIGSSe.1165 7.1.1. Deposition Technologies. The most common technologies to produce laboratory-scale NC-based sintered thin film solar cells are spin coating, doctor blading, spray deposition, dip coating, and electrophoretic deposition. Spin coating makes use of low-viscosity inks and generally requires multiple deposition steps, with a low temperature (∼150 °C) thermal process in between to ensure fixation of the layer before the next one is applied. This low-temperature thermal process is typically applied using a hot plate in an air atmosphere. Doctor blading, also known as knife coating or tape casting, involves applying a relatively high-viscosity ink on a surface limited by two barriers (to control the film thickness), where the barriers are often created by means of a tape. The material is applied by moving the blade over the barriers, relative to the substrate. Spray deposition uses relatively lowviscosity inks and substrates that are typically heated at a temperature close to the boiling point of the solvent.1166,1167 Generally, spray pulses are used to keep the substrate temperature constant and to prevent the formation of drops over the substrate. Dip coating1168 and electrophoretic deposition1169,1170,1169 are two bath-based techniques which involve the use of NCs suspended in a liquid, whereby the NCs are deposited on the substrate immersed in the bath with or without the aid of an electrostatic field. In fact, electrophoretic deposition of CdS NCs was already used to produce CdS/Cu2S solar cells back in 1979.1171 7.1.2. Ink Composition. Control of the composition of the absorber layer is a key parameter to maximize the device performance. This is especially true and particularly challenging in multinary CIGSSe and CZTSSe materials, where the presence of highly conductive secondary phases such as Cu2−xSe needs to be avoided.337,1172,1156 Futhermore, the stoichiometry of the final material needs to be carefully adjusted to tune the band gap, to optimize the amount of electronically active defects, and to minimize recombination at the CdS−absorber interface.167,1173 In this regard, both CIGSSe and CZTSSe solar cells are generally Cu poor: [Cu]/([In]+[Ga]) = [Cu]/([Zn]+[Sn]) = 0.8. While the optimum Ga composition in CIGSSe is around [Ga]/([In]+[Ga]) = 0.2, CZTSSe is required to be zinc rich: [Zn]/[Sn] = 1.2.167,1173,1174 With respect to the chalcogen, sulfides have associated lower processing costs and toxicity, but selenides (or partially 5962


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in the appropriate ratios in the precursor solution, most works on the use of NC-based inks to produce sintered thin film solar cells use sulfide NCs and introduce an annealing step in a Se-containing atmosphere. This leads to the partial or full selenization of the precursor layer by chalcogen exchange. This chalcogen replacement has an associated significant volume expansion (14%), which helps film densification and is beneficial to minimize voids and cracks in the film.969,1181 Such reactive sintering is also considered to favor the formation of larger crystal domains, when compared with the annealing of sulfide or selenide NCs in an inert atmosphere or in an atmosphere containing the same chalcogen as in the NC.1191 The use of metal NC-based precursors has an even larger associated volume increase with sulfurization or selenization.1183 However, if the NC surface is not protected, with citrate for instance,1183 metal NCs generally contain an oxide layer, which demands an additional reduction step and/or the use of severe sulfurization/selenization treatments, with highly toxic H2Se or H2S gases, to be converted to the proper composition. Generally, the Se-containing atmosphere is created by introducing a small amount of elemental Se pellets into a closed ampule or graphite box containing the substrate, in a furnace. A better control of the atmosphere composition is obtained in 2- and 3-zone furnaces, which allow for independent control of the Se vapor pressure and the film annealing temperature. High-vacuum evaporation chambers, with precisely controlled vapor flows, have also been used in this direction. However, vacuum conditions may increase the fabrication costs, and low sintering pressures may result in the evaporation of other elements, thus leading to difficulty in controlling the composition of the final layer. The use of gasphase selenization using toxic gases (such as H2Se) also hinders low-cost scaled up production, which is one of the presumed advantages of solution-based processes. While the use of sulfide absorber layers would have an advantage in terms of material cost and processing, their efficiencies have always fallen below those of selenides. To avoid this limitation, inks containing excess amounts of Se have been also formulated and used to produce CISe and CZTSe absorbers.1181 Excess selenium can be added to CIGS precursor inks either as elemental Se particles or as Se2−-containing salts. In this direction, Maes et al. recently demonstrated that (N2H5)2Se can stabilize CIS NC inks and induce the formation of a pure CISe phase.1181 The drawback of introducing an excess of Se in the film is the difficult composition gradation by means of the chalcogen composition, which is a common strategy to produce high-efficiency CIGSSe and CZTSSe devices. Besides the potential segregation of phases in highly offstoichiometric ink compositions and the incorporation/conservation of the chalcogen(s) with the appropriate amount(s) and distribution, the main challenge to tune the final film composition, especially in CZTS, is to control the elemental losses that occur during high-temperature annealing. Specifically, Sn can be lost in the form of SnS(g) and/or SnSe(g) at annealing temperatures ≥500 °C. To limit the degree of elemental losses, annealing temperatures are typically kept below 550 °C and the Se partial pressure is maximized by incorporating an excess of this element in a graphite box or sealed ampule.1035 Some experimental works also consider the introduction of small amounts of Sn or SnSe to compensate for the loss of this element during the annealing step.

efficiencies were achieved. However, other reports have demonstrated that the selected elemental precursors can have an influence on the composition of the final absorber film. For example, Hages et al. found that the final Ge loss during selenization of CZTGeSSe absorbers, produced from Cu2Zn(SnyGe1−y)S4 NCs, seemed to depend on the Ge precursor used in the synthesis, either GeI4 or GeCl4.1035 Most works use hydrophobic NCs and nonpolar, organic solvents such as toluene or dichlorobenzene (DCB) to fabricate solar cells. However, more economical and environmental friendly processes have also been investigated, where hydrophilic CZTS NCs that are coated with polyvinylpyrrolidone (PVP) and/or ethylene glycol can be dispersed in polar solvents such as ethanol or water, with the consequent solar cell devices reaching PCEs close to 2%.1002 In the same direction, relatively large-scale sonochemical processes (up to 200 g per batch) for surfactant-free CIGSSe and CIS NCs have also been developed and have reached PCEs up to 2.62%.979,1180 NC-based inks have not only been formulated using ternary or quaternary Cu chalcogenide NCs, but also by using metal or oxide NCs, combinations of binary and ternary Cu chalcogenides, and combinations of binary/ternary Cu chalcogenides with salts or molecular precursors.1181−1185 In spite of the presumed advantage in composition control of NC-based inks, inks based on combinations of binary/ternary NCs (or combinations of NCs with salts) are stated to provide easier control of the final absorber composition.1186−1188 In this regard, Cao et al. used a mixture of Cu2SnS3, CuxS, and ZnS NCs to produce 8.5% CZTSSe solar cells,1186 Cai et al. used mixtures of metal chloride salts (i.e. CuCl, InCl3) with CuS and In2S3 NCs, coupled with the addition of thiourea, to produce CIS absorbers,1189 and Chesman et al. controlled the Ge composition in CZTGeSSe absorbers by mixing CZGeS and CZTS.364 Apart from tuning the NC composition, or using combinations of NCs with salts, another strategy to control composition is by introducing ligands that contain one or more of the constituent elements of the final layer. This was demonstrated by Jiang et al, where NCs were used in combination with molecular metal chalcogenides to produce CISe, CIGSe, and CZTSSe absorbers.981 Band-gap graded CZTGeSSe solar cells were also produced from CZGeS NCs that contained controlled amounts of Sn2S64− MCC ligands bound to the surface.1190 7.1.3. Sintering Atmosphere. An annealing step is used to sinter the NCs into large crystal domains that facilitate charge transport, reduce surface recombination defects, and overall, increase photogenerated charge carrier lifetimes. While the use of NCs with controlled composition is presumed to be advantageous in adjusting the final film composition and obtaining highly homogeneous films, the annealing step can actually strongly affect these parameters. Therefore, careful adjustment of the annealing time, temperature, and atmosphere is necessary to maintain/reach phase purity and proper defect control in CIGSSe and CZTSSe films. Besides the growth/ conservation of the desired phase, by removing or avoiding the formation of secondary phases, the main compositional effects of the annealing treatment are the incorporation/exchange of the chalcogen and the removal of high-vapor-pressure elements, such as Sn and Ge, which need to be prevented or anticipated and compensated. While an a priori key advantage of NC-based solution processing technologies is the incorporation of all the elements 5963


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The incorporation of Ge in CZTGeSSe poses another problem in this direction, as Ge and GeS2 are volatile compounds with a greater volatility than SnS.1034,1036,1035 Hages et al. reported efficiencies of up to 9.4% for CZTGeSSe solar cells, which were produced using quinary CZTGeS NCs with a Ge content of 30 at. %. This high efficiency was reached through minimization of bulk Ge loss during the high-temperature selenization step of the absorber processing.1035 A drawback of using high Se partial pressures has been the formation of a Se-rich “fine-grain” layer between the sintered CZTSSe layer and Mo back-contact of the film. In situ, energydispersive XRD and fluorescence analysis studies were recently conducted to investigate the selenization process of CZTS NC thin films and to provide important insights into this process.1192,1193 The studies revealed that selenization takes place through two simultaneous reactions: a direct and fast formation of large grains of selenides, starting with copper selenide grains on top of the NC layer as soon as Se is incorporated into the film, which are subsequently transformed into CZTSe through the outdiffusion of the cations to the surface of the NC layer. This is accompanied by a slower replacement of S by Se in the NC layer to form a bilayer structure, where the top layer is fully sintered into micrometer grains and the bottom layer consists of unsintered nanometer size kesterite grains. The thickness of this unsintered bottom layer has been found to depend on the selenium partial pressure, the selenization time, and temperature profiles.1034,337,1035 High Se partial pressures also result in a catastrophic complete selenization of the Mo layer, which can be prevented by using thicker Mo films or appropriate selenium diffusion blocking layers.1194 There has been important progress in the direction of optimizing layer crystallinity using milder annealing conditions, which in turn reduces the energy and time costs. A particularly interesting strategy is the use of metastable NCs (Figure 43) to promote crystallization at lower temperature,246,1195,1196 without negatively affecting the device performance.556 Alternatively, the use of amorphous nanoparticles was also demonstrated to provide easier sintering, while, at the same time, allowing the nanoparticle synthesis temperature to be reduced.1197 To promote layer crystallization, extrinsic additives are generally incorporated into the absorber film before or during crystallization. In this direction, the incorporation of small amounts of alkali metals and, particularly, sodium1198,1199 has become a widely accepted requirement to reach high device performances. Besides promoting crystallization, sodium strongly increases the hole concentration of the absorber layer and improves the minority carrier lifetime by defect passivation, resulting in enhanced fill factors, larger open circuit voltages (Voc), and overall higher PCEs.1198,1200,1201 In the same direction, potassium has been demonstrated to be an interesting alternative to sodium, which enabled flexible PV devices with remarkable PCEs to be fabricated.1202 Taking advantage of its small size, sodium is typically incorporated into the absorber layer during the annealing step, by diffusion from the soda-lime glass (SLG) and through the Mo back contact. However, this method does not allow facile control of the amount of sodium that diffuses from the substrate. At the same time, this strategy is not valid with flexible substrates. In vacuum-based CIGS and CZTS solar cells, sodium is alternatively incorporated using additional vacuum-deposition steps to form a thin layer of a sodium compound such as NaF or Na2Se over the Mo (before CIGS

Figure 43. Cross-sectional SEM images of nanorod films annealed at three different temperatures. (a) The film annealed to 355 °C still shows nanorods of original size (scale bars, 200 nm). (b) The film annealed to 375 °C already shows some larger grains, but also still some small crystallites (marked by white arrow) near the size of the original nanorods (scale bar, 200 nm). (c) Finally, the film annealed to 400 °C shows only large grains (scale bar, 200 nm). Reproduced with permission from ref 246. Copyright 2014 Macmillan Publishers Ltd: Nature Communications.

deposition) or over the deposited absorber. Alternatively, sodium is incorporated by more complex deposition processes involving the codeposition of sodium during the absorber growth.1200,1203 While numerous works based on solution processing also make use of additional steps, such as soaking the film in a sodium salt solution (e.g. refs 1204 and 1012), solution-based technologies allow for facile incorporation of controlled amounts of Na in a unique manner by just adding the sodium-containing compound into the ink. Sodium compounds can also be incorporated directly at the NC surface to enhance their effect. This was demonstrated by Zhou et al., who introduced sodium at the CZTS NC surface by means of CF3COONa in oleic acid, in which a significant enhancement of the PCE was obtained, and this was related to the increased carrier concentration and elongated minority carrier lifetime, induced by defect passivation.1201 Specifically, a 50% enhancement in the PCE (6%) was achieved from the surface passivation of CZTS NCs with sodium, as compared to the unpassivated CZTS NC device (4%) in this work. Another strategy to promote grain growth is the use of Cu-rich phases such as CuS or CuSe, which have lower melting temperatures and thus facilitate material diffusion and crystal coalescence. While these phases can facilitate grain growth, they could also ruin the device by causing electrical shorts; thus, 5964


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Figure 44. Cross-sectional scanning electron microscope images for “champion” CZTSSe and CZTGeSSe (30% Ge) devices. Below the sintered CZTSSe and CZTGeSSe absorbers, the “fine-grain” layer (red) followed by MoSe2 (green) and Mo (blue) layers are indicated. Reproduced with permission from ref 1035. Copyright 2015 John Wiley & Sons, Ltd.

they should not be present in the final absorber layer. In this scenario, the use of Cu-rich NCs within an overall Cu-poor precursor has been demonstrated as an excellent strategy to promote grain growth while achieving a proper absorber composition. In this direction, Cai et al. produced CISSe solar cell from mixtures of salts and NCs in water and showed that the CuS NCs acted as nucleation sites, which accelerated grain growth and allowed grain sizes 7 times larger to be achieved.1189 Moreover, in this work, the use of water as a solvent involved an obvious cost reduction, and the lack of organic solvents also helped to prevent the formation of a carbon layer in the final absorber material. Jeong et al. and Seo et al. independently showed that the intentional incorporation of a CuSe phase, within a multiphase NC-based ink, also helped to crystallize CISe absorber layers.1205,1206 External elements, such as Sb and Bi, have also been used to promote crystal growth by the formation of lower melting point phases.1207,1208 Salts of these elements can be introduced directly within the ink or can be used with the additional role of displacing organic ligands from the NC surface, as shown by Carrete et al.350,1167 In general, the incorporation of extrinsic additives into the NC or the ink may enable a facile, quantitative, and versatile approach to tune the final material functional properties. These additives can be used to control composition, improve crystallinity, passivate defects, or adjust the charge carrier concentration. 7.1.4. Graded Solar Cells. The introduction of a chemical potential gradient within the absorber layer has been demonstrated to be an excellent strategy to direct minority carriers toward the proper electrode, and away from high recombination interfaces, thus significantly improving the device efficiency.1152,1209 Band-gap grading can be introduced by compositional grading within the absorber layer. This can be obtained by changing the Ga composition within CIGSSe solar cells1210−1213 or by introducing Ge gradients in CZTGeSSe solar cells.369,1190,1176,1035 These compositional graded layers can be relatively easily produced by solution-processing strategies using multistep deposition processes such as spin coating, dip coating, and pulsed spray deposition. In addition, the [S]/[Se] ratio can be tuned within the layer in both cases to obtain a chemical potential gradient. This chalcogen compositional gradient can be obtained during selenization.1214 7.1.5. Surface Ligands/Carbon. While colloidal synthesis methods require the presence of organic ligands to control NC growth, to provide colloidal stability, and sometimes also to adjust composition, they generally need to be removed before

the annealing treatment to facilitate crystal growth and to prevent the incorporation of large amounts of carbon in the final absorber layer. Organics hinder mass transfer between adjacent NCs during annealing and, thus, limit grain growth. When NCs with organic ligands or inks containing organic binders are used, a large amount of carbon is found in a layer next to the Mo after annealing in an oxygen-free atmosphere.556 Typically, a bilayer structure (Figure 44) is obtained after selenization, which consists of an upper highly crystalline layer and a Cu-, carbon-, and Se-rich unsintered bottom layer with smaller grains, that limits PCE.1174,1195 This carbon rich interlayer is difficult to remove without highly severe treatments, which may result in an uneven growth of the crystals and the formation of porous structures.1195 The accumulation of carbon beneath the large grained CZTSe layer occurs as a consequence of the diffusion of cations out of the carbon-containing NC layer.1192 The source of carbon in the unsintered layer has been attributed to capping ligands with long hydrocarbon chains, such as OLA or DDT, that are present on the NC surface prior to annealing. After annealing, the long hydrocarbon chains carbonize and leave these carbon residues behind in the film. While the type of ligand may not be seen as a major concern in sintered solar cells, as compared with nanocrystalline solar cells, Martin et al. recently demonstrated an important effect of ligands on the formation of pyrolized graphitic carbon within CZTSSe absorber layers and, thus, on the related PV performance.1215 They compared sintered layers made of NCs capped with OLA and DDA (Figure 45) and found that although OLA is longer than DDA, OLA as a ligand leads to higher performance devices. This was attributed to the formation of larger CZTSSe crystals in the OLA containing layer during annealing. The double bond in OLA led to the formation of a more ordered graphitic carbon, that readily phase separated from the CZTSSe grains, facilitating grain growth and resulting in the commonly observed bilayer structure. Pyrolized graphitic carbon, which is electrically conductive and also reacts with chalcogens to produce heterocyclic moieties, was incorporated into the device absorber layer. However, the DDA-related carbon underwent much less phase separation during annealing and had smaller grains with a more disjointed morphology.1215 In view of the ligands being problematic, several works have also been devoted to the synthesis of organic ligand-free NCs for PV applications. For example, Huang et al. recently described a procedure to produce large grained CZTS thin films, without the usual fine-grained underlayer, from CZTS NCs produced in formamide and thiocetamide and presumably 5965


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Figure 45. Left: Scheme of the ligand effect on the formation of graphitic carbon in sintered CZTS nanoparticle-based solar cells. Right: a and d show the top-down SEM images for the OLA- and DDA-based films, respectively. b and e show the cross-sectional SEM images for the OLA- and DDA-based films, respectively, where the white line represents the EDS line-scan region for each. The black scale bars present in the SEM images are 500 nm. Reproduced with permission from ref 1215. Copyright 2015 American Chemical Society.

capped with S2−.1216 Ligand-free NC-inks have also been obtained from top-down ball milling processes or mechanical alloying (e.g. refs 1217−1219), with PCEs up to 4.8% achieved for CISSe.1219 In parallel to the development of processes to obtain ligandfree NCs or NCs with easily decomposable capping molecules such as triphenylphosphate,1012 a multitude of ligand exchange/ displacement processes have been developed. These processes can be used to further control composition by means of metal chalcogenide complexes (MCCs),981,1190 to introduce crystal growth promoters,350 or to passivate surface defects.1201 Ligand exchange strategies using 5-amino-1-pentanol have also been employed to allow benign polar solvents such as ethanol, n-propanol, or even water to be used.1220 In addition, these polar solvents allow for the controllable incorporation of dopants such as sodium, potassium, and antimony. These ligand exchange processes are carried out with the NCs in solution or after their deposition in a layer. When performed in solution, the colloidal stability of the NCs is compromised, which is a main drawback for deposition techniques such as spin coating, spray deposition, or dip coating that use low-viscosity inks/baths, but it is not dramatic in the higher viscosity inks required for doctor blading, for instance. When ligands are displaced from NCs that are already supported, a limitation exists on the thickness of the layer that can be treated. This is an important drawback for relatively thick film processing technologies such as doctor blading, but aside from the time increase, it is not a major drawback for multistep deposition processes such as spin coating or dip coating. For example, Riha et al. used EDA to displace TOPO capping ligands from dip-coated films by dipping the substrate into a 0.01 M solution of EDA in acetonitrile.1221 7.1.6. Other Architectures/Components. Besides optimization of the absorber layer, a large amount of work has focused on the optimization of other cell components in Cu chalcogenide NC-based sintered thin film solar cells, such as the replacement of CdS for ZnS to fully eliminate Cd from the device1002 and the replacement of heavy glass substrates by light and conformable metal foils.1184,1002 The use of superstrate type architectures in several cases using nanostructured oxide supports may limit the annealing temperature and prevent the use of Se atmospheres; thus, the final layer composition needs to be controlled exclusively in the

ink.1222,1223 This type of architecture is considered within section 7.2. 7.1.7. Other Absorbers. Besides CIGSSe, CZTSSe, and CZTGeSSe, NC-inks and absorber layers of other Cu chalcogenides have also been investigated. For example, Cu2−xS NCs have been used to produce solar cells with 0.24% PCE by Liu et al.1224 The introduction of alternative elements (such as Al, Fe, or Zn) into CGS or CIS has also been a main strategy to increase absorption in the visible and IR part of the solar spectrum and, thus, develop novel PV materials. For example, Brik et al. studied the properties of Al-doped CGS,1225 and Vahidshad et al. studied Fe- and Zn-doped CIS NCs.977 Aside from CIGS and CZTS, ternary Cu chalcogenides containing Sn, Sb, or Bi to form Cu-A-X semiconductors (A = Sn, Sb, Bi; X = Se, Te, S) are probably the most popular absorber materials in Cu-based thin film solar cells.1226, 272,307,1227,783,219,308,291,1228,404,405,1229,1230,953,1231−1233,406, 272,404,1234,448,1235

In particular, the Cu−Sb−S compounds are composed of highly abundant and low-toxicity elements and have direct band-gaps with energies that can be tuned by adjusting the composition, from 1.5 eV for CuSbS2 to 1.0 eV for Cu2SbS3. Using DFT calculations, Yang et al. showed that CuSbS2 has superior defect physics and that it has an extremely low concentration of recombination centers within the forbidden gap, especially under the S rich condition.269 In addition, Yang et al. demonstrated that highly crystalline and phase-pure CuSbS2 films with large grain sizes could be successfully obtained using a hydrazine-based solution-process, and PCEs up to 0.5% could be reached in this report.269 The selenide form, CuSbSe2 was prepared in a separate report, and PCEs up to 1.3% were achieved in this work.448 While the production of thin films of these materials is mainly achieved by salt-based solution processing strategies,1236 NCs and their corresponding NC-based inks have also been developed. For example, Cu12Sb4S13 and Cu3SbS4 NCs were produced and characterized by J. van Embden et al.292 Cu2SnS3 NC-based absorbers have also been produced, which showed an appropriate direct band gap ∼1 eV, a high optical absorption coefficient of >104 cm−1, and PCEs of 0.135%.1237 Cu3BiS3 is another excellent candidate for PV,309,310,1234, 311,1238 which has a direct band gap of about 1.4 eV and a high optical absorption coefficient of >104 cm−1. While Cu3BiS3 NCs have been synthesized and characterized, no PV NC-based 5966


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devices have been reported.577 Several other Cu chalcogenides, such as CuAlS2 and CuAlSe2,1239 have also been proposed for solar cells, but no attempt to fabricate devices of these materials by NC-based solution processing technologies has been reported.

The introduction of an ultrathin insulating oxide coating on a mesoporous TiO2 film, to decrease the recombination rate, simultaneously represented a substantial improvement on CIS ETA solar cells.1246 In terms of processing, the use of metal salts was found to be advantageous to infiltrate CIS inside the pores of nanostructured TiO2 and to reach higher coverages. However, the growth of the CIS absorber by AL-CVD, the initially most commonly used growth technology, required temperatures up to 400 °C and additional annealing in sulfur vapor at 500 °C and in oxygen at 200 °C to suppress the concentration of ionic defects.1241,1243 It was soon realized that such nanostructured solar cell architectures are especially convenient for the ample range of solution-processing technologies. In this direction, CuSbS2 was grown within a 3D architecture from salts, although no device performance was reported.1247 Page et al. fabricated a Cu2−xS-based ETA solar cell, by means of cation exchange of chemical bath deposited CdS, that reached up to 0.07% PCE.1248 Grasso et al. also produced a nanocomposite CIS BHSC with a 5% PCE, by chemical spray deposition as an alternative method to AL-CVD.1242 Fully spray deposited TiO2/ In2S3/CIS devices using deposition temperatures up to 300 °C in air and PCEs up to 7% were reported by Goossens and Hofhuis.1249 Using electrodeposition and an annealing step at 350 °C, Valdés et al. fabricated CISe BHSCs reaching moderate efficiencies.1250 More recently, Nguyen et al. fabricated CISbased BHSCs using nanocrystalline titania, spray pyrolyzed CIS, and a chemical bath deposited Inx(OH)ySz buffer layer, reaching up to 3.2% PCE.1251 Santhosh et al. used chemical spray pyrolysis and low-temperature processing to produce ETA solar cells, combining both CIS and Cu2S as absorbers, that reached 3.82% PCE.1252 Park et al. produced CZTS BHSCs by spin coating CZTS precursor thin films, prepared from metal chlorides and thiourea, followed by sintering at 500 °C with the subsequent addition of CdS by CBD and achieved PCEs of 5.02%.1253 When compared with sputtered CZTS films, the solution processing of the absorber provided an advantageous nanoporous structure, which allowed CdS contact across the complete layer. The performance of such nanostructured solar cells strongly depends on the scaffold nanostructure. For example, O’Hayre et al. reported that the efficiency of TiO2/CIS nanocomposite solar cells increased with an increasing TiO2 particle size, which was associated with improved charge transport with increasing particle size.1254 Following SSSC designs, a blocking layer of the same wide-band-gap oxide material is generally applied between the porous layer and the electrode to prevent the direct contact of the p-type absorber with the n-type electrode.968 Such direct contact would strongly increase recombination, as observed by Motoyoshi et al., who studied the effect of the stacking sequence of the CIS and TiO2 layers on the electronic properties of TiO2−CIS solar cells.1255 To provide rapid avenues for charge carrier extraction, ZnO and TiO2 nanowire and nanotube arrays have been frequently used as ordered porous structures to deposit/grow the nanocrystalline absorber layer. Wang et al. used electrochemical deposition to cover TiO2 nanotube arrays with CISe.1256 Lee et al. fabricated CZTS-based superstrate solar cells, by spin coating molecular precursors onto CdS-coated ZnO nanorods and sintering at 250°C, and achieved 1.2% PCEs.1257 Ionic and molecular precursors are generally preferred to prepare the absorber layer in such kind of solar cells because the introduction of NCs within a 3D porous scaffold is not

7.2. Sintered Nanostructured Solar Cells Based on NCs

The use of nanostructured supports instead of a planar geometry provides both a more effective charge carrier separation and a higher light absorption. This is accomplished by increasing the volume of absorber at a small distance of a charge collecting interface and introducing light scattering. Nanostructured scaffolds were initially used to produce crystalline Cu chalcogenide-based solar cells with extremely thin absorbers, the so-called ETA solar cell. The first ETA solar cells were reported by Kaiser et al. in 2001 and were based on using a CIS absorber layer, TiO2 as the n-type wide band-gap semiconductor and CuSCN as the p-type semiconductor.1240 In ETA solar cells, a transparent p-type conductor such as CuSCN is used to transport holes toward the electrode, and thus, the absorber plays a minor role on the charge transport. An alternative architecture is the so-called 3D p−n junction or 3D bulk heterojunction solar cell, which we will call just the bulk heterojunction solar cell (BHSC). In this architecture, a nanoporous scaffold is filled up with the absorber, which transports one of the charges toward the electrode, avoiding the use of the hole conductor and increasing the total volume of absorber material. A related design is that of semiconductorsensitized solar cells (SSSCs), also called quantum dot sensitized solar cells (QDSSCs), which will be discussed in section 7.4. Compared with the SSSCs, the BHSC is considered more stable even without expensive sealing against oxygen and water. Besides, in solid-state BHSCs, the photocurrent response is 4 orders of magnitude faster than in SSSCs using liquid electrolytes.1241−1243 In the BHSC, the driving force for transport is related to selectivity of contacts, resulting in concentration gradients of electrons and holes throughout the cell, i.e. diffusion.1242 In the first ETA and BHSCs, the absorber layers were, in most cases, annealed/crystallized at temperatures up to 500 °C. However, identical architectures, making use of nanoporous scaffolds, are commonly used as supports of unsintered NC layers, as discussed in section 7.3. Overall, there is a very difuse boundary between the different types of SCs, including sintered thin films, ETA, and BHSC, and unsintered nanocrystalline solar cells, such as nanocrystal solar cells and SSSCs. It should also be pointed out that, while nanostructured, the absorber layer of most sintered ETA and BHSCs is not produced from NCs. However, these solar cell architectures have still been included as a subsection in this review to provide a more complete overview. Nanoporous TiO2/Cu1.8S BHSCs were already reported in 2002 by Reijnen et al.1244 They used atomic layer chemical vapor deposition (AL-CVD) to grow the Cu chalcogenide absorber, while maintaining processing temperatures below 300 °C.1244 Using the same technology, solar cells based on a nanometer-scale, interpenetrating network between n-type TiO2 and p-type CIS were developed by Nanu, Schoonman, and Goossens.1241,1243 An Al2O3 tunnel barrier and an In2S3 buffer layer in between the two materials allowed for a better alignment of the conduction bands and resulted in PCEs up to 4% being achieved. In this direction, the introduction of appropriate buffer layers, such as In2S3, has been regarded as essential to suppress interfacial recombination.1245 5967


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efficient. However, ordered arrays facilitate the deposition of ex situ formed colloidal NCs. In some works, the surface of the oxide arrays can be treated with 3-mercaptopropionic acid, for instance, to promote the uniform deposition of NCs.968 This point will be further discussed in section 7.3. 7.2.1. Outlook of Sintered Solar Cells. Overall, sintered solar cells made of NCs have reached relatively large efficiencies compared with unsintered devices. Currently, deeper analysis of the underlying physics is needed to elucidate their limitations, especially in CZTSSe solar cells.1258−1260 While the solution-processing of large area devices such as PVs has associated important cost advantages, when compared with vacuum-based processes, the benefit of using NC-based inks (instead of salts or molecular-based inks) to produce sintered solar cells is not evident. NCs-based methods are still behind in efficiency compared to salt-based processes, and they cannot compete in terms of cost, if similar annealing processes in a Se atmosphere are used. The preorganization of ions within NCs may be considered to provide an advantage, in terms of composition homogeneity in the final film, but the annealing step used to sinter the NCs into larger crystallographic domains involves an atomic rearrangement and even exchange (selenization) at a much larger scale, where printed ionic or molecular precursors can also be considered homogeneous. Thus, the main competitive advantage of NC inks over ionic/molecular precursors is lost during this crystallization treatment. Also, the sintering process represents a significant part of the total device fabrication cost, especially if controlled atmospheres with toxic components such as Se are required. Furthermore, sintering of the NC-based layer makes it incompatible with some substrates, and it may modify the composition at the nanoscale by inducing phase segregation, evaporation of components, and creation of defects. Thus, technologies that avoid this annealing step have obvious cost advantages over sintered solar cells. At the same time, in alternative strategies not requiring annealing, the use of nanometer-scale crystalline domains, produced at relatively low temperature by colloidal synthesis methods, for example, have a much clearer advantage. An overview of the application of Cu chalcogenide NCs in sintered solar cells is provided in Table 18, with details on the reported cell architecture, NC precursor, deposition and processing conditions, and important device related parameters, such as efficiency (η), Voc, Jsc, fill factor (FF), and the light intensity, which were used to test the resultant PV devices.

nanocrystalline layer for absorption and carrier photogeneration but do not rely on them for the charge transport of both types of carriers (e.g. sensitized solar cells and hybrid solar cells) will be considered in following sections. When using NCs with sizes in the quantum confinement regime, NCSCs are commonly referred to as quantum dot solar cells (QDSCs). Notice that this term sometimes also embraces other solar cell architectures making use of QDs, such as the QD sensitized solar cell (QDSSC), that will be discussed in section 7.4 under the general name of semiconductor sensitized solar cells (SSSCs). Beyond solution processability and third generation concepts, the use of NCs with size in the quantum confinement regime, thus having discrete and size-dependent electronic structure and electron wave functions extending well beyond the NC limits, intends to facilitate charge transport within the absorber layer. At the same time, QDs have enhanced absorption cross sections and can provide extended carrier lifetimes. In NCSCs, PCEs are generally limited by the short diffusion length of photogenerated charge carriers within the absorber layers, which contain a large density of grain boundaries. Therefore, the passivation of traps at the NC surface, and at the materials interfaces, and the promotion of charge transport between NCs are key parameters that need to be optimized.1298 PbS QDs are currently the most investigated material in NCSCs. The main advantages of PbS QDs are their large exciton Bohr radius (18 nm), their ease of synthesis with controlled parameters in the extended quantum confinement size range, and their appropriate band gap for solar energy conversion when the NCs are in the quantum confinement regime.633,1299,1300 NCSCs were initially based on simple Schottky configurations, with a p-type QD absorber layer in between a relatively large work function TCO, such as ITO, and a low work function metal, such as aluminum, calcium, magnesium, or silver. This created a rectifying junction which introduced a depletion region, with a built-in field, that promoted photogenerated charge separation and transport in opposite directions.633,1301−1305 This simple configuration presented the additional advantage of having a limited number of interfaces but the inconvenience of illumination at the nonrectifying side of the junction. However, its main drawback was the constraint in PCE that is imposed by the Fermi level pinning at the metal−semiconductor interface, which limited the Voc well below the predicted value on the sole account of the semiconductor band gap.633,1306 To overcome this limitation, the current most efficient NCSCs make use of the so-called heterojunction configuration, which includes n- and/ or p-type doped wide-band-gap semiconductors, such as TiO2 and/or MoO3 as selective charge extraction contacts or filters between the absorber layer and the current collectors.633,1152 With this configuration, PbS NCSCs have recently reached up to 10.7% PCEs.1307 The main limitation of these devices is still the relatively low lifetime of the charge carriers and the reduced diffusion length, which constrains the absorber layer thickness ( 420 nm; 80 mW/cm2 (350 W) Xenon (500 W)

54.8

809.5

1180

1020

2450

7.93 mL (per cm2)

4011

332

11.09 (per cm2)

4147

700

36

1513

2465

81

3620

757.5

Xenon; 100 mW/cm2 (300 W)

Xenon; λ ≥ 420 nm (350 W)

Xenon; λ > 430 nm (300 W) Xenon; λ ≥ 420 nm (300 W) Xenon; λ ≥ 420 nm (500 W) Xenon; λ ≥ 420 nm (300 W) Xenon; λ ≥ 420 nm (300 W) Xenon; λ > 400 nm (300 W) Xenon (300 W)

1857

Xenon; λ ≥ 430 nm (350 W) Xenon (550 W) 6250

450

120

Activity (μmol/h·g)

Xenon (300 W)

Xenon (400 W)

Light source

3.23 at 420 nm

41

20.9 at 420 nm

20 at 420 nm

0.25 at 400 nm

22.6 at 420 nm

14.7 at 420 nm

19.1 at 420 nm

31.8 at 420 nm

14.2 at 420 nm

14.9 at 420 nm

3.7 at 420 nm

Quantum efficiency (%)

1738

325

1737

340

1736

1735

1734

1733

1732

1731

1730

1729

1728

1727

1726

1725

1724

1723

1722

1721

1720

1719

Ref



colloidal

l × t (30 nm × 15 nm) d (50−100 nm)

CIGS NCs

ZnS:Cu HNS

CGIS NCs

H2

H2

H2

H2 H2

Cu NCs

Cu2S NCs

CIS NCs

CZTS film

CGIS NCs

CIS NMP

CNiTS NCs

H2

H2

H2

H2

H2

H2

Cu2S-ZnS NCT

H2

H2

ZCIS NHS

H2

WS2-Au-CIS NCT

CZTS-PtCo NHS

H2

H2

colloidal

Au: l × w (99.6 nm × 17.2 nm) Cu2FeSnS4: t (34.3 nm) d (12 nm)

Au@Cu2FeSnS4 CSN

H2

6012

d (50 nm)

d (3.5−4.0 μm)

hydrothermal

d × l (100 nm × 5−20 μm) d × t (200 nm × 1.2 nm) d (12−15 nm)

electrospinning

hydrothermal

hot-injection

electrodeposition

solvothermal

solvothermal

electroplating

wet chemical

colloidal

CIS t (5 nm)

w × l (8 nm × 50−150 nm) t (1−5 nm)

hydrothermal co-precipitation

0.2−0.6 μm 2−4 nm

TiO2/CuS/NiS NCT CuS/CdZnS NCT

colloidal

colloidal

4 nm

sonochemical

Synthesis method

Particle size

Cu-based photocatalyst

Evolution

Table 23. continued

Ru (1.0 wt %)

Pt

Ru (1.0 wt %)

Pt (1 wt %)

Cocatalyst

Na2S (0.1 M), Na2SO3 (0.1 M) aqueous solution Na2S (0.35 M), Na2SO3 (0.25 M) aqueous solution Na2S (0.35 M), Na2SO3 (0.25 M) aqueous solution Na2S (8.4 g), Na2SO3 (3.15 g) aqueous solution Na2S (0.35 M), Na2SO3 (0.25 M) aqueous solution methanol (10 mL)/water (40 mL) solution Ethanol (20 mL)/water (80 mL) solution Na2HPO4/NaH2PO4 (0.2 mol/L) solution Na2S (0.05 mol/L), Na2SO3 (0.3 mol/L) aqueous solution Na2S (0.25 M), Na2SO3 (0.25 M) aqueous solution Eosin Y and TEOA aqueous solutions at pH 9

Na2S (0.35 M), Na2SO3 (0.25 M) aqueous solution

aqueous solution of methanol (40 mL) Na2S (0.1 M), Na2SO3 (0.02 M) aqueous solution

Na2S, Na2SO4, water

Na2S (0.1 M), Na2SO3 (0.1 M) aqueous solution Na2S (0.1 M) aqueous solution

Reactant solution

50/50

active surface area 0.3 cm2 25/25

200/100

20/50

2/10

1.5 mg/cm2

30/30

100/50

10/50

10/−

50/80

100/100

15/100

10/50

Catalyst mass (mg)/reactant solution volume (mL)

irradiation at AM 1.5G Xenon; λ = (385 −740) nm (300 W) Xenon; λ > 420 nm (300 W) λ ≥ 420 nm

Xenon (300 W)

Xenon; λ = (400 −800) nm (300 W) Halide; λ > 400 nm (150 W) Xenon; 100 mW/cm2 (300 W) Xenon; 100 mW/cm2 (300 W) Xenon (300 W)

Xenon (300 W)

Xenon; 193 mW/cm2 (300 W) Xenon, λ ≥ 400 nm (300 W) Xenon, λ ≥ 420 nm (300 W) Xenon (500 W) - Mercury; λ > 420 nm (1000 W) - LED; 450 nm (30 W) Xenon; 100 mW/cm2 (150 W)

Light source

Ref

2028

316

50.6

383

1750

1749

1748

87 per 1 cm2

1746

1294

1745

1744

1743

1006

1252

1741 1742

1740

1739

925

1747

8.48

6.9 at 560 nm

Quantum efficiency (%)

273.25

1430.4

1250

38.9

2232

250

1850

90

800 3520

3380

6.7

1728




Chemical Reviews

Review

Figure 69. Single-excitation and two-step photocatalytic water splitting processes using one type of NC with both O2 and H2 evolution sites (a); two types of NCs, one for each reaction, contacted by means of a redox shuttle (b); two types of NPs separated by a membrane and combined with two redox couples (c); a supported photoactive material (d) or a semiconductor-sensitized material (e) where one reaction, e.g. oxygen evolution, takes place and which is electrically connected to, for instance, a Pt electrode for hydrogen evolution; or two supported photoactive materials or sensitized layers electronically connected (f, g). Cartoons inspired by refs 1751 and 1752.

Figure 70. Scheme of MPA-coated CIS QDs loaded with Ru as a hydrogen evolution catalyst (left) and supported on TiO2 nanoparticles (right). Reproduced with permission from ref 1734. Copyright 2012 Elsevier B.V.

9.1.1.2. Photoelectrochemical Water Splitting. Supported Cu chalcogenide NCs and nanostructured materials have also been tested for the PEC water splitting (see Table 24). CIS QDs were used by Li et al. as a sensitizer of a TiO2 photoanode for water splitting, where an energy conversion efficiency of 1.9% was under 100 mW cm−2 simulated solar illumination (at +0.23 V bias in a PEC cell) with an aqueous sulfide/sulfite electrolyte, compared to 1.2% efficiency with no bias.872 In subsequent work, the same group reached saturation photocurrents of up to 16 mA cm−2 and light-to-chemical energy conversion efficiencies up to 13% at a +0.34 V bias.1756 In this study, the same architecture was used, but a CdS layer was deposited by SILAR on top of the TiO2/CIS-QD layer, which was produced by attaching colloidal CIS QDs to TiO2 by means of MPA. The CdS coating was reported to extend the absorption spectra of the CIS QDs and to facilitate photogenerated charge carrier separation. Recently, Guijarro et al. reported the use of CIGS photocathodes, produced from the printing of CIGS NCs, for solar water splitting and reached stable photocurrents, saturating at 8.0 mA cm−2 and onsetting at 0.6 V vs RHE.1757 The initial studies looked at directly grown CIS films,1759 but more recent studies have investigated CIS doped films (with Zn, Sb, and Ni) and tested their potential for hydrogen production.1731 The authors observed that the incorporation of the dopants resulted in higher H2 evolution rates compared to CIS alone, with Zn providing the highest performance. As an example, Luo et al. used a solution-processed nanocrystalline CIS layer to reach water splitting photocurrents of 3.5 mA cm−2

crystal structure, kesterite or wurtzite, does not have a significant effect on the H2 generation rates, but the composition, and particularly the Cu/(Zn+Sn) ratio, strongly influenced the rate, with the Cu rich CZTS NCs providing the highest performances. In subsequent work, the same authors demonstrated that the incorporation of PtCo on CZTS resulted in even higher evolution rates, of up to 1850 μmol h−1 g−1.1006 Recently,. Gonce et al. analyzed the photocatalytic hydrogen evolution of Cu2XSnS4 (X = Zn, Ni, Fe, Co, Mn) nanofibers using a dye as the photosensitizer, reaching up to 2028 μmol h−1 g−1 for Cu2NiSnS4.383 Ha et al. demonstrated Au/Cu2FeSnS4 NCs provided reasonable photocatalytic hydrogen generation rates without the addition of a catalyst, largely owing to a Au-related plasmonic enhancement.387 Recent work has further looked into the mechanisms of these reactions, as a function of the facet types and particle shape, with Li et al. reporting some interesting observations of hydrogen evolution on Cu2WS4 decahedra-shaped particles.325 Some approaches have investigated modifications of the reaction chemistry. For example, Iwashina et al. used a Z-scheme for the complete water splitting reaction, instead of using a sacrificial agent.368 In this work, reduced graphene oxide (RGO) was used as the electron mediator between the H2 evolution material, TiO2, and the O2 evolution materials (of which CIS, CGS, CZTS, and Cu2ZnGeS4 loaded with Pt were investigated). This Z-scheme system achieved an efficiency of 0.023% with initial H2 evolution rates close to 400 μmol h−1 g−1 for CGS. 6013


6014

d (300−500 nm)

100−200 nm

d (50−150 nm)

600−650 nm

t (20−30 nm)

12.7 nm

5.1 nm

CIS NCs

CITS powder

CIS NRA

Cu9S5 HNS

CIS NRA

CISe NCs

(Cu2Sn)x/3Zn1−xS NCs

G/n-CZTSe/Au

Ag/AgCl Pt

20−25 nm

CZTSe NCs

solvent

Ag/AgCl Pt

colloidal

CZTS: 12 nm Ag2S: 4 nm

CZTS-Ag2S NCs

G/FTO/CZTSAg2S

Ag/AgCl Pt

solid-state reaction G/ITO/CAgGS/ Ru

Pt

SCE

WS2-Au-CuInS2

CAgGS powder

electroplating

t (5 nm)

CIS NCs

Ag/AgCl Pt

G/ITO/Gr/CIS/ GrO

hydrothermal

t (15−20 nm)

CIS NCs

Pt

Ag/AgCl Pt

G/FTO/Mo/CIS/ Ag/ AgCl/ CdS/AZO/TiO2 sat. KCl

Ag/AgCl Pt

1765

∼5 (at 1.2 V vs Ag/AgCl) NaCl (1 M) aqueous solution

Solar energy efficiency: 2.81%

0.58 (at 0.5 V vs Ag/AgCl)

Xenon; λ > 420 nm (300 W) Eu(NO3)3 (1 M) aqueous solution

Xenon (300 W)

341

4.5 (at −0.9 vs Ag/ IPCE: 7.7% at 480 nm under −0.4 V AgCl) vs Ag/AgCl

Xenon; λ > 420 nm (300 W) K2SO4 (0.2 mol/L) aqueous solution at pH 6.9

1755

1745

0.005

Xenon; λ > 420 nm; 100 mW/cm2 0.5 M Na2SO4

Na2S (0.1 mol/L)

1764

368

1763

1748

Xenon; 100 mW/cm2 2.47 (at 0.18 vs Ag/AgCl)

Solar energy conversion: 0.023%

IPCE: 20% at 520 nm at 0 V vs RHE

Xenon; 100 mW/cm2 3.5 (at −0.3 V vs (450 W) RHE)

K2SO4 (0.1 mol/L), Na2HPO4 (0.025 Xenon; λ > 420 nm mol/L), KH2PO4 (0.025 mol/L) (300 W) aqueous solution at pH 6.9

solution of Na2SO4 (0.5 M) and KH2PO4 (0.1 M) at pH 5.0

0.3 (at 0.95 V vs Ag/AgC)

Onset potential: 0.63 VRHE Solar conversion efficiency: 1.63% IPCE: 45−50% at 400−700 nm

1762

Xenon; 100 mW/cm2 0.31 (at 0.3 V) (300 W) −9.3 (at 0 VRHE)

1735

Xenon; 100 mW/cm2 22.4 (at 0 V vs Ag/ Onset potential: −0.7850 V AgCl)

phosphate buffer solution (0.2 mol/L AM 1.5G of Na2HPO4 /NaH2PO4)

Na2SO4 (0.2 M) aqueous solution

Ag/AgCl Pt TiO2/ (Cu2Sn)x/3Zn1−xS G/Mo/CZTS/ CdS/In2S3/Pt


Ag/AgCl Pt

TiO2/ZnS/CdSMn/CISe/ZnS

1737

1761

flat band potentials Energy conversion efficiency: 1.35% 1.12 and 1.06 V vs NHE

1760

1155

1759

Xenon; 100 mW/cm2 19.07 (at 0.74 V vs Hydrogen generation efficiency: Ag/AgCl) 11.48% at 0.5 V vs Ag/AgCl

Xenon; (300 W)

10.5 (at 0 V vs Ag/ IPCE of 57.7% at 480 nm AgCl)

Xenon; λ>420 nm; 150 mW/cm2 (300 W)

Na2S (0.5 M), Na2SO3 (0.5 M) aqueous solution Na2SO4 (0.5 M) aqueous solution

3.52 (at 0 VAg/AgCl)

100 mW/cm2

Na2S (0.25 M) aqueous solution

KOH (1 mol/L)

Hg/ Pt Hg2Cl2

Ag/AgCl Pt

Ag/AgCl Pt

0.072 (at −0.45 V) Photoconversion efficiency: 0.24%

1731

Xenon; 100 mW/cm2 −8.58 (at 1.0 V vs (300 W) Pt)

K2SO3 (0.25 M), Na2S (0.35 M) aqueous solution

1758

1756

872

Xenon; 500 mW/cm2 16.9 (at 0.29 V vs hydrogen generation efficiency: 3.2% SCE) (0.29 V vs SCE)

Voc: 0.57 V Conversion efficiency: 13% at +0.34 V

Voc: 0.52 V

Na2S (1 mol/L)

Ag/AgCl graphite Na2SO3 (0.1 mol/L) aqueous solution Xenon; 50 mW/cm2 (500 W) rod (cathode)

Ag/AgCl Pt

Pt

16

1

Ref

100 mW/cm2

100 mW/cm

2

Light source

Photocurrent density (mA/cm2) (potential) Other



Electrolyte

G/FTO/TiO2/CIS Ag/AgCl Pt

G/FTO/Cu9S5

G/FTO/ZnO/ ZnS/CdS/CIS

SiO2/Mo/CITS

G/FTO/CIS

solid-state reaction G/FTO/CZTS

solvothermal

electrodeposition

solvothermal

colloidal

SILAR

hydrothermal

ion exchange hydrothermal method

mechanical ballmilling

solvothermal

CBD

CZTS powder

CIS films

t (5 μm)

G/ITO/Sb:CIS

t (1.38−1.86 μm)

CIS:Sb films

CZTS films

G/ITO/ZnO−CIS SCE

CIS t (10−20 nm) hydrothermal and cation exchange

Pt

ZnO@CIS NRA

solvothermal

G/FTO/TiO2/ CIS/CdS

3.5 nm

CIS QD

Pt

G/FTO/TiO2/CIS

4.3 nm

CIS QD

solvothermal

NC size

Cu-based NC

Working electrode Ref elec- Counter Synthesis method structure trode electrode

Table 24. Summary of Cu Chalcogenide NCs Used in Photoelectrochemical Water Splitting (PEC)



Review

0.01 (at 0.3 V vs RHE) 100 mW/cm2

1766

with a FTO/Mo/CIS/CdS/AZO/TiO2/Pt architecture.1763 W. Bo et al. used graphene/CIS/graphene oxide nanocomposite films for PEC water splitting, reaching up to 2.47 mA cm−2 photocurrent densities. Other Cu chalcogenide nanostructures, such as CIGS,1755 CZTS,341,380 CZTSe,1765 or Cu2CoSnS4 nanofibers380 have also been tested for PEC water splitting. Of particular significance, Yu et al. demonstrated that CZTSAg2S colloidal NCs, incorporating p−n nanoheterojunctions, facilitated charge carrier separation and thus increased photocurrent densities.341 Lee et al. added to the mechanistic understanding by comparing the PEC properties of n-type and p-type nanocrystalline CZTSe photocathodes, concluding that the n-type material provided larger photocurrents but much lower stability than the p-type material.1765 As in the case of PV, highly crystalline layers usually provide improved performances over nanocrystalline ones due to better charge transport properties and reduced density of recombination sites. In this regard, highly crystalline, Cu chalcogenide layers for PEC water splitting have been produced by electrodeposition,1748 sputtering,1767,1768 and coevaporation1769 and from micron-sized crystals produced by solid-state methods.1770 NCs find similar advantages here as in the PV field, namely their potential for low-cost production and for higher interface areas for better charge carrier separation. To take advantage of the latter, ordered nanorod/nanowire arrays have also been tested as supports for the photoabsorber/ photocatalytic materials. Li et al. investigated ZnO/CIS nanorod arrays prepared by a combination of hydrothermal and cation exchange methods for PEC water splitting reaching photocurrents close to 17 mA cm−2 and hydrogen generation efficiencies of 3.2% at +0.29 V in a Na2S electrolyte.1758 Yu et al. used ZnO/ZnS/CdS/CIS core−shell nanowire arrays produced by cation exchange reactions to reach photocurrents up to 10.5 mA cm−2 in a Na2S Na2SO3 electrolyte.1760 Instead of ZnO, Guo et al. used hierarchical TiO2 to produce TiO2/CIS nanoarrays (Figure 71) which provided photocurrent densities up to 19.07 mA cm−2 and efficiencies for hydrogen generation of up to 11.48% in a KOH electrolyte.1761 Also noteworthy is work by Cheng et al. where a WS2/Au/CIS photocatalyst was produced by loading WS2 nanowire arrays with Au NCs and afterward coating them with CIS NCs.1745 All the components in such heterostructures were observed to contribute to reach H2 production rates of up to 38.9 μmol h−1 g−1 in a Na2S- and Na2SO3-containing electrolyte. The influence of shape was further investigated by Sheng et al., where CISe nanospheres and TiO2 nanotube arrays were used to fabricate photoelectrodes with the architecture TiO2/ZnS/CISe/Mn-CdS/ZnS, which yielded photocurrents of up to 22.4 mA cm−2 and hydrogen evolution rates of 7.93 mL cm−2 h−1.1735 In related work, Zhang et al. supported (Cu2Sn)x/3Zn1−xS NCs on TiO2 nanotube arrays and studied the photocurrent dependence on the conduction band offset determined by composition.1762 Interestingly, Tan et al. used Cu/CuS-ZnO/ZnS branched nanowires for PEC water splitting and observed that sonication provided higher photocurrent densities, which was due to a self-bias associated with the piezoelectric voltage created at the ZnO.1294 A wide range of viable constructs is possible with, for example, Chong et al. reporting the high performance of TiO2/CdSe/Cu2Se cascade architecture on a TiO2 nanotube array structure (as photoanodes), where stable photocurrent densities of up to 20 mA cm−2 were demonstrated.1771

d (80 nm) CIS NRA

template-assisted

G/CEp/Au/CIS/ CdS/ZnS

Ag/AgCl Pt

Na2SO4 (0.5 M) aqueous solution

1750 7.8 × 10 Xenon; λ > 420 nm (300 W) Na2SO4 (0.5 M) aqueous solution d (3.5−4.0 μm) CIS NMP

hydrothermal

G/ITO/CuInS2MoS2

Ag/AgCl Pt

−3

NC size Cu-based NC

Table 24. continued

Working electrode Ref elec- Counter Synthesis method structure trode electrode

Electrolyte

Light source

Photocurrent density (mA/cm2) (potential) Other

Ref

Chemical Reviews

6015


Chemical Reviews

Review

Figure 71. Scheme of hierarchical TiO2/CIS nanoarrays, the measured photocurrent density, and the calculated photoconversion efficiency. Reproduced with permission from ref 1761. Copyright 2014 Royal Society of Chemistry.

9.1.2. Photocatalytic Degradation of Pollutants. The photocatalytic degradation of pollutants, and particularly dyes, has been extensively investigated, both for its intrinsic interest as an environmental remediation strategy and also as a rapid assessment of the material potential to take advantage of solar energy. In the first direction, it is worth keeping in mind that solar energy can be used not only to produce electricity or solar fuels, but also to help improve water and air quality by means of photocatalytic processes. An effective environmental remediation through photocatalysis requires materials or composites to be able to absorb light in the visible part of the spectrum and also to be able to catalyze the pollutant decomposition reaction. In particular, the quality of water, a vital component for life and also essential in most industrial activities, has been jeopardized in numerous locations with high population densities and intensive industrial activity. One main water contaminant caused by human activity is nonbiodegradable synthetic dyes currently used in the textile industry, paints, printing inks, papers, plastics, and a large number of consumer products. Synthetic dyes, with richer and longer lasting colors, replaced natural dyes used since 3500 BC. Currently these chemical dyes are leaving one of the deepest water footprints worldwide, particularly in India and China. Textile dyeing and treatment is alone estimated to contribute up to 20% of the total industrial water pollution. Synthetic dyes strongly affect the aquatic ecosystem and, consequently, the whole food chain not only by blocking sunlight and stopping reoxygenation but also because some of them contain toxic elements, such as chromium. Other dyes are known as carcinogenic agents and can cause chronic toxicity to animals which come in contact with them by oral or skin administration. Several physical, biological, and chemical removal techniques, such as adsorption, coagulation, flocculation, membrane filtration, ozonation, electrochemical, radiolysis, bacterial, algal, fungal, and advanced oxidation process techniques, have been applied for their removal. However, this is not an easy task, especially taking into account the important cost constraints and that a large variety of compositions exist, with more than 3500 different types of textile dyes alone. In this regard, the use of low-cost photocatalytic materials activated by solar light, such as Cu chalcogenides, represents a potential strategy for dye and more generally organic or even inorganic pollutant degradation.1772−1781 Besides, the development of photocatalysts and even PV materials with new compositions, shapes, or crystallographic phases has often been supported on performance feedback obtained from measures of dye degradation rates. The motivation is that these measurements are easy to follow using widely-available UV−vis spectroscopy. However, in general, dyes are inappropriate as model compounds for the evaluation of photocatalytic activity,1782 and the results of these analyses

should not be directly correlated to the material potential for photocatalytic water splitting, PV, or other relevant applications. Table 25 shows the results of the photodegradation tests found in the literature using Cu chalcogenides. The photodegradation of dyes and other organic molecules having an optical fingerprint was tested using a variety of Cu chalcogenide compositions and shapes, including CuS NCs,1783,1784 CuS nanotubes,1785,1786 CuS hierarchical nanostructures,802,1787,1788 CuS, Cu2S, and Cu1.8S nanostructured flowers,801,1789,1790 Cu2S microsponges,1791 CuS nanostructures,759,1792−1798 Cu2S NPs,1799−1802 ZCIS NPs,850 ZCIS nanorods,898 CuSe nanoflakes,1803 solid and hollow CISe NPs,1804 Cu2Se nanowires,1805 Cu2FeSnS4 nanostructured spheres, 388 CuSe 1−x S x nanoflakes, 1806 and Cu 2 ZnGeS 4 NCs.365 Interband materials, such as Sn-doped CGS and CIS NCs1155 and Cr-doped CGS,1598 have also been studied and are a particularly interesting class of materials, which have enhanced performances mediating dye photodegradation. In some systems, improved photodegradation properties were obtained when combining Cu chalcogenide NCs with RGO,1810,1816,1817,1828 or with polymers such as pectin,1812 or by using semiconductor−semiconductor heterostructures such as CdS/Cu2S nanorods,1818 CuS-MoS2 nanocomposites,1832 or CuS/ZnS core/shell NCs.1823 Cu chalcogenide NCs have also been used as optical or/and catalytic sensitizers of wideband-gap semiconductors, such as TiO2. In this direction, several architectures have been tested, including CuS QDs supported on TiO2 nanotubes,1808 CuxS/TiO2 nanocomposites,1809 CuS-TiO2 nanofibers,1833 CZTS-TiO2 composites,1834 CIS-TiO2 nanotube arrays,1822 CISe-TiO2 nanotube arrays,1835,1836 and CuS/ZnO heterostructured nanowire arrays;1811,1827 Cu2S@ZnO heterostructures,1813,1815 and nanocomposites;1820 composites formed by RGO, CIS, and TiO2,1814 and RGO-CIS-ZnO nanocomposites;1837 and CuS-WO3 composites.1826 In addition, semiconductor−metal nanostructures, such as Cu2S-Au1807 or CZTS-Au heterostructures,1819 or semiconductor−semiconductor−metal nanostructures, such as Ag-CuBiS2-TiO2 composites,1830 have been produced and tested for the degradation of dyes and other pollutants. The plasmonic properties of Cu chalcogenides have also been exploited to improve the performance for the photocatalytic degradation of pollutants or dyes. As an example Han et al. demonstrated that plasmonic Cu2−xSe-graphitic C3N4 were highly efficient for the degradation of methyl blue in aqueous solution.1838 9.1.3. CO2 Photoreduction. The sunlight driven photocatalytic reduction of CO2 can be a partial solution to solving a double problem, the greenhouse effect and the storage of sunlight into more directly usable forms of energy, such as alcohols or hydrocarbons.1839,1840 In this direction, Cu chalcogenide NCs have been used for the photoelectrochemical 6016


d (4.8 nm)

d (3.6 nm)

ZCIS NCs

CuS QD

6017

SILAR

hydrothermal

hydrothermal

d (5 nm)

d (30 nm)

d (400−600)

d (1.5−2 μm)

d (20 nm)

d (100−200 nm)

CuS/rGO NCT CuS NCs

CuS NCs

Cu2S NMP

Cu2S NMP

CIS NCs

CGTS NCs

hydrothermal

sonochemical

solvothermal

d (18 nm)

40−80 nm

d (10−25 nm)

d (2−5 μm)

CuS NCs

CuS NCs

CuS NCs

CuS powder

hydrothermal

wet chemical

CuS NCs

CGS powder

solid-state reaction and ball-milling solid-state reaction

electrodeposition

coprecipitation

microwave assisted solvothermal hydrothermal

CuS NTA

ZCIS NRA

colloidal

photochemical

stepwise crystallization

l × w (5.2 nm × 7.6 nm) d × l (100 nm × dozens of μm) 18 nm

CuxS powder

hydrothermal

d (1.8−2.4 μm)

CuS NCs

colloidal

solvothermal

d (9 nm)

Cu2S NCs

solvothermal

d (7 nm)

Au−Cu2S NCs

CuS NMP

RhB

MB

MB

MB

MB

MB

MB

2,4-D

MB

phenol

MB

AO7

MB

MB

RhB

MO MB

MB phenol

RhB

RhB

1,4-dioxane

MB RhB

MB

eosin Y

l × id × od (several μm hydrothermal × 235 nm × 300 nm) 1 μm solvothermal

CuS NTA

Pollutant safranine

Synthesis method

wet chemical

Particle size

CuS NCs

Cu-based NCs Light source

catalyst, solution of MB (50 mL, 80 mg/L), H2O2 (2.5 mL) catalyst (30 mg), MB aqueous solution (40 ml, 20 mg/L) MB aqueous solution (50 ppm), H2O2, catalyst (50 mg) catalyst (10 mg), solution of RhB (200 mL, 2.5 mg/L), H2O2 (0.1 mL, 30%)

catalyst (7 mg), aqueous (18 mL), RhB solution (3 mL, 0.1 mM) catalyst (20 mg), MB solution (50 mL, 20 mg/L) catalyst (1 mg), aqueous solution of MB (10 mL, 4 mg/L) catalyst (film 2 cm × 2 cm), aqueous AO7 solution (50 mL, 1 × 10−5 M) initial dye concentration = 100 mg/L, catalyst dosage = 0.5 g/L catalyst (30 mg), aqueous solution of phenol (30 mL, 100 mg/L) catalyst (yielded from one experiment), MB solution (100 mL, 10 mg/L) 2,4-D water solution (60 mL, 10 mg/L, pH 3.2) catalyst (100 mg), MB solution (200 mL, 10 mg/L) catalyst (100 mg), MB solution (100 mL, 1 mg/L) catalyst (20 mg), MB (50 mL, 2 × 10−5 M)

catalyst, MO/MB solution (25 mL, 0.0125 mM), H2O2 (30%, 4 ml/L)

catalyst (25 mg), MB/phenol aqueous solution (0.3 mM, 100 mL)

catalyst (50 mg), aqueous dioxane solution (50 mL, 2.4 mM) catalyst (10 mg), aqueous RhB solution (2.0 × 10−5 M, 50 mL) catalyst (5 mg), RhB solution (0.02 mM)

H2O

81

Xenon; λ ≥ 420 nm (300 W) Xenon; λ ≥ 650 nm (500 W) Tungsten; λ ≥ 400 nm (100 W) UV light

∼90

30

100

∼42 natural light Xenon; λ > 400 nm (20 W)

40

87

80

30

180

90

180

40

150

80

180

120

visible light

94

87

65

96.2

91

100

97

60

81

80

120 360 180 300 120

∼75 ∼85 99 99 91 40

120

60

180 180 300

90

60

Degradation time (min)

∼100

100

60 55 60

90

∼15

Degradation rate (%)

Xenon (500 W)

Xenon; λ ≥ 420 nm (350 W) daylight lamp (18 W)

Xenon; λ > 400 nm (500 W) Xenon; λ > 400 nm (300 W) solar light

Halolite; λ > 430 nm (500 W) Mercury; UV light

UV light

Mercury; λ > 420 nm (500 W) Mercury (450 W)

Halogen; λ > 420 nm (100 W) Mercury; UV light (300 W)

BUV-30; λ = 254 nm (30 W) catalyst (100 mg), aqueous solution of eosin UV lamp; λ =365 nm Y (10−4 M) catalyst (30 mg), MB solution (40 mL, natural light 20 mg/L), H2O2 (1.3 mL) catalyst (5 mg), aqueous MB/RhB solution Halogen; λ > 420 nm (50 mL, 5 μM) (100 W)

Reactant solution

Table 25. Literature Survey of the Photocatalytyic Degradation of Pollutants Using Cu Chalcogenide NCs

1793

1817

1792

1816

1815

1598

1155

1814

1813

1791

1812

1811

1810

1786

898

1809

1808

850

801

1807

1807

802

1785

1783

Ref



6018

solvothermal

solvothermal

ca. 2.3 μm

d (2.5−3 μm)

CZTS: ca. size (12 nm) colloidal Pt: ca. size (2 nm) t (305 nm) electrochemical

d (22 nm)

CuS NMP

CuS NMP

CZTS-Pt NCs Cu7S4 films

Ag2S:Cu NCs

CuS-Cu2S NCs

CuS-rGO NCT Cu2S NCs

d (30 nm)

grain size (11 nm)

hydrothermal

d (∼10−20 nm)

CuS/WO3 composite CuS NCs

hydrothermal

d (1.5−2 μm)

Cu1.81S NMP ZCuS NCs

MB MG

MB

solvothermal

chemical reduction

MB

MB

MB

Ofloxacin antibiotic

RB

MB MO R6G RhB

MB

RhB

MB

MO CR

RhB

2-CP

CO2

MB

MB

MB

MB

MB RhB MB + RhB

Pollutant

hydrothermal

SILAR

hydrothermal

coprecipitation

colloidal solvothermal

d (20−25 nm) d (5−7 nm)

CISe NCs Cu2S NCs

solvothermal

electrodeposition

CBD

t (few nm)

l × d (2 μm × 60−70 nm) d (25 nm)

CuS NCs

CIS NCs

CuS/ZnO NCT ZCIS NRA

CZTS: l × w (28 nm × colloidal 8 nm) Au: d (2.1 nm) Later size (200-800 hydrothermal nm) t (15-40 nm) 35 nm mechanical

cation exchange

CdS-Cu2S NRA CZTS-Au NHS CuSe NCs

Synthesis method

solvothermal

Particle size

CuS NCs

Cu-based NCs

Table 25. continued Light source

light catalytic experiment

Xenon; λ > 420 nm; 20 mW/cm2 (150 W) natural light

sunlight visible light Fluorescent (250 W) daylight lamp (18 W)

Xenon; 100 mW/cm2 (300 W) Xenon; λ = (420−800) nm; 35.5 mW/cm2 (500 W) Xenon (500 W)

visible light

natural sunlight

LED illumination; λ = 405 nm λ = 420 nm

catalyst (10 mg), aqueous (50 mL), RhB Xenon (300 W) solution (10 ppm) aqueous (40 mL), RB solution (10 mL, Tungsten (200 W) 10−4 M), H2O2 (1 mL, 30%) catalyst (1 g/L), Ofloxacin antibiotic aqueous UV light; λ = 365 nm solution (50 mL, 10 mg/L at pH 6.5) (125 W) Halogen (1000 W) catalyst (2.5 g/L), MB solution (75 mL, Tungsten-Halogen 2 × 10−5 M) (100 W) catalyst (2 cm × 2 cm), MB aqueous solution Xenon; λ > 400 nm (50 mL, 10 mg/L) (500 W) catalyst (30 mg), MB suspension (50 mL, Tungsten (150 W) 20 mg/L), H2O2 (30%, 2 mL) catalyst (50 mg), MB solution (10−5 M, UV light; λ = 365 nm 250 mL) (48 W) visible light; λ > 420 nm catalyst (10 mg), aqueous solution of dye sunlight (25 mL, 10−5 M)

catalyst (50 mg), RhB aqueous solution (200 mL, 10 mg/L), H2O2 (30%, 2 mL) MO (30 mL, 20 mg/L) aqueous solution of CR (50 mL, 0.5 × 10−3 M) catalyst (yielded by one experiment), MB solution (100 mL, 10 mg/L) catalyst (1 mg), aqueous solution of RhB (1 ml, 10−5 M) catalyst (20 mg), aqueous (100 mL), MB solution (3 mL, 400 mg/L), H2O2 (3 ml) catalyst (5 mg), dye solution (30 mL, 5 mg/L), H2O2

2-CP (100 mL, 20 mg/L), Na2SO4 (0.01 M)

NaOH, water

solution of MB in H2O (10−5 M), methanol, NaCl (∼30 mM) solution of MB (2.5 mL, 2.4 × 10−5 M), alcohol (methanol or ethanol):water (1:6) catalyst (10 mg), MB solution (40 mL, 4 × 10−5 M), H2O2 (1 mL) catalyst, MB solution

catalyst (20 mg), aqueous (100 mL), dye natural light solution (2.8 mL, 400 mg/L), H2O2 (30%, 2 mL)

Reactant solution

61.95 90.25

99.7 99.3

99.27

100 40

18 100

60

120

60

∼80 100

120 140

150

150

25

48

120

60

90 135

56

180

30

25

10

50 60 56 30


90 95

81

98.1 8 98 95

93

50

83

93 90

96

81.95

96.5

99

45

96 94 93 85


1829

1801

1828

1827

1826

1825

1824

340

1789

1788

1823

1790

1804 1799

1794

1822

1821

1820

1803

1819

1818

1787

Ref



solvothermal

d × l (140−300 nm × 660−840 nm) d (142 nm)

1 μm

d × l (100 nm × 5−20 polyrol μm) l × t (200−600 nm × hydrothermal 15−50 nm)

Cu2Se NWA

CuS NMP

Cu NWA

6019

CuSe1−xSx NCs

solutionimmersion hydrothermal

solvothermal

5−6 μm

CITS NMP

Cu2S NCs

hydrothermal

w (40−110 nm)

CuS NCs

Synthesis method

colloidal

Particle size

CuBiS2 NCs

Cu-based NCs

Table 25. continued

MB

RBl MB MO MO

MB

MG

catalyst (10 mg), aqueous (15 mL), MO solution (0.04 mM), H2O2 (0.1 mL) catalyst (10 mg), MB solution (40 mL, 2 × 10−5 M), H2O2 (1 mL, 30%)

catalyst (100 mg), RhB solution (100 mL, 10−5 mol/L) catalyst (20 mg), MG solution (20 mL, 2.5 × 10−5 M) G/ITO/CdS/Cu2S slide, MB solution (0.025 M) catalyst (5 mg), aqueous dye solution (20 mL, 10−4 M), H2O2

4-nitrophenol RhB

aqueous solution of pollutant, catalyst

nitrobenzene

Reactant solution

catalyst (10 mg), AB 1 solution (50 mL, 10 ppm) catalyst (20 mg), AB 1 solution (100 mL, 5 ppm) aqueous solution of pollutant, catalyst

MO MV RhB AB 1

Pollutant

Xenon; UV−vis; 100 mW/cm2 (300 W) Xenon; λ > 400 nm; 20 mW/cm2 (150 W)

natural light

Xenon; λ > 400 nm (450 W) Tungsten; visible light (150 W) Tungsten; visible light (150 W) Xenon; λ > 420 nm (300 W) Xenon; λ ≥ 420 nm (500 W) Xenon

Halogen; UV light (150 W)

Light source

99

15

120

300

40

∼100 80−90 90 80−90 ∼100

20

180

60

60

30

180 60 80 5


∼100

73

100

100

99

9.39 85.03 70.16 99


1806

1294

1795

1831

1805

388

1796

1796

1830

Ref



1845

1848

(see Table 26) and photocatalytic reduction of CO2 (see Table 27). Yang et al. used Zn:CIS nanoarrays for the photoelectrochemical synthesis of methanol from CO2 hydrogenation.1821 The authors demonstrated that Zn doping was the key to optimizing the band alignment, and thus, photocurrent densities up to 2.02 mA cm−2 and methanol production rates up to 0.032 M h−1 were reached. In the same direction, Singh et al. studied a TiO2-CIS composite photocatalyst for the photocatalytic reduction of CO2 into solar fuels.1841 The authors identified energy levels responsible for photocatalytic reduction of CO2 using scanning tunneling microscopy/spectroscopy. They experimentally demonstrated selectivities above 70% for ethane production and up to 4.3% efficiency in converting UV into fuels using concentrated sunlight, which is above the values obtained for TiO2−Pt nanocomposites. A range of fuels have been generated using Cu chalcogenide NCs for the mediated photocatalytic reduction of CO2. For example, Manzi et al. studied the photocatalytic CO2 reduction to CO and methane using Cu2SPt hetero-NCs (Figure 72)1842 and achieved evolution rates up to 3.02 and 0.13 μmol h−1 g−1, respectively. In a separate report, Kar et al. used Cu2S and CuS nanorods for the photoreduction of CO2 to CH4, reaching production yields up to 38 μmol m−2 h−1 under AM 1.5G irradiation.1843 In the same line, Adam et al. used nanostructured CIS films for the photoelectrocatalytic CO2 reduction to CO, reaching faradaic efficiencies above 20 %.1844 Yuan et al used CIS-graphene nanocomposite thin film electrodes for the photoelectrochemical reduction of CO2 to methanol.1845 Liu et al used CuxAgyInzZnkSm solid solutions, including Ru and Rh based co-catalysts for the photocatalytic CO2 reduction to methanol, reaching methanol yields up to 118.5 μmol g−1 h−1.1846 Nanostructured CIS thin film photocathodes were also used for the solar-driven photoelectrochemical reduction of CO2 to methanol by Yuan et al., reaching faradaic efficiencies up to 97 % at an overpotential of 20 mV.1847 The same authors also studied the effect that the presence of pyridine has on the surface of CIS in this reaction.1848 9.1.4. Other Photocatalytic Applications. Very few works have explored the potential of Cu chalcogenides for other photocatalytic reactions beyond water splitting and pollutant degradation and CO2 reduction. As an example of alternative reactions, Cui et al. demonstrated the advantages provided by the plasmonic properties of copper sulfide nanostructures in the field of photocatalysis.1849 They showed that Cu7S4−Pd heteronanostructures improved sunlight energy conversion efficiency in photocatalytic reactions such as the Suzuki coupling reaction, hydrogenation of nitrobenzene, and oxidation of benzyl alcohol.

−0.05 (at 0.59 V)

2.6 (at 0.59 V)

Xenon; 100 mW/cm2

Xenon; λ = (350−650) nm; 100 mW/cm2

−0.2 (at 0.54 V) 100 mW/cm2

Graphite

G/ITO/CIS

G/ITO/CIS/Gr

electrodeposition

electrodeposition

crystal size (38.4 nm)

25 nm

CIS film

CIS NCs

SCE

Graphite

G/ITO/CIS electrodeposition 0.5 μm CIS NMP

SCE

Working electrode structure Synthesis method Particle size

SCE

Graphite

0.1 M acetate buffer solution, pyridine (10 mM), pH 5.2 0.1 M acetate buffer solution, pyridine (10 mM), pH 5.2 0.1 M acetate buffer solution, pyridine (10 mM), pH 5.2

Photocurrent density (mA/cm2) (potential)

methanol concetration 1.18 mM after 10 h at potential of −0.54 V methanol concentration 1.7 mM after 7 h at potential of −0.59 V methanol concentration 0.9 mM after 10 h at potential of −0.59 V

1847

Review

Cu-based NC

Table 26. Photoelectrochemical Reduction of CO2

Ref electrode

Counter electrode

Electrolyte

Light source

Others

Ref

Chemical Reviews

9.2. Other Catalytic Applications

Beyond their excellent optical and optoelectronic properties that make them interesting materials for photocatalytic applications, Cu chalcogenides are also an exceptional catalyst in a number of reactions not involving light. Actually, coppersulfide clusters have an important biological role as catalytic centers in nitrous oxide reductase, which reduces nitrous oxide to dinitrogen.1850−1854 Technologically, electrocatatlytic properties of Cu chalcogenides are also used in several biomedical sensors,1855 although this application will be discussed in section 13. In this subsection, we will focus on nonbiologically related catalytic applications of Cu chalcogenide NCs. 6020


Chemical Reviews

Review

Table 27. Photocatalytic Reduction of CO2 Cu-based Particle photocatalyst size

Synthesis method

Evolution

Cocatalyst Reactant solution

Catalyst mass (mg)/reactant solution volume (mL)

CAgIZS NCs

100 nm

coprecipitation

methanol

RuO2

50/−

Xenon; λ > 118.5 400 nm (1000 W)

CuS NWA

w (50 nm) electrochemical anodization

methanol

catalyst (film 2 × 2 cm2)

sunlight (AM 1.5)

NaHCO3 (0.21 g, 2.5 mmol), Na2S (2.5 mmol/L) aqueous solution water droplets (5 μL), 5.5 atm. 1% CO2

Light source


38 μmol/m2·h

Ref 1846

1843

several energy conversion technologies. In particular, highperformance electrocatalysts for ORR and OER are an essential component in rechargeable metal−air batteries and regenerative low-temperature fuel cells.833,1856−1864 In these devices, the use of a catalyst based on scarce and high-cost platinum group metals represents a severe cost constraint that strongly limits deployment. ORR and OER activity depend, among other factors, on O2 absorption and activation. Multivalence off-stoichiometric compounds with high defect densities have been shown to be excellent candidates for these applications. While mostly transition metal oxides have been used due to their low cost and good stabilities in alkaline media, some chalcogenides have also been recently proposed. Among them, Cu chalcogenides show outstanding properties. As an example, Liu et al. demonstrated the exceptional electrocatalytic ORR properties in alkaline media of Cu2Se nanowires, with performance above that of Pd.833 Notably, Seredych et al. used CuxSy-nanoporous carbon composites for ORR in alkaline medium and demonstrated this material to outperform Pt in terms of current density and stability.1856 A lot of research has focused on optimization of the catalytic activity by modifying the electrodes. Wang et al. used Cu-deficient plasmonic Cu2−xS nanoplates as electrocatalysts for ORR.563 They demonstrated that supporting the nanoplates on carbon black and reduced graphene oxide increased activity.563 They also tested different doping levels to probe whether an increase of the concentration of free holes increases ORR activity, with the best performances reached using off-stoichiometric compounds. Cu2S nanoplates can also be used for OER (Figure 73),1864 which is the bottleneck in the water splitting process. This is because the transfer rate of the four electrons involved is slow and the activation energy barrier for O−O bond formation is high, requiring overpotential in substantial excess of its thermodynamic potential. Ir- and Ru-based catalysts are generally used for this reaction. Alternative Cu2+ complexes have also been tested, which has motivated the exploration of Cu2S NCs. In this regard, An et al. observed that Cu2S NCs outperformed CuS NCs and that the addition of glycine promoted performance, reaching activities and durabilities close to those of state-of-the-art commercial Ir/C catalysts.1864 9.2.2. Other Catalytic Reactions. Cu chalcogenides have also been used for synthetic organic chemistry, demonstrating excellent activities and selectivities. In this direction, colloidal NCs compete with molecular catalysts, with the main advantage being that all or most catalytic atoms are exposed. In NCs, part of the atoms are trapped within the NC and thus do not directly participate in the reaction. The advantage of NCs is the ability to easily separate them from the product simply by size/mass or magnetism. Additionally, NCs offer the opportunity for multiple functionality including light absorbance

Figure 72. Scheme of the light induced cation exchange process used to produce Cu2S-Pt nanorods and of the photocatalytic CO2 reduction into CO and CH4, and experimental results of the CO evolution from CdS-Cu2S NRs after illumination with a 447 nm laser or a broad spectrum Xe lamp. Reproduced with permission from ref 1842. Copyright 2015 American Chemical Society.

Figure 73. Proposed reaction mechanism for Cu2S NCs in water oxidation. Reproduced with permission from ref 1864. Copyright 2015 American Chemical Society.

9.2.1. Oxygen Evolution and Oxygen Reduction Reactions (OER/ORR). OER and ORR are particularly important reactions. In this direction, the development of highly active, low-overpotential, stable, and low-cost electrocatalysts based on abundant elements is a critical challenge in 6021


Chemical Reviews

Review

Table 28. Other Catalytic Reactions of Cu Chalcogenide NCs Cu-based NC

Particle size

Synthesis method

Active material

Catalytic activity

Ref

Cu2Se NWA

d × l (50−70 nm × several μm)

solid-liquid phase chemical transformation

Cu2Se

ORR

833

Cu-BTC(MOF)/ GrO rGO/Cu2−xS

ORR

1856

ORR

563

GC/CuxS

OER

1864

OER

1771 380 1876 1841

CuxSy/(nanoporous C) composite Cu2−xS NCs CuxS NCs

d × t (21 nm × 2.5− 3.5 nm) d × t (11.0 nm × 3.2 nm)

Cu2Se film

colloidal liquid state transformation process electrochemical deposition

CCoTS NCs Pt-CuS NHS CIS NCs

d (150−250 nm) 5−15 nm d (3.8 nm)

electrospinning colloidal colloidal

Cu2Se/CdSe/ TiO2 GC/CCoTS Pt-CuS/C TiO2−CIS

CIS film

average grain size (10−15 nm) d × l (60−70 nm × 1.5 μm)

sol-gel

G/ITO/CIS

HER MOR CO2 reduction to methane, ethane, acetaldehyde, paraffins CO2 reduction to CO

CBD, accompanied by template-assisted method

Zn:CIS

CO2 hydrogenation to methanol

1821

Cu−S cluster CuS NMP

d (3.5 μm)

hydrothermal

N2O reduction to N2 reduction of trichloroacetic acid

1850 1855

Cu2−xSe NCs

d (65.5 nm)

water-based PVP-assisted

reduction of 4-NP to 4-AP

1879

CuS NCs CuS NMP

18.29 nm d (4.0−4.3 μm)

cellulose pyrolysis thermal decomposition of AP

1874 1875

CuS NCs CuS NCs Cu2Se NCs Cu7S4 NCs

2−3 nm 2−3 nm 12−14 nm 14 nm

wet chemical elemental-direct-reaction route microwave irradiation microwave colloidal solvothermal

Cu−S CILE/Gr-CuS/ Hb/CTS Cu2−xSe/rGO/ PVP CuS CuS

d × t (23 nm × 2 nm) d (10 nm) od × l (150−200 nm × 1 μm) d (50−100 nm)

colloidal

synthesis of dihydropyrimidinones synthesis of xanthenes Suzuki and Sonogashira cross coupling reactions Suzuki coupling reaction, hydrogenation of nitrobenzene, oxidation of benzyl alcohol hydrogenation of phenylacetylene

1870 1871 1872 1849

CuS NCs

CuS CuS Cu2Se Cu7S4@M (M = Pd, Ag, Au) C/CuS-Pd

colloidal hydro-/solvothermal

CuS CuS

oxidation of hydrosulfide ions MB degradation

1877 759

coprecipitation

CuS

ultrasonic assisted removal of chrysoidine G

1779

Zn:CIS NA

CuS NCs CuS NTA CuS NCs

1844

1873

ration (1−4%) reduces the initial and final decomposition temperature by 39 and 108 °C, respectively, helping the modulation of the ammonium perchlorate burning behavior. Also notable is the work by Ding et al., where CuS-Pt heteronanostructures provided selective electrocatalytic activity toward methanol oxidation, a key reaction in direct methanol fuel cells.1876 Morever, the degradation of dyes and pollutants with a clear optical footprint has been also targeted using Cu chalcogenides, both for its intrinsic interest and as an easy to follow paradigmatic reaction to test catalytic properties of newly developed materials. In this regard, CuS NCs were used for the catalytic oxidation of hydrosulfide ions,1877 CuS and Cu1.8S nanostructures were employed for the degradation of methylene blue,759,799,1878 CuS NCs on activated carbon were used for the degradation of chrysoidine G,1779 and Cu2−xSe NCs embedded with RGO on nanofibers were tested for the reduction of 4-nitrophenol.1879 An overview of the other catalytic reactions of Cu chalcogenides is provided in Table 28.

for photocatalysis, multiple consecutive reactions at particular reaction sites, magnetic functionality for separation, or induction heating or plasmonic enhancement.1865−1869 As an example, Chaudhary et al. demonstrated that CuS NCs exhibited good catalytic activity for the Biginelli reaction with an excellent yield in the production of dihydropyrimidinones.1870 The same group reported the synthesis of xanthene derivatives using CuS NCs.1871 Singh et al. used tetragonal Cu2Se nanoflakes for the Suzuki and Sonogashira cross coupling reactions.1872 They proposed the excellent activities observed to be related to the formation of surface CuBr in the presence of TBAB, which activates both coupling reactions. In a separate study, Wang et al. demonstrated the selective hydrogenation of phenylacetylene over CuS nanoplates containing highly dispersed Pd.1873 In the energy conversion area, Chaki et al. used CuS NCs for cellulose pyrolysis, which is the key step in the conversion of biomass into liquid or gaseous biofuels, with obvious advantages in terms of transportation, storage, and ready usage.1874 The observation was that pure cellulose decomposition starts at 205 °C in the presence of 3% CuS NCs instead of at 295 °C in their absence. The catalytic activity of CuS nanostructured particles in the thermal decomposition of ammonium perchlorate was further reported Zeng et al., and this is significant, as it is the most common energetic material and oxidizer in composite solid propellants.1875 CuS incorpo-

10. ENERGY STORAGE APPLICATIONS 10.1. Batteries

A battery is composed of several electrochemical cells and is able to convert stored chemical energy into electrical energy. Each cell consists of a cathode and an anode, which are separated by an electrolyte solution. The undergoing chemical 6022


Chemical Reviews

Review

(Figure 74).1888 In this study, a two-step reaction was observed which supported an initial formation of intercalation compounds, such as LixCuS (0 < x ≤ 0.16) and LiyCuS (y ≈ 1), in the first stage of the electrochemical reaction. A conversion reaction of the LixCuS insertion phase was observed in the second process, as indicated by the presence of Li2S and Cu metal in significant quantities. In this reaction scheme, the high electronic conductivity of the sulfide phases1889 allows the LiyCuS/Cu interface to occur anywhere in the particle, and not just exclusively near particles that are electronically connected (i.e. via the carbon) to the current collector. In addition, the rapid diffusion of Li+ in Li2S allows for the formation of a thick Li2S film and these kinetics permit the diffusion of Li+ through the Li2S to the Li2S/LiyCuS interface. As a result, Cu particles or dendrites will grow away from the LiyCuS particles (Figure 74c), whereas the Li2S phase will grow into the LiyCuS phase. Another possibility is that Li+ transport through LiyCuS can also occur, which could cause the Li2S interface to grow directly on the Cu particle. This may prevent further growth of the Cu particle (by blocking the Cu+ diffusion) and encourage growth of other Cu nuclei on the LiyCuS surface or in the particle, with the latter resulting in particle cracking. Furthermore, the transformations that occur in this system are complicated by the decomposition of the unstable LiyCuS phase to form Cu1.96S and Li2S, a process mediated by the diffusion of both the Cu1+ and Li+ ions.1886 It is not clear why the Cu1.96S is not reduced further, which suggests that the decomposition process results in domains of Cu 1.96 S surrounded by Li2S, so that electrical contact is lost with a significant fraction of this phase. The myriad of binary compounds with different stoichiometries (i.e. Cu2S, Cu7S4, Cu1.96S, Cu1.85S, and CuS), and their complicated interaction with Li+ ions, inspired Wang et al. to study the electrochemical behavior of electrodes, that were prepared with nanomaterials with different Cu:S molar ratios.1890 In this work, the authors concluded that stoichiometries with higher Cu content demonstrated better cycling and rate performance capabilities, due to their higher conductivities and the unique displacement reaction between Cu2S and Li2S. This displacement reaction is enabled by the high ionic mobility of Cu ions and the similarity of their crystal structures, which allows Li to be inserted and removed reversibly during the charge−discharge process.1891 McDowell et al. presented an atomic-level view of the structural and morphological transformations of Cu2S NCs during their reaction with Li.1892 Using in situ HR-TEM imaging, the authors observed the displacement reaction mechanism in Cu2S NCs, in which Cu metal was extruded out of NCs during the formation of a single Li2S domain. The new Li2S particle retained the original shape and morphology of the Cu2S particle and the extruded Cu formed either particles at the surface, or larger dendrites if many Cu2S NCs were in contact. In addition, a rearrangement of the sulfur sublatttice in Cu2S after its reaction with Li was observed, in that the close-packed sulfur planes shifted from hcp-type stacking (Cu2S) to fcc-type stacking (Li2S). However, copper sulfides suffer from poor capacity retention, which is their main disadvantage. The poor electrochemical performance of the Li/Cu/S system is attributed to (i) large volume changes (loss of structural stability); (ii) decomposition of the electrolyte; and (iii) dissolution of sulfur species in the electrolyte.1893 At voltages higher than 1 V, reversible Li+ intercalation/deintercalation processes occur. However, the

reaction in the electrodes liberates electrons, which generates a current when they are connected externally. The amount of electrical energy produced by the battery is a function of the capacity (mA h g−1) and the potential of the cell (V), parameters which are intrinsically related to the electrode materials.1880,1881 Among the different battery technologies, lithium ion batteries (LIBs) are the most widespread technology for energy storage, with applications from portable electronics to power sources for automotive applications. There are still many remaining challenges to be overcome, particularly in terms of improving the performance of LIBs, such as the requirement for cheaper and safer electrodes that have high reversible capacity. During the last 15 years, the search of highperformance electrode materials has been intensely focused on nanostructures.1882,1883 Nanostructures can offer improved energy storage capacity and cyclic stability due to (i) the huge surface area and large number of reaction sites (better electrolyte immersion), (ii) the short distance for mass and charge diffusion (reduced diffusion path length of Li+ ions), and (iii) the mitigation of the volumetric changes during electrode operation (structural stability).1884 In the search for high-performance electrode materials, Cu chalcogenide NCs have been identified as promising candidates, owing to their high theoretical capacities and electrical conductivities. An overview of the different Cu chalcogenide NCs which have been used as electrode materials for LIBs is discussed in the following section, and details of the relevants parameters of these cells are presented in Table 29. 10.1.1. Cu−S. Binary Cu-chalcogenides exist in a wide range of stoichiometric compositions and crystal structures, which are tightly related with their electrochemical performance. CuS is a promising cathode material for use in LIBs, due to its high specific capacity of 561 mA h/g, good electronic conductivity of 1 × 10−3 S cm−1, flat discharge curves, and the use of low-cost, abundant materials.1885 According to the electrochemical curve, the discharge profile of CuS presents two distinct voltage regions, characterized by two plateaus, at 2.05 eV and 1.68 eV, which have been attributed to the following reactions:1886,1887 CuS + x Li+ + x e− → LixCuS (first process) Cu yS + 2Li+ + 2e− → Li 2S + yCu (where y ≈ 2) (second process)

The first process involves an Li insertion reaction to form LixCuS, which further reacts with more Li to generate Li2S and Cu metal. In addition, an intermediate phase of Cu2S has also been proposed, whereby LixCuS in the first process reacts with more Li to generate a Cu2S or Cu1.96S phase, with slight varations in the composition of this intermediate phase being reported in the literature: 2LixCuS + 2Li+ + 2e− → 2Li 2S + Cu 2S (intermediate phase)1887 1.96LixCuS + (2 − 1.96x)Li+ + (2 − 1.96x)e− → Li 2S + Cu1.96S (intermediate phase)1886

Yamakawa et al. presented a detailed study to understand the reaction mechanisms during the interaction of Li with CuS 6023


Chemical Reviews

Review a

Table 29. Overview of Cu chalcogenide NCs That Have Been Used in Energy Storage Applications Active material

Working Electrode*

Cu2S nanowire array Cu2S coral-like

E-D on Cu foil (70:15:15)

Cu2S/tubular mesoporous carbon composite

(80:10:10)

Cu2S/carbon composites

(80:10:10)

Cu2S

(80:10:10)

Cu2S

(70:20:10)

Cu2S

Grown on Cu foil (70:20:10)

Cu2S Cu2S Cu2S

Cu foam sulfurization (75:15:10)

Cu2−xS

E-D on Cu foil

Cu2S CuS thin film flake morphology carbon coated CuS

(60:20:20)

(80:10:10)

CuS CuS CuS nanowire bundles

E-D on Cu foil (80:10:10) (75:15:10)

tubular dandelion-like CuS structures CuS rods

(87:?:?) (70:20:10)

hollow CuS

(70:15:15)

CuS/Graphene “Double-Sandwich-Like” CuS@RGO CuS NWs/RGO

(70:20:10) (70:20:10)

CuS nanoplate-based sphere like hierarchical structures CuS nanorods (below 10 nm) Ultrathin CuS nanosheets CuxS/Cu Nanotubes

Cu2−xSe nanosheets Cu2Se

(75:10:15)

Discharge Capacity - mA h g−1 (current density - mA g−1)

Electrolyte Cu−S 1M LiPF6 in DEC:EC (1:2 vol.) 1 M LiTFSI in DOL:DME (1:1 vol.) 1 M LiTFSI in DOL:DME (1:1:1 vol.) 1 M LiPF6 in EC:DEC:EMC (1:1:1 vol.) 1 M LiPF6 in EC:DEC:EMC (1:1:1 vol.) 1 M LiCF3SO3 in DME:DOL (1:1, vol.) 1 M LiPF6 in EC:DMC:DEC (1:1:1 vol.) + vinylene carbonate 1M LiTFSI in DOL:DME (1:1:1 vol.) 1 M LiPF6 in EC:DMC:DEC (1:1:1 vol.) 1 M LiTFSI in DOL:DME (1:1, vol.) 1 M LiTFSI in DOL:DME (1:1, vol.) 1M NaCF3SO3 in TEGDME 1 M LiTFSI in tetramethylene sulfone 5 M LiNO3 and 0.001 M LiOH. (Aqueous) 1 M LiClO4 DOL:DME (1:2, vol.) 1 M LiPF6 in EC:DMC (1:1, vol.) 1 M LiPF6 in EC:DEC (1:1 w/w) with 30% FEC 1 M LiPF6 in EC:DMC:DEC (1:1:1 vol.) 1 M LiTFSI in DOL:DME (1:1, vol.) 1M LiPF6 in EC:DEC (1:1, vol.) 1 M LiPF6 in EC:DMC (1:1, vol.) 1M LiPF6 in EC:DMC:DEC (1:1:1 vol.) + VC (2%) 1 M LiPF6 in EC:DMC:DEC (1:1:1 vol.)

(80:15:5)

1 M LiPF6 in EC/DMC (1:1, vol.)

(70:20:10) (80:10:10) 70:20:10

1 M LiPF6 in EC:DMC (1:1, vol.) 1M LiPF6 in EC:DEC (1:1, vol.) 1 M LiPF6 in EC:DMC (1:1, in wt %) Cu-Se 1 M LiPF6 in EC:DMC (1:1, vol.) 1 M NaClO4 in EC:DMC (1:1, vol.) 1 M NaClO4 in EC:DMC (1:1, vol.) 1 M NaClO4 in EC:DMC (1:1, vol.) 1 M NaClO4 in EC:DMC (1:1, vol.) Cu-Te 1 M LiPF6 in EC:DEC (1:1, vol.) 1 M LiPF6 in EC:DEC (1:1, vol.) 1 M LiPF6 in EC:DEC (1:1, vol.) 1 M LiPF6 in EC:DEC (1:1, vol.)

Cu2Se

(80:10:10) Selenization of Cu grid PLD on SS

CuSe

PLD on SS

CuSe2

PLD on SS

Cu2−xTe nanosheets Cu2−xTe nanosheets Hollow Cu2−xTe NPs Cu1.67Te nanocubes

(70:20:10) (70:20:10) (70:20:10) (70:20:10)

6024

Cycles

Electro-chemical Window (V vs Li+/ Li)

Ref

230 (2C) 280 (1C)

100 100

0.01−2.5 1−3

1908 1890

270 (0.2C) 225 (10C) 300 (100)

300

1.0−3.0

1103

100

0.01−3

1939

30 (100)

100

0.01−3

1939

250 (C/10, C/5, C/2)

50

1.2−2.5

1896

340*** (0.1 mA cm− 2)

500

0.02−3

1940

313 (100)

100

1−3

1891

0.41*** mAh cm−2 (0.25C)

100

0.02−3

1941

200 (1C)

150

0.8−3

1895

200 (C/2)

50

1.2−2.6

1911

220 (50) 190

20 20

0.4-2.6** 1.8−2.5

1918 1881

45 (10)

100

1.8−2.6

1905

376 (2C) 90 (10 mAh/g) 518 (0.2C) 196 (4C) 425 (1120)

100 30 90

1.2−3.0 1.8−2.6 0.02−2.7

1347 1905 1942

2nd

1−3

1909

472 (100)

100

1−3

1891

75 (5)

Fading 40 25 100

1.5−2.6

1902

0.01−3 0.001−3.0

1915 1917

620 (0.5C) 320 (4C) 1C = 560 mAh g−1 ∼90 (100)

100

0.02−3.00

1943

430 10 fading

0.01−3.00

1904

∼400 (0.1C) 642 (200) 282 (5)

250 360 50

0.01−3.0 0.005−3 0.001−3.0

1910 70 1944

210 (50) 113.6 (0.1C)

100 100

0.01−3 1.8−2.5**

828 1945

177.7

100

1−2.5**

1919

118.9

100

1.8−2.5**

1919

196.6

100

1.8−2.5**

1919

100 220 180 120

5000 130 130 130

0−2 0−2 0−2 0−2

144 144 144 144

296 (50) 748.6 (0.2C)

(400) (100) (100) (100)


Chemical Reviews

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Table 29. continued Active material

Working Electrode*

CTS NPs mesoporous CTS spheres

(90:0:10) (80:10:10)

CTS microspheres/RGO

(80:10:10)

CTS microspheres

(80:10:10)

CTS cabbage like

(80:10:10)

CTS NPs/RGO

(80:10:10)

CTS NPs

(70:20:10)

(H3O)2(enH2)Cu8Sn3S12 Hollow CTS microspheres

(70:20:10) (80:10:10)

CZTS flower-like CZTS NPs Porous CZTS film

(80:10:10) (80:10:10) deposited on Mo sheet (80:10:10)

CZTS NPs CZTS film 3D CZTS microstructure Cu2ZnSnS4 nanocrystals

RFMS Ni foam (80:10:10) (80:10:10)

CuSbS2 bricks

(80:10:10)

CuSbS2 Blocks

(80:10:10)

Flower like Cu3BiS3

Cu2MoS4 Cu2MoS4/GO

(6:2:2) (6:2:2)

CIZS@graphene

(80:10:10)

CuFeS2 spike-like nanorods

(70:15:15)

CuFeS2 spike-like nanorods

(70:15:15)

CuFeS2 hexagonal plate

(40:20:40)

Discharge Capacity - mA h g−1 (current density - mA g−1)

Electrolyte

Cu-Sn-S 1 M LiPF6 in EC:DEC (1:1, vol.) 431 (60) 1 M LiPF6 in EC:DMC:DEC 436 (100) (1:1:1, in wt %) 425.6 (100) 1 M LiPF6 in EC:EMC:DMC (1:1:1 vol.) 1 M LiPF6 in EC:EMC:DMC 99.8 (100) (1:1:1 vol.) 621 (100) 1 M LiPF6 in EC:DMC:DEC (1:1:1, in wt %) 1 M LiPF6 in in EC:DMC (1:1, 560 (100) vol.) 436 (100) 1 M LiPF6 in EC:DMC:DEC (1:1:1, in wt %) 1 M LiPF6 in EC:DMC (1:1, vol.) 563 (100) 1 M LiPF6 in EC:EMC:DMC 190 (100) (1:1:1; vol.) Cu-Zn-Sn-S 1 M LiPF6 in EC:DMC (1:1, vol.) 150 (200 ) 1 M LiPF6 in EC:DMC (1:1, vol.) 100 (15 ) 1 M LiPF6 in in EC:DMC (1:1, in 680 (100) wt %) 1 M LiPF6 in in EC:DMC (1:1, in 288 (100) wt %) 1 M LiPF6 in EC:EMC:DMC 668 (100) (1:1:1 vol.) 1 M LiPF6 in EC:DEC (1:1, vol.) 786 (100) 1 M LiPF6 in EC:DMC (1:1, vol.) 100 (15) Cu-Sb-S 1 M LiPF6 in EC:EMC:DMC 85.7 (1:1:1 vol.) 1 M LiPF6 in EC:EMC:DMC 225 (100) (1:1:1 vol.) Cu-Bi-S 1 M LiPF6 in EC:DMC: DEC 180 (100) (1:1:1 vol.) Cu-Mo-S 1 M LiPF6 in EC:DEC (1:1, vol.) 100 (20) 1 M LiPF6 in EC:DEC (1:1, vol.) 150 (20) Cu-In-Zn-S 1 M LiPF6 in EC:DEC:DMC 494 (2000) (1:1:1, in wt %) Cu-Fe-S 1 M LiPF6 in EC:EMC:DMC 63.4 (0.2C) (1:1:1 vol.) 1 M LiTFSI in DOL:DME (1:1, 425.3 (0.2C) vol.) 1 M LiPF6 in EC:EMC:DMC 1100 (14) (2:1:2 vol.)

Cycles

Electro-chemical Window (V vs Li+/ Li)

Ref

15 50

0.1−3 0−2.5

1921 689

100

0.01−3

1923

100

0.01−3

1923

50

0−2.5

937

100

0−2.5

937

50

0.01−3

1924

100 50

0.01−3.0 0.01−3

1492 1922

50 30 40

0−3 0.8−3 0−4

1929 1946 1927

30

0−3

342

100

0.01−3

1928

100 30

0.01−3 3.0−0.8

1930 1946

50

0.01−3

1378

15

0.01−3

1935

50

0−2

313

10 10

1.5−2.8 1−3

322 322

480

0.05−3

1932

50

1−3

1938

50

1−3

1938

1

1.0−2.8

1876

a

Notes: The working electrode column is intended to describe the deposition method of the active material onto the substrate, and in the case of a prior slurry preparation to paint the electrodes, the values in brackets provide its formulation in terms of the ratio between constituents (active material:conductive carbon:binder). Electrochemical windows marked with ** refer to V vs Na+/Na.

On the other hand, the dissolution of sulfur species in the electrolyte can be minimized by replacing carbonate-based electrolytes with ether-based 1895,1896 or solid electrolytes.1897,1898 For instance, a battery using a Cu2S electrode exhibited a good cycling performance, with capacity retained at about 200 mA h g−1 after 150 cycles and coulombic efficiency of 98.4% in 1,3-dioxolane (DOL)/1,2-dimethoxyethane (DME) (1 M LiTFSI) electrolyte, whereas its capacities rapidly dropped to 0 mA h g−1 after five cycles in ethylene carbonate/dimethyl carbonate (1 M LiPF6) electrolyte.1895 To understand the poor cycling stability of carbonate-based

lithiated phases undergo full conversion reactions when cycled at lower voltages (below 1 V).1887 The formation of lithium polysulfides causes severe capacity fading upon cycling, due to their dissolution in the electrolyte and the consequent drift away from the electrode.1886 Additionally, structural integrity is lost and the electrical pathways within the electrodes become isolated, enlarging capacity fading. Several approaches have been investigated, in order to avoid this irreversible reaction, by limiting discharge voltage to 1 or 1.5 V to try to improve cycling.1886,1894 Nonetheless, the capacity is highly reduced from the theoretical value when the discharge voltage is cut-off. 6025


Chemical Reviews

Review

which is higher than the theoretical capacity, and this increase was attributed to a possible surface activation process in the ultrathin CuS nanosheets. The Cu2S nanowire arrays reported had higher discharge/charge capacity and reversible lithium capacity, compared to Cu2S thin films, as the nanostructures improved the reversible reaction of Cu2S with Li+, due to their better flexibility for accommodating the strain of Li+ insertion/ extraction, and shorter path lengths for both electronic and ionic transport.1908 An alternative route to improve the capacity fading is to fabricate composites with a carbon matrix, which can act as an electrically connecting media and as a buffer layer for volume changes, which occur during the lithium insertion and extraction process. Recent studies in Li−S systems have shown that porous carbon-decorated sulfur can effectively alleviate the escape of polysulfite from the cathode and can improve cycle stability.1912,1913 This idea was applied to the Cu2S−Li system, in which Cu2S-tubular mesoporous carbon composites were tested as cathode materials for LIBs.1103 In this system, the formation of Cu2S occurred as an intermediate during the initial charge−discharge cycles of a Cu foil with pristine sulfur impregnated tubular mesoporous carbon (Figure 75). This configuration prevented the aggregation of NCs and maintained the structural integrity of active materials, which enhanced the electronic conductivity.1104,1914 Many examples of carbon-based nanocomposites with Cu chalcogenides can be found in the literature.1915 As an example, CuS/graphene nanocomposites exhibited improved cycling stability and capacity retention, compared to bare CuS, in which the enhanced electrochemical performance was attributed to the good electronically connecting channels offered by graphene.1916 Ren et al. reported capacities of ∼750 mA h g−1 for the CuS-RGO double sandwich-like nanocomposites.1917 The improved electrochemical performance was attributed to the outstanding conductivity of RGO nanosheets, as well as the shorter Li+ ion diffusion in the double-sandwich-like CuSgraphene nanocomposite. Cu chalcogenide compounds have also been briefly explored as anode materials for sodium ion batteries (SIBs). In this regard, Kim et al. studied the electrochemical performance of Cu2S as an anode for SIBs and obtained a capacity of 220 mA h g−1 after 20 cycles and coulombic efficiencies close to 100%.1918 In addition, analysis of the structure and composition explained the discharge process, with the following reaction scheme: Cu2S + xNa → NaxCu2S (x < 2). 10.1.2. Cu−Se. Xue et al. prepared copper selenides with different stoichiometries (CuSe2, CuSe, and Cu2Se) and evaluated their electrochemical performance.1919 In this work, it was revealed that different electrochemical reaction

Figure 74. Schematics illustrating possible reaction pathways for CuS (the black box represents the current collector). (a, b) Cu1+ diffusion (brown arrows), away from the LiCuS/Li2S interface to the LiCuS/Cu interface occurs. Li+ diffusion (black arrows) to the LiCuS/Li2S interface occurring both through the Li2S and directly from Li+ the electrolyte. (c) Cu nuclei formation away from the current collector. Adapted with permission from ref 1888. Copyright 2009 American Chemical Society.

electrolytes, Shi et al. studied the electrochemical performance of the Cu2S electrode in the presence of carbonate electrolytes, with different molecular structures.1899 The authors noted significant differences between linear- (dimethyl or ethyl methyl) and cyclic- (propylene, vinylene, and fluoroethylene) carbonates, which were found to be linked to their polarity and, hence, the ability to interact with polysulfides. In terms of the discharge capacity, it was reported to be 200 mA h g−1 for the linear-carbonates after 50 cycles, whereas it faded completely (0 mA h g−1) for the cyclic-carbonates after 10 cycles.1899 Wang et al. reported the in situ preparation of a CuS cathode and used liquid electrolyte mixtures, containing 1 M LiClO4 and a solvent mixture of DME/DOL.1900 The good cycling behavior and high capacity retention of their cells were attributed to the large 3D net structure and the production of mixed ion electron conductors during cycling.1901 It was also speculated that mixed conductors could not only increase the conductivity, but also prevent the dissolution of the products formed during the charge−discharge process.1347 One of the most explored strategies to overcome the poor electrochemical performance is the use of well-designed nanostructured morphologies.1883,1884 Numerous reports have studied the electrochemical performance of different Cu2−xS nanostructured morphologies, such as hollow spheres,1902 porous spheres,1903 hierarchical sphere structures,1881,1904,1905 stick-like hierarchical structures, 1906 nanowires, 1907,1908 flakes,1347,1881 tubular dandelion-like structures,1909 nanorods,1910 and polyhedrons.1911 For example, ultrathin CuS nanosheets presented a discharge capacity of 506 mA h g−1 and a corresponding charge capacity of 378 mA h g−1 in the 1st cycle, and this decreased to 321 mA h g−1 in the 20th cycle.70 In addition, it was noted that the capacity increased to 642 mA h g−1 at the 270th cycle at a current rate of 200 mA g−1,

Figure 75. Electrode with firmly anchored polysulfide ions, which are held by the copper ions. Note: CuxS is an intermediate, not the final product. Reproduced with permission from ref 1103. Copyright 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim. 6026


Chemical Reviews

Review

the high reversible capacity and excellent stability were associated with the cabbage-like morphology.937 10.1.5. Cu-Zn-Sn-S. CZTS has also been considered as an electrode material for LIBs, in which both Sn and Zn are electrochemically active elements toward Li.1926 Based on previous studies of the lithiation of related binary compounds, Yin et al. proposed the following reaction scheme:1927

mechanisms were observed for each phase: CuSe2: 5CuSe2 + 14Li+ + 14e− ⇌ Cu 2Se + Cu3Se2 + 7Li 2Se5

CuSe: 5CuSe + 4Li+ + 4e− ⇌ Cu 2Se + Cu3Se2 + 2Li 2Se

Cu 2ZnSnS4 + 8Li+ + 8e− ↔ 2Cu + Zn + Sn + 4Li 2S

Cu2Se:

Sn + x Li+ + x e− ↔ LixSn (0 ≤ x ≤ 4.4)

Cu 2Se + 2Li+ + 2e− ⇌ Li 2Se + 2Cu

It was noted that the same discharge products (Cu2Se and Cu3Se2) were observed for the CuSe2 and CuSe phases, and that these compounds exhibited similar electrochemical behavior despite their differences in stoichiometry and crystal structure. In addition, the authors reported significant capacity losses which were correlated with large structural and volumetric changes that were induced by the conversion reactions. Similar to Cu2S, Cu2Se compounds undergo a displacement reaction, in which Li is gradually inserted into the face-centered cubic Se matrix, occupying the same positions as copper and forming Li2Se. This Li2Se structure has very similar lattice parameters to Cu2Se, and the volumetric changes are almost negligible, thus allowing for good capacity retention during cycling. 10.1.3. Cu−Te. Among the binary Cu chalcogenides, there are very few reports of Cu2−xTe nanostructures, which have been tested as anodes for LiBs. In this regard, Han et al. prepared nanocubes, nanosheets, and hollow nanoparticles and compared their electrochemical performance.144 The highest capacity and the best cycling performance were obtained for Cu2−xTe nanosheets, which had a capacity of 220 mA g−1 after 120 cycles, compared to a theoretical capacity of 226.86 mA g−1 and the following reaction mechanism: Cu2−xTe + 2Li+ + 2e− → (2−x)Cu + Li2Te. Another merit of Cu2−xTe is its high density and discharge voltage, which could lead to a high volumetric capacity. 10.1.4. Cu−Sn−S. Tin-based nanomaterials have been extensively used as anodes for LIBs due to their high theoretical performance.1920 Wu et al. reported the use of Cu2SnS3 polycrystalline nanoparticles for the first time as electrode for LIBs.1921 In this report, the authors considered this material as an excellent candidate, due to its large interlayer distance (2.280 Å) and tunnel size (3.921 × 5.587 × 4.210 Å3) in the crystal structure, which is considerably larger than the diameter of the Li ions (1.36 Å) and could be effective in facilitating the diffusion of Li ions through the crystal structure. Moreover, Sn is also an active element toward Li alloying. A variety of Cu2SnS3 nanostructures were subsequently investigated, including mesoporous spheres,689,1922 cabbage-like nanostructures,937 hollow microsphere,1923 small nanoparticles,1924 and a crystalline Cu−Sn−S framework.1925 Based on the electrochemical measurements and structural analysis, the Li storage mechanism was described by the following equations:

Zn + x Li+ + x e− ↔ Li xZn (0 ≤ x ≤ 1.5)

The electrochemical performance of hierarchical porous films, 1927 flower-like nanostructures, 1928,1929 and nanocrystals342 has been investigated. Several strategies have been proposed to improve stability and cyclability. In particular, Lin et al. prepared porous CZTS thin films on a Ni foam and coated them with lithium phosphorous oxynitride, in which this configuration allowed for a release of the volumetric expansion and reduced the dissolution of sulfides.1930 A gradual capacity fading was observed, despite the high capacities obtained after 100 cyles (668 mA h g−1). Shortly after this report, Jiang et al. demonstrated the importance of material morphology on the electrochemical performance by using 3D CZTS microflowers, in which capacities up to 786 mA h g−1 were obtained after 100 cycles at reasonably high currents, 100 mAg−1.1928 10.1.6. Cu-In-Zn-S. This system was investigated based on the alloying ability of In and Zn with Li, and the increase of electrical conductivity associated with the presence of Cu.1931 In order to further improve its electrochemical performace, Tang et al. prepared CuInZnS-decorated graphene nanosheets, which had very good capacity (∼500 mA h g−1) and stability (no fading after 480 cycles) at high current rate (2000 mA g−1).1932 10.1.7. Cu-Bi-S. Zeng et al. reported the electrochemical performance of Cu3BiS3 as an electrode for LIBs for the first time.313 The authors defined the whole reaction process as: Cu3BiS3 + 9Li ⇌ 3Li 2S + Li3Bi + 3Cu

However, the fast capacity decay suggested the occurrence of side-reactions (with the electrolyte) and irreversible reactions, where the capacity after the 50th cycle was reported to be about 180 mA h g−1. 10.1.8. Cu-Sb-S. CuSbS2 has also been considered as a potential electrode for LIBs. Theoretically, this material provides the same advantages as binary copper sulfides, with the additional presence of Sb, which is electrochemically active toward Li.1933 However, there were only two similar reports evident in the literature, in which the authors produced nanobricks1934 or nanoblocks1935 and tested their electrochemical performance. While the initial capacity was high (10901934−8771935 mA h g−1), the reversibility of the reactions involved was not achieved, as large capacity fading was observed. 10.1.9. Cu-Mo-S. Yu et al. proposed the use of layered Cu2MoS4, which was inspired by the well-known Li intercalation in Mo cluster sulfides.322 In this report, the authors carried out electrochemical and structural characterization and described the reaction mechanism as:

Cu 2SnS3 + 6Li+ + 6e− ↔ 2Cu + Sn + 3Li 2S

Sn + x Li+ + x e− ↔ LixSn (0 ≤ x ≤ 4.4)

However, the reversibility of these reactions was not complete and full regeneration could not be achieved. On the contrary, Qu et al. achieved a capacity of 621 mA h g−1 after 50 cycles for Cu2SnS3 cabbage-like nanostructures, in which

Cu 2MoS4 + 2Li+ + 2e− → Li 2Cu 2MoS4 6027


Chemical Reviews

Review

Table 30. Oveview of Cu Chalcogenide NCs Which Have Been Used as Electrode Materials for Supercapacitors Active material

Notes on Working Electrodea

Electrolyte

Specific capacitance (F g−1) (current density (A g−1))

CuS NTs CuS NTs

(85:10:5) Ni foam 0.5 M NaOH in water (85:10:5) Ni foam 0.5 M NaOH/0.5 M Na2S/0.5 M S powder in water CuS@CNT (80:10:10) Ni 2 M KOH in water foam CuS NWs Cu foil 3 M KCl in water CuS (75:15:10) Ni 6 M KOH in water substrated CuS-MWCNTs (75:15:10) Ni 6 M KOH in water substrated CuS nanoplates FTO 1 M LiClO4 in water CuS nanosheets (75:15:10) Ni 6 M KOH in water foam CuS/Acetylene (75:15:10) Ni 6 M KOH in water black foam CuS Cu foil 1 M KCl in water Flower-like CuS (unknown) 2 M KOH in glycerol CuS NCs (75:20:5) Stainless 1 M Na2SO4 in water steel

Cycles (Capacity retention %)

Electrochemical Window (V vs)

Ref 1485 1485

500 (15) 2175 (15)

1000 (80)

−1.3−0.1 (Hg/HgCl2) −1.3−0.1 (Hg/HgCl2)

110 (2.9)

1000 (∼100)

0−0.5 (Hg/HgCl2)

1951

305 (5 mA cm−2) 925.1 (1)

5000 (87) 600 (∼60)

0−0.5 (Ag/AgCl) −0.2−0.4 (Hg/HgO)

1947 1953

2831 (1)

600 (90)

−0.2−0.4 (Hg/HgO)

1953

72.85 (3) 920 (1)

600 (72.5)

−0.6−0.3 (Ag/AgCl) −0.2−0.4 (Hg/HgO)

1949 1952

2981 (1)

600 (92)

−0.2−0.4 (Hg/HgO)

1952

1443 (1) 597 (1) 62.77 (5 mV s−1)

2500 (76%) 1000 (80)

−0.7−0.2 (Hg/HgCl2) −1.1−0.4 (Hg/HgO) 0−0.5 (Ag/AgCl)

1948 1954 1950

a

Note: Working electrode intends to describe the substrate used to deposit the active material. In the case of a prior slurry preparation, the values in brackets provide its formulation in terms of the ratio between constituents (active material : conductive carbon : binder).

Cu 2MoS4 + 4Li+ + 4e− → Li4MoS4 + 2Cu

which take place at the electrode (pseudocapacitors). In particular, large surface area carbon-based materials have been used for EDLCs. However, their electrochemical performance is poor to satisfy the demand of high power density and energy. Alternative electrodes of carbon-based materials such as pseudocapacitive materials have been used, which have fast and reversible electrochemical reactions. A particularly interesting family of materials are the metal sulfides, due to the different valence states of metal constituents in sulfides and the high theoretical capacity of sulfur and its compounds. Several reports have explored the possibilities of CuS as electrode material for supercapacitors (see Table 30), in which the electrode configuration ranged from electrochemical deposition of CuS on Cu foil,1947,1948 FTO,1949 or stainless steel;1950 and slurry-based preparations, in which the active material, conductive carbon, and a binder are coated onto Ni foam1485,1951,1952 or Ni substrates.1953 Qian et al. proposed the use of CuS nanotubes as electrode and a polysulfide electrolyte for supercapacitors, in which a high power performance was provided in this configuration.1485 In particular, the authors reported specific capacities 17 times higher (2175 F g−1 at 15 A g−1 current density), compared to a CuS-CNT composite (122 F g−1 at 1.2 A g−1).1951 The good electrochemical performance of the supercapacitor was explained by (i) the large surface area and good electron transfer passage, due to the porous nature of the CuS nanotubes (made up of small NCs); (ii) the fast redox reaction of S2−/Sx2− at the electrode/electrolyte interface, which allowed for rich faradic capacitance; and (iii) the low charge transfer resistance of the CuS nanotube electrode in polysulfide electrolyte.1485 In another report, high power performance and good rate capability were also obtained with the use of an acetylene black/CuS nanosheets composite, in which the layered CuS nanosheets allowed for a short diffusion pathway, and a high electronic conductivity was obtained with acetylene black and a synergetic effect between both constituents of the composite.1952

Cu 2MoS4 + 8Li+ + 8e− → 4Li 2S + 2Cu + Mo

However, capacity fading was observed, which revealed the lack of reversibility of the reactions, and this was associated with the formation of Li2S and the inability of Cu to reincorporate into the lattice. In order to improve cycle performance, Cu2MoS4-graphene oxide nanocomposites were also tested.1936 While initial capacity was lower, stability was highly improved, as the functional groups of graphene oxides can be used as trapping sites for Cu ions, reducing its dissolution and improving overall material stability. 10.1.10. Cu-Fe-S. CuFeS2 is also a promising electrode material for LIBs, due to its high theoretical capacity (587 mA h g−1) and good electrical conductivity (103 S m−1). However, there are very few reports in the literature which have studied its electrochemical properties. In the late 1980s, CuFeS2 was used as an additive to improve the discharge characteristics of the Li/CuO system.1937 Almost three decades later, Ding et al. tested CuFeS2 as a cathode for LIBs; however, its cyclability was not explored.1876 Wang et al. subsequently studied the properties of CuFeS2 nanorods as an electrode for LIBs using different electrolytes, in which a discharge capacity up to 180 mA h g−1 was obtained at high scan rates (10 C) in ether-based electrolytes.1938 In addition, the reaction mechanism was considered to be: CuFeS2 + x Li → LixCuFeS2 (1.76 V) LixCuFeS2 + x Li → Cu + Fe + x Li 2S (1.52V) 10.2. Supercapacitors

Supercapacitors are electrochemical systems that have the ability to provide high capacity, high power density, and long cycle life. These systems can store and release energy by charge separation at the interface between the electrode and the electrolyte (electrical double-layer capacitorsEDLCs), and/or by fast and reversible Faradaic electrochemical reactions 6028


Chemical Reviews

Review

11.1. Quaternary Copper Chalcogenides

11. THERMOELECTRIC APPLICATIONS Thermoelectric (TE) devices allow the direct and reversible conversion between heat and electricity.1955 TE devices have multiple uses in the fields of thermal energy harvesting and temperature control. The heat flow induced by an electrical potential can be used in portable fridges, car seat climate control systems, and to precisely cool optoelectronic and electronic devices, such as fiber optic lasers. As generators, their use can be extended to power autonomous electronic systems, to radioisotope thermoelectric generators for space missions, and to improve the energy efficiency of domestic and industrial energy generation processes. In the early part of the twentieth century, Altenkirch showed that to improve the performance of a TE material, it is necessary to increase both the magnitude of its Seebeck coefficient (S, also known as thermopower) and its electrical conductivity(σ), and to reduce its thermal conductivity (κ).1956 These intrinsic material parameters are grouped in the so-called TE figure of merit (ZT), defined as ZT = σS2Tκ−1, where T is the absolute temperature. The coefficient of performance of a TE device solely depends on the operating temperature, the temperature gradient and the ZT. The currently available TE materials struggle to simultaneously display large S, high σ and low κ, due to the strong interdependency of these parameters.1957 Over the last 30 years, the materials science community has been actively searching for different approaches to maximize the TE performance by material design, with the majority of efforts pointing toward a reduction of the thermal conductivity (κ). Heat is transported by charge carriers (electronic thermal conductivity, κel) and phonons (lattice thermal conductivity, κlatt), κ = κel + κlatt. While κel is proportional to σ via the Wiedemann−Franz law, the lattice thermal conductivity depends on the ability of phonons to propagate through the material; thus, it is is the only parameter that is not directly related to the other parameters through the electronic structure of the material. In order to minimize lattice heat transport, scientists have tried to find materials with intrinsically low thermal conductivities, based on the decrease of κlatt with increasing atomic mass, interatomic distances, and/or lattice periodicity. These characteristics are usually found in multinary compounds with complex structures and large unit cells.1958 A complementary strategy to further reduce κlatt is to introduce grain boundaries to scatter medium- and long-wavelength phonons.1959 To-date, nanocomposites composed of a crystalline matrix with nanoinclusions of other phases have provided the best TE performance,1960 due to a more efficient phonon scattering at interphases of different materials with an acoustic impedance mismatch. 1961 The most common approaches to produce bulk nanocomposites are ball-milling and the precipitation of secondary phases from metastable solid solutions. The assembly of precisely designed NC building blocks could be envisioned as an alternative approach to produce TE nanomaterials with high TE efficiency, as it provides the possibility to produce bulk nanocomposites with high accuracy and the required versatility.159,1962−1964 In this regard, many groups have placed special interest in developing Cu chalcogenide NCs and tested their TE properties. The fundamental reasons to explore Cu chalcogenide NC as potential candidates for TE are discussed in the following subsections, complete with an overview of the work done on bulk nanomaterials which are produced by the bottom-up assembly of NCs.

In 2009, Liu et al. proposed chalcopyrite-like Cu-doped quaternary chalcogenides as a novel class of wide-band-gap thermoelectric materials.1020 The idea was based on a convenient slab distribution of the VBM states, with structural units with different functionality. In particular, Cu2CdSnSe4 can be envisioned as being composed of tetrahedral [Cu2−Se4] electrically conductive slabs, which are separated by tetrahedral [Cd−Sn−Se4] electrically insulating slabs. This configuration preserves the hole mobilities in a complex crystal structure (high electrical conductivity), which provides low thermal conductivities. In addition, the homogeneous distribution of the CBM states along the whole unit cell resulted in a poor electron mobility, due to the complexity of the crystal structure. The relatively high hole-to-electron mobility ratios are partially behind the high Seebeck coefficient in this material. Conveniently, the partial substitution of Cu(I) by Cd(II) creates hole carriers and enhances the electrical conductivity, while the disorder (introduced by the partial substitution) increases the phonon scattering and further reduces the thermal conductivity. This finding triggered a series of investigations into the study of other compounds within this family, such as Cu2ZnSnSe4,353,1965−1969 Cu2ZnGeSe4,370,1970 Cu2(ZnFe)SnSe4,1971 Cu2FeSnS4,381 Cu2CoSnSe4,1972 Cu2HgSnSe4,373 and Cu2HgGeSe4,374 with details of the tested TE materials and their performance presented in Table 31. With the idea of combining the advantageous properties of nanostructured materials with this family of tetrahedrally coordinated semiconductors, a variety of off-stoichiometric Cu2-II-IV-VI4 NCs were developed in order to optimize their TE properties. The first reports of quaternary Cu chalcogenides for TE materials were based on the Cu2CdSnSe4 system.357,358 By varying the Cu-to-Cd ratio, the electrical conductivity could be increased and its thermal conductivity could be simultaneously reduced due to the presence of interstitials. To the best of our knowledge, the best ZT obtained for this system to-date is 0.71 for Cu2.15Cd0.85SnSe3.9 at 685 K.357 Chen et al. subsequently reported that the partial replacement of Cu with Ag to form Cu2−xAgxCdSnSe4 resulted in an improvement in the ZT, up to 0.78 at 750 K.1973 Cd-free systems such as Cu2ZnSnSe4, Cu2ZnSnS4, or Cu2ZnGeSe4 have also been investigated. For example, consolidated NCs with a composition of Cu2.15Zn0.85GeSe3.9 achieved a ZT up to 0.55 at 723 K.370 Wurtzite Cu2ZnGeSe4 NCs were prepared and consolidated in another report, but the instability of this phase only permitted a low-temperature study and achieved a maximum ZT below 0.1 at 390 K.371 With regard to Cu2ZnSnSe4, Fan et al. achieved a ZT up to 0.44 at 723 K for Cu2Zn0.03Sn1.10Se2.98; however, the low content of Zn in the material meant that the material could be considered as Cu2SnSe3 slightly doped with Zn.346 Chen et al. reported a ZT of 0.7 at the same temperature by optimizing the composition of Cu2.1Zn0.9SnSe4, and the improved results in this report were based on the further reduction of the thermal conductivity and higher Seebeck coefficients.1023 For sulfur-based NCs, Cudoped CZTS NCs achieved a low ZT of 0.14 at 700 K,339 despite their low thermal conductivity and relatively high Seebeck coefficient, and this result was associated with their intrinsic low electrical conductivity.989 11.2. Ternary Copper Chalcogenides

Cu2SnSe3 (CTSe) is a p-type semiconductor with a direct band gap of 0.84 eV and a 3D distorted diamond-like structure.1974 6029


6030

OLA

EDT

Cu3Sb1−xSnxSe4−ySy

EDT

Cu3SbSe4:Sn

Cu3SbS4:Ni

HAD

OLA OLA HDA, OA

Cu2ZnSnSe4 Cu2ZnGeSe4 Cu2ZnGeSe4

Cu3SbSe4:Sn

OLA OLA

Cu2ZnSnSe4 Cu2ZnSnSe4:Cu

HDA, OA OA, DDT HDA, OA HDA, OA OLA, DDT OLA

OLA

Cu2ZnSnS4

Cu2HgGeSe4 CuFeS2 Cu2GeSe3 Cu2SnSe3 Cu2SnSe3 Cu2SnSe3:Co

OLA

Cu2−xAgxCdSnSe4

HDA

OLA HDA, OA

Cu2CdSnSe4 Cu2CdSnSe4

Cu2HgSnSe4

Surface functionalization

Cu-based NC

−, HP, 350 °C, 80 MPa, 1 h −, HP, 275-400 °°C, 600 MPa, 1 h hydrazine, HP, 400 °C, 75 MPa, 30 min

−, HP, 350 °C, 600 MPa, 1 h

∼100 200-300 20−30

∼50−300

500 °C, 2 h , Ar, press, RT, 2 tons −, HP, 300 °C, 500 MPa 500 °C, 2 h, Ar, press, RT, 2 ton, 5 min 500 °C, 1 h, Ar, HP, 500 °C, 40 MPa, 5 min hydrazine, HP, 623 K, 48 MPa, 20 min hydrazine, SPS, 723−823 K, 45 MPa, 5 min

500 °C, 2 h, Ar, press, RT, 2 tons, 5 min

hydrazine, HP, 350−550 °C, 60 MPa, 30 min −, Press 500 °C, 1 h , Ar, HP, 40 MPa, 5 min

−, SPS, 310 °C, 60 MPa, 10 min hydrazine, HP, 400 °C, 60 MPa, 30 min

hydrazine, SPS, 678 K, 50 MPa, 5 min

1.20

∼1.2

1.2

1.51 1.50 1.2 ∼1.45

1.1 0.98

hydrazine, HP, 550 °C, 80 MPa, 30 min 500 °C, 2 h, Ar press, RT, 5 tons

hydrazine, SPS, 730 K, 45 MPa, 5 min

Eg (eV)

Consolidation (LE/annealing, c. type, T, P, t)

35

14 30 35 6.4 30 15 5 35 ∼20

25 25−30 10-25

20 20−50

10.6

30−50

20−30 15

NC size (nm) Cu2.1Cd0.8Sn1Se3.4 Cu2Cd1Sn1Se4 Cu2.05Cd0.95Sn1Se4 Cu2.15Cd0.85Sn1Se3.9 Cu2.3Cd0.70Sn1Se3.8 Cu2Ag0Cd0.44Sn0.86Se3.34 Cu1.9Ag0.14Cd0.38Sn1.06Se3.58 Cu1.8Ag0.22Cd0.4Sn1.11Se3.58 Cu1.7Ag0.35Cd0.33Sn0.67Se3.58 Cu1.6Ag0.46Cd0.36Sn0.7Se3.51 Cu2Zn0.98Sn1.21S4.36 Cu2.19Zn0.80Sn0.75S3.53 Cu2ZnSnSe4 Cu2Zn1Sn1Se4 Cu2.05Zn0.95Sn1Se4 Cu2.1Zn0.9Sn1Se4 Cu2Zn0.03Sn1.10Se2.98 Cu2Zn1Ge1Se4 Cu2Zn1Ge1Se4 Cu2.15Zn0.85Ge1Se3.9 Cu2Hg1Sn1Se4 Cu2.3Hg0.7Sn1Se3.8 Cu2.3Hg0.7Ge1.0Se4 Cu1Fe0.99S2.08 Cu2Ge1Se3 (WZ-ZB) Cu2Sn1Se3 Cu2SnSe3 Cu2.21Sn0.4Se3 Cu1.92Co0.16Sn0.69Se3 Cu1.82Co0.16Sn0.76Se3 Cu1.94Co0.15Sn0.71Se3 Cu3SbSe4 Cu3Sb0.98Sn0.02Se4 Cu3SbSe4 Cu3Sb0.98Sn0.02Se4 Cu3SbS4 Cu2.25Ni0.75SbS4 Cu3Ni0.05Sb0.95S4 Cu3Sb0.98Sn0.02Se4 Cu3Sb0.98Sn0.02Se3.5S0.5 Cu3Sb0.96Sn0.04Se3.5S0.5 Cu3Sb0.94Sn0.06Se3.5S0.5 Cu3Sb0.90Sn0.10Se3.5S0.5

Composition

523 523 523 700

673

575

733 500 730 730 598 715

723 400 730 730 723

700 700 300 723

∼723 685 685 685 685 750

T (K)

Table 31. Summary of the TE Properties of Cu Chalcogenide Materials Produced from the Bottom-up Assembly of NCs

∼22000 ∼1600 ∼4200 ∼6600 ∼6600 ∼18000 ∼30000 ∼12000 ∼8000 ∼8000 36.7 1388 ∼21700 ∼4000 ∼8000 ∼10000 18000 10953 ∼2000 20382 1000 10000 ∼25000 326.5 3,784 5756 ∼11900 ∼18000 ∼18000 ∼8000 ∼7500 ∼5000 ∼25000 ∼10000 ∼29500 69300 229500 285600 ∼28500 ∼21000 ∼36300 ∼50000 ∼77000

σ (S/m) ∼180 ∼235 ∼200 ∼175 ∼155 ∼175 ∼140 ∼200 ∼260 ∼250 990 301 ∼80 ∼300 ∼240 ∼240 175 81 ∼200 121 ∼260 160 ∼105 816.3 192 225 ∼299 ∼170 ∼170 ∼225 ∼245 ∼300 ∼200 ∼300 ∼230 171.1 171.6 122.4 ∼230 ∼242 ∼200 ∼180 ∼147

S (μV/ K)

∼0.5 ∼0.85 0.439 0.38 0.7 ∼1.9 ∼0.85 ∼0.60 ∼0.65 ∼0.65 ∼0.95 ∼1.15 ∼1.3 ∼1.0 ∼0.90 ∼0.97 ∼1.20 ∼1.0 ∼0.9 ∼1.0 ∼1.0 ∼1.2

∼0.55 ∼0.55 ∼0.55 0.91 ∼0.54 ∼0.80 0.40

∼0.70 ∼0.40 ∼0.30 ∼0.20 ∼0.30 ∼0.60 ∼0.70 ∼0.45 ∼0.40 ∼0.46 0.97 0.645

κ (W/ mK) ZT

∼0.2 0.34 0.264 0.26 0.3 0.34 ∼0.42 0.63 ∼0.47 ∼0.47 0.25 0.5 0.62 1.05 ∼0.13 0.37 0.22 1.1 ∼0.95 ∼0.98 ∼1.05 ∼0.90

∼0.45 ∼0.60 0.70 0.44 0.05 ∼0.10 0.52

∼0.65 ∼0.17 ∼0.40 0.71 ∼0.40 ∼0.65 ∼0.55 ∼0.75 ∼0.78 ∼0.60 ∼0.026 0.14

Ref

1986

293

1861

294

1406 966 215 939 1979 1980

373

346 371 370

1024 1023

339

1973

358 357



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viewed as a three-dimensional Cu−Se framework of distorted [CuSe4] tetrahedra with inserted one-dimensional arrays of [SbSe4] tetrahedral. The largely degenerate VBM is mainly formed by a hybridization of Cu 3d and Se-4p states, while the CBM consists mainly of Sb-5s and Se-4p hybridization. In this configuration, the Cu-Se framework provides paths for hole transport with considerably high mobility.226,227 Furthermore, the complexity of the lattice structure provides low thermal conductivities. However, the carrier density is on the order of 1018 cm−3 at ambient temperature,228 which is not optimum for a high-performance TE material, and thus, an extrinsic p-type dopant should be properly introduced. Different doping strategies have been pursued with the intention of optimizing carrier concentration and, hence, the TE performance. For example, the ionic radius of Sn4+ (0.71 Å) is similar to that of Sb5+ (0.62 Å), and by replacing into the crystal structure (Sb by Sn), holes can be provided. Li et al. prepared Cu3Sb0.98Sn0.02Se4 NCs and achieved a ZT up to 1.05 at 690 K.1861 Co-doping with Sn and S resulted in a slightly larger ZT, specifically 1.1 for Cu3Sb1−xSnxSe3.5S0.5 at 650 K.1986 An alternative dopant is Ni, which provides extra holes and, thus, p-type doping. Two different types of Ni-substituted Cu3SbS4 NCs (Cu3−xNixSbS4 and Cu3NixSb1−xS4) were studied, in which the Cu2.25Ni0.75SbS4 NCs showed an enhanced electrical conductivity, partial variations on the Seebeck coefficient, and reduced lattice thermal conductivity, and yielded a maxmimum ZT of 0.37 at 520 K.293 The TE properties of chalcopyrite CuFeS2 NCs have also been explored, due to its relatively small band gap (0.53 eV) and its use of low-cost, earth abundant elements. In particular, Liang et al. compared CuFeS2 NCs (consolidated at 300 oC and 500 MPa) with the corresponding bulk material.966 The high density of interfaces plus the role of organic ligands (potentially still present) in the NCs provided very low electrical and thermal conductivities, compared to the bulk material. In addition, a higher ZT was obtained for the CuFeS2 NCs (0.26), when compared with the bulk material (ZT = 0.003).

This 3D pattern is partially similar to the two-dimensional layer-like bond network found in I2-II-IV-VI4; thus, it also provides a low thermal conductivity. Furthermore, the Cu−Se bond network stabilizes the structure and forms an electrically conductive framework controlling the hole transport, while the Sn orbitals barely contribute to the p-type carrier transport but allow electrical conductivity to be tuned via partial substitution with a group III element.1975 It is also noteworthy that these ionic substitutions promote phonon scattering, due to the creation of atomic mass fluctuations.301,1976 In this regard, doping with In,1975 Ga,1977 and Mn1978 and isoelectronic alloying with Ge301 on the Sn site have been demonstrated to be effective in enhancing the ZT value. Several reports have studied the possibility of further reducing the thermal conductivity by reducing the material crystal domains to the nanoscale for the Cu2SnSe3 system. As an example, intrinsic Cu2SnSe3 NCs939 showed a significant reduction on the thermal conductivity as expected, but the lack of control on the charge carrier concentration prevented the expected, theoretical TE performance from being achieved.1979 In a recent report, Cu2SnSe3 NCs were synthesized in which the partial replacement of Sn atoms by Co allowed the carrier concentration and therefore electrical conductivity to be enhanced, which rendered Co:Cu2SnSe3 nanostructured TE materials with a ZT up to 0.63 at 715 K.1980 Another system investigated in the I2−IV-VI3 family is Cu2GeSe3. Ibáñez et al. studied the TE properties of orthorhombic (OTR) Cu2GeSe3 NCs and polytypic NCs (which had both zinc blende (ZB) and wurtzite (WZ) domains).215 In this report, similar power factors (σS2) were observed for both NCs, but lower thermal conductivities and a 2.5-fold increase in the thermoelectric figure of merit were obtained for the polytypic WZ/ZB sample, compared to the single-phase OTR sample. The different thermal conductivities were attributed to the high density of the heterojunctions in polytypic WZ/ZB sample and the random cation distribution of the WZ phase. This further extended phonon scattering to short wavelengths, while, at the same time, it could be more efficient than homojunctions because of mismatches in the acoustic impedances. A less explored group of Cu chalcogenide compounds with interest in the TE field is that of I-V-VI semiconductors such as Cu12Sb4S13, Cu3SbA4, and Cu3SbA3 (A = S,Se). Of these, tetrahedrite Cu12Sb4S13 is the most promising due to its natural occurrence and use of nontoxic, earth abundant elements. In addition, tetrahedrites have complex crystal structure with a large number of atoms per unit cell, which provides intrinsically low thermal conductivities in the material. The compound Cu12Sb4S13 exhibits a low electrical resistivity with a positive Seebeck coefficient in the entire temperature range, indicating that the majority of carriers are holes. Undoped tetrahedrites have high carrier concentrations, and doping at the cation site aims to reduce the hole concentration.1981 The most common dopants at the Cu site are Co,1982 Mn,1983,1984 Fe,222 and Zn.222 James et al. prepared Cu12Sb4S13 NCs in a solvothermal approach and tested their TE performance, in which ZT values up to 0.63 at 720 K were reached.1985 By performing the reaction under 300% excess sulfur, famatinite (Cu3SbS4) impurities with detrimental electronic properties were obtained. The electronic properties were largely improved after mixing the resulting material with Zn-rich natural mineral tennatite, with ZT values of 0.85 being achieved after this treatment. Another largely studied material is famatinite Cu3SbSe4. It crystallizes in a ZB-type tetragonal superstructure, that can be

11.3. Binary Copper Chalcogenides

Since 1827, binary Cu chalcogenides have been explored as materials for TE applications, due to their huge versatility, which is tightly related to their interesting and complex phase transition behavior, intrinsically low thermal conductivities, and high thermopower.1987 Numerous works developed with Cu2−xS and Cu2−xSe presented outstanding results (ZT ∼ 2).98,1988,1989 However, these results were obtained at temperatures of around 1000 K, which makes NCs not the ideal building block for these systems. To the best of our knowledge, there is just one report studying the TE properties of Cu2−xSe NC thin films.1990 In this work, the authors investigated the transport properties of Cu2−xSe films, prepared with symmetric and asymmetric organic ligands, and they demonstrated how the polar binding headgroup and organic linker coupling symmetry affected the electrical conductivity and Seebeck coefficient. Although there was no study of the thermal properties of the films, the large thermal impedance mismatch between organic and inorganic NCs anticipated low thermal conductivity values.

12. SENSORS Sensors are analytical devices which interact with the environment to generate a specific output. Depending on the type of interaction, many different sensors can be designed. 6031


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Table 32. Summary of the Different PL-Based Sensors Found in the Literature System CuInS2 CuInS2 CuInS2 CuInS2 CuInS2

CuInS2 CuInS2 CuInS2 CuInS2 CuInS2/ZnS CuInS2 CuInS2 CuInS2 CuInS2 CuInS2 CuInS2 CuInS2 CuGaS2

Analyte Cd2+ Cu2+ Cu2+ Thrombin Dopamine Catechol Pyrogallol Gallate DNR (daunorubicin) Heparin Heparinase Zn2+ Co2+ dicyandiamide alkaline phosphatase (ALP) Glucose fluoride anions Parathion-methyl adenosine-5′-triphosphate (ATP) Acid phosphonates (ACP) l-Noradrenaline (LNE)

Functional group MPA MPA MPA-Cd2+ Fibrinogen MPA/3-aminophenyl boronic acid

3′-NH2-modified capture ONTs L-cysteine L-cysteine 8-aminoquinoline Thioglycolic acid (TGA) 3-aminophenyl boronic acid Tryptophan MPA/3-aminobenzeneboronic acid MPA/3-aminobenzeneboronic acid/ glucose MPA-Pb2+ MPA-Cu2+ MPA-Cu2+/ ATP mercaptoacetic acid

Detection scheme

Detection range

LOD

0.8−72.0 μM 0.2−10 μM 0.1−10.0 μM 6.7 × 10−7 to 0.39 μM 0.5−40 μM 1−100 μM 2.5−80 μM 2.0−80 μM 0.033−0.088 μM 0.05−15 μM 0.2−5 μg mL−1 5−1000 μM 0.3012−90.36 μM 2.0−2000 μM 8.4 to 168 nU mL−1 5−8000 μM 0.1−700 μM

0.19 μM 0.1 μM 0.037 μM 8.7 pM 0.2 μM 0.5 μM 1.5 μM 0.5 μM 0.019 μM 0.012 μM 0.07 μg mL−1 4.5 μM 0.16 μ 0.6 μM 3.6 nU/mL 1.2 μM 0.029 μM

PL PL PL PL PL

Turn-on Turn-off Turn-off Turn-off Turn-off

1996 1996 1996 2011 2008

Ref

PL PL PL PL PL PL PL PL PL

Turn-off Turn-off Turn-on Turn-on Turn-off Turn-off Turn-off Turn-off Turn-on

2012 2006 2006 2002 2000 2009 2014 2010 2010

0.10−38.00 μM 0.5−20 μM

0.06 μM 0.2 μM

PL Turn-on PL Turn-on

2003 2005

6.4−192 nU mL−1 0.5−100 μM

3.1 nU mL−1 5.0 × 10−7 M

PL Turn-off PL Turn-off

2005 2013

outstanding properties make them ideal candidates as probes for optosensing. QDs probes have also been used for detection of metal ions, small molecules, proteins, and nucleic acids.1994 The interaction of metal ions with QDs can either enhance or quench the fluorescence of QDs, and this direct response offers the opportunity to develop very simple and effective probes. In addition, it was proven that this interaction could be mediated by the capping agent present on the QD surface.1995 Since this observation, there has been a growing interest in the search for ideal QD-capping agent combinations to develop environmentally friendly metal ion sensors. MPA-capped CIS QDs were used for the detection of Cu2+ and Cd2+ by monitoring their PL quench and enhancement, respectively.1996 The interaction of Cu2+ with CIS QDs induced a quenching in the PL, due to the reduction of Cu2+ to Cu+ on the surface of the QDs via nonradiative electron−hole recombination through an effective electron transfer process.1997 In the case of Cd2+, the PL enhancement is attributed to the passivation of S2− surface defects due to the interaction with Cd2+.488,1998,1999 Based on these results, the sensor sensitivity for Cu2+ detection was increased, by replacing the original CIS QDs with Cd2+ modified CIS QDs.1996 This is similar to TGA-capped CuInS2/ZnS QDs, which were used for the detection of Co2+.2000 In the presence of Co2+, the PL of CuInS2/ZnS QDs was quenched due to the reaction between Co2+ and the thiol bonds of TGA, as was previously reported for the detection of Co2+ with CdS.2001 Consequently, TGA was detached from QDs, which increased their surface defects and their nonradiative transitions, hence quenching the PL of the QDs. More complex biosensors can be formed by a hierarchical reaction mechanism with different molecules, in a PL turn-on and turn-off QD-complex system. In these systems, one of the most important parameters is the adequate selection of the QD capping agent and, hence, the surface functionalization. These sensors can be fabricated by functionalizing the QDs with a selective ion receptor, where the receptor and its

Various sensing techniques, such as electrochemical detection, functional NC-amplified optical assays, colorimetry, fluorescence, and electrochemiluminescence, have been studied todate. The use of NCs as active sensing elements, transducing components, or even electrodes has allowed for the development of more sensitive and specific sensors. NCs can enhance sensing properties because of their increased surface-to-volume ratio (i.e. larger active sensing area available for the interaction with the target molecules) and their unique size-dependent properties, such as quantum confinement effects or stronger photon and phonon quenching. In addition, the NC surface can be carefully functionalized with different capping agents, with surface functionalization proving to be very important for the design and performance of sensitive nanosensors. A summary of the different PL-based sensors found in the literature is provided in Table 32. 12.1. Fluorescence-Based Sensors

Fluorescence-based detection is the most common method utilized in biosensing because of its high sensitivity, simplicity, and diversity.1991,1992 Generally, fluorescence-based sensors come in a variety of schemes: (a) fluorescence quenching (turn-off), (b) fluorescence enhancement (turn-on), and (c) fluorescence resonance energy transfer (FRET). QD PL arises from the electron-hole (exciton) recombination, which is strongly affected by the changes of surface states or ligands of the QDs. QD PL sensors are based on fluorescence changes, which are induced by the direct physical adsorption or chelation of ions and small molecules on the surface of the QDs.1991 Semiconductor QDs offer significant advantages for sensing, compared with organic dyes. In particular, QDs have size-tunable absorption and emission, broad and intense excitation bands enabling flexibility in excitation, high fluorescence quantum yields, resistance against photobleaching, and large two-photon action cross sections, when compared to established organic dyes.1993 These 6032


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Figure 76. PL-based nanosensor to detect Zn(II) and 8-aminoquinoline by the turn-on/-off of the PL in CIS QDs. Adapted with permission from ref 2002. Copyright 2014 Springer.

solution. This reaction caused the fluorescence quenching, which was then used as a fluorescence probe for the detection of vicinal diols such as dopamine. Moreover, the conjugation between the F-CIS QDs and dopamine further increased the thickness of the surface capping layer and reduced the hydrophilicity of the F-CIS QDs, all of which led to a quench in the fluorescence of the F-CIS QDs. In a separate report, CIS NCs with the same functional group (3-aminophenyl boronic acid) were also used for the detection of dicyandiamide, which was based on the fluorescence quenching of the F-CIS QDs.2009 CIS QDs with a different functional group (3-aminobenzene boronic acid, APBA) were prepared and used as a fluorescence nanosensor for the detection of glucose and fluoride anions.2010 In this report, the APBA functional group could covalently bridge the glucose molecule and this induced the aggregation of APBA-CIS QDs to form a large assembly. The counterbalance of the electrostatic repulsion between QDs and the covalent cross-linking of glucose introduced an elastic tension in the bonds, which stretched the interface of the thiol-modified QDs and created surface states that quenched the fluorescence. However, the authors observed that the addition of fluoride anions into the APBA-CIS QD/glucose system led to the assembly being destroyed, because the fluoride anions strongly interacted with the boron, thus resulting in the quenched fluorescence being dramatically turned-on. A separate report investigated the surface passivation of CIS QDs with fibrinogen, to form a fluorescent probe for the detection of thrombin in human serum.2011 A slightly different approach was used for the detection of daunorubicin (DNR) and for the NIR imaging of prostate cancer cells, by using PC3M prostate cancer cell lines on which the Mucin 1 protein (MUC1) was overexpressed.2012 In this report, novel DNRloaded MUC1 aptamer-NIR CIS QD conjugates were developed, which can be used as a targeted cancer imaging and sensing system. DNR is one of the most widely used anticancer drugs, which can inhibit the proliferation of cancer cells by intercalating into the DNA structure in cell nuclei. In this particular report, CIS QDs were coated with 3′-NH2modified capture ONTs composed of an A10 spacer and a (TCG)7 ONT, which can hybridize with the MUC1 aptamer(CGA)7, and this creates the double-stranded GC-rich region that accommodates multiple DNR molecules. The authors noted that changes in the fluorescence intensity were observed when DNR intercalated into the double-stranded regions, in which sequential decreases in the PL of the QDs occurred with increasing concentrations of DNR. Thus, the observed quenching of the QD fluorescence indicated that DNR could be selectively transported to the targeted cancer cells by binding to the MUC1-QD conjugate, which specifically binds to prostate cancer cells.

interaction with transition metals can modify the QD PL. For example, the functionalization of CIS QDs with 8-aminoquinoline led to the suppression of its PL, in which 8-aminoquinoline interacted with CIS through a hole-transfer mechanism, and this disrupted the radiative recombination process. However, the PL could be recovered by the chelation of Zn2+ with the lone pair electrons of 8-aminoquinoline (Figure 76).2002 The same idea was used in MPA-capped CIS QDs, in which the PL intensity could be quenched by the addition of Pb2+, which stripped MPA from the QD surface and changed the photophysical properties of the QDs.2003 By hydrolyzing parathion-methyl with organophosphorus hydrolase, dimethylthiophosphoric acid (DMPA) was produced, which has a stronger coordinative interaction with Pb2+, and this allowed the PL to be recovered. CIS QDs were also used to detect an important biomarker and indicator of prostate cancer, acid phosphatases (ACP).2004 In this case, the PL of MPAcapped QDs was quenched by Cu2+. However, the authors noted that the addition of adenosine-5′-triphosphate (ATP) could effectively turn-on the PL because of the strong binding between phosphate groups and the metal ion. In addition, it was found that the ATP molecule could be converted into three phosphate molecules by the hydrolysis of ATP. However, this hydrolyzation process led to the PL quenching, due to the dissociation of the Cu2+−ATP complex.2005 The importance of capping ligands on the QD surface was also demonstrated for the detection of heparin. For example, the PL of MPA-capped CIS QDs showed no variation in the presence of heparin, whereas the PL of L-cysteine capped-CIS QDs was quenched in the presence of heparin.2006 This can be explained by the interaction of the amino groups of L-cysteine with the negative sulfate and carboxylate groups in heparin, through electrostatic interactions and hydrogen bonding. Sequentially, the addition of heparinase can hydrolyze the heparin, which results in a turn-on of the PL in the QDs.2006 In the detection of alkaline phosphate (ALP), the PL of tryptophan-functionalized CIS QDs were initially quenched with Cu2+, but were recovered with the complex formation of pyrophosphate (PPi) and Cu2+.2007 However, its detachment from the QD surface resulted in the PL being consequently turned-off, due to the conversion of PPi into two phosphate molecules via the hydrolyzation of ALP and, consequently, the dissociation of the PPi-Cu2+-PPi complex. The surface functionalization of CIS QDs with 3-aminophenyl boronic acid (i.e. F-CIS QDs) has been used for the detection of molecules with a vicinal diol structure, such as dopamine, catechol, pyrogallol, or gallate.2008 In these four cases, the F-CIS QDs containing boronic acid functional groups were found to be reactive toward vicinal diols and formed fiveor six-membered cyclic boronate esters in an alkaline aqueous 6033


Chemical Reviews

Review

Table 33. Summary of Nanosensors Based on Fluorescence Resonance Energy Transfer System

Analyte

Porous network of CuS CuS Cu2−xSe Cu2−xSe

ssDNA Bovine serum albumin DNA thrombin

NC surface

Dyes

Detection range

LOD

Detection scheme

Ref

citrate polydopamine polydopamine

carboxyfluorescein (FAM) Rhodamine B TAMRA TAMRA

0−20 nM 20−500 nM 2−60 nM 1−40 nM

0.8 nM 10 nM 0.5 nM 1 nM

FRET Fluorescent dyes PIET PIET

2019 2020 2021 2021

Table 34. Summary of the Sensors Designed Based on CL Tag

Analyte

Detection Range

Limit of Detection

Ref

Luminol-H2O2-Cu2+ CuS Luminol-H2O2-Cu2+ CuS-Au Luminol-H2O2-Cu2+ 80CuS-Au

sequence-specific DNA sequence of DNA single-nucleotide polymorphisms (SNPs) in genomic DNA

0.02−1.0 × 10−10 M 0.20−1.0 × 10−13 M 0.02−1.0 × 10−10 M

5.5 × 10−13 M 4.8 × 10−15 M 1.9 × 10−17 M

2023 2024 2025

from the surface of the CuS NCs through competitive binding toward CuS. In a different work, polydopamine-embedded Cu2−xSe NCs were found to act as acceptors (fluorescence quenching ability) and were used to construct a photoinduced energy transfer (PIET) pair with DNA-conjugated fluorescent organic dyes donors.2021

While the vast majority of PL-based sensors were based on using CIS QDs, there is one report in the literature employing CuGaS2 (CGS). In this report, mercaptoacetic acid cappedCGS NCs were used as a fluorescence probe for the detection of L-Noradrenaline (LNE).2013 The CGS NCs were demonstrated to have a high sensitivity to LNE with a detection limit of 5 × 10−7 mol L−1, and upon irradiation, the excited electrons transferred from the CGS QDs to LNE, which lead to PL quenching. The authors determined that the LNE molecules were absorbed at the CGS QD surface, due to the hydrogen bond-like interaction between the −COOH group of the mercaptoacetic acid and the −OH group of the LNE molecule and that this bonding interaction may have led to the observed PL quenching. In general, the most common PL-based sensors in the literature are based on fluorescence quenching (turn-off) or enhancement (turn-on). However, Cu chalcogenide NCs have also been used as acceptors in fluorescence resonance energy transfer (FRET), where the PL is provided by a fluorophore. A summary of nanosensors based on the FRET configuration is outlined in Table 33. The FRET electrodynamic phenomenon occurs through a nonradiative process, where a donor fluorophore initially absorbs the energy due to the excitation of incident light, and transfers the excitation energy to the acceptor.2015 Previous studies have shown that the strong binding affinity of biomolecules, which are covalently linked to the NC surface by a fluorophore, can possess fluorescence quenching properties, thus providing an efficient way for sensing/detection.2016,2017 It has been reported that Cu2+ can interact with coordination groups such as carboxyl and hydroxyl through coordinate bonding, and Cu2+ is known to be a highly efficient fluorescence quencher.2018 Furthermore, the interactions between dyes and CuS NCs can be potentially tuned through variation of the dye structures. In this direction, CuS NC networks were mixed with the fluorescent carboxyfluorescein (FAM)-labeled single-stranded DNA (ssDNA) probe and were observed to efficiently quench the fluorescence of the dye-labeled ssDNA.2019 The fast quenching kinetics and efficiency were explained by the strong interactions between ssDNA and CuS, the excellent fluorescence quenching ability of CuS, and the large surface area and porous network property of CuS NCs.2019 In a recent work, Rhodamine B (550 nm), fluorescein (495 nm), Rhodamine 6G (551), and 7-hydroxycoumarin (461) were chosen as fluorescence dye, in which their fluorescence was quenched after forming complexes with CuS NCs, due to the FRET mechanism.2020 The authors observed that the fluorescence could be recovered upon the addition of the proteins or antibiotic bacteria, which displaced the dyes

12.2. Chemiluminescence-Based Sensors

In chemiluminescence (CL)-based sensors, a compound involved in a light emitting reaction is used as a label. The analyte of interest directly participates in a CL reaction, or undergoes a chemical transformation in such a way that one of the reaction products is a coreactant of a CL reaction. In particular, CL assays have been extensively used for detection of nucleic acid and hydrogen peroxide (H2O2). In protic solvents, an oxidation system and an oxidative catalyst are required in alkaline conditions (i.e. pH of 10−13), in which H2O2 is the most frequently used oxidizing agent. Transition metal cations, either free or complexed to organic or inorganic ligands, can catalyze the luminol CL oxidation reaction.2022 In this scenario, CuS NCs have been used as a nanotag to enhance the sensitivity of a CL-based sensor.2023 The CL intensity of luminol-H2O2-Cu2+ was used to determine the target DNA concentration, upon the concentration of dissolved Cu2+ ions from hybrids that were tagged with CuS NCs. Following the same concept, a sensor with a higher sensitivity, high selectivity, and lower detection limit was reported by further amplification of the signal, in which Au NCs were loaded with approximately 100 CuS NCs.2024 A similar setup was later used for the quantification of single-nucleotide polymorphisms, in which higher sensitivity was achieved by incorporating Au NCs loaded with approximately 80 CuS NCs, which provided Cu2+ ions for CL detection.2025 A comparison of the tags, analytes, and detection ranges and limits for CL-based sensors is provided in Table 34. 12.3. Electrochemical-Based Sensors

Electrochemical sensors are typically based on a reaction which produces or consumes electrons, thus producing an electrochemical signal that can be measured by an electrochemical detector.2026 In particular, amperometric biosensors can measure changes in the current on the working electrode, due to direct oxidation of the products of a biochemical reaction. While complexes of Cu(I) ions with various ligands (especially sulfur ligands)2027 were well recognized as systems presenting high activity toward the decomposition of H2O2,2028 the electrocatalytic activity of Cu2S NCs for the redox reaction of H2O2 was not reported until 2007.794 In this case, a glassy carbon electrode was modified with multiwalled carbon 6034


Chemical Reviews

Review

Table 35. Summary of the Electrochemical Sensors Found in the Literature System

Analyte

Cu2S-MWCNT Cu2S-MWCNT Porous CuS NTs AuNPs/CuS-Gr Rose-like Cu2S CuS-RGO CuS/Graphene Porous CuS NTs Porous CuS NTs Cu2S-OMCs Cu2S-DWCNT Cu2S-DWCNT Cu2S-MWCNT Cu2S-MWCNT Sol-gel CuS ultrathin film

H2O2 Glucose glucose DNA Cholesterol H2O2 APL glucose glucose H2O2 Glucose H2O2

CuS-MWCNT

Working electrode

Detection range

LOD

50 nM−10 μM 10 μM−1 mM 0.05−5 μM 0.001−1 nM 0.01−6.8 mM 5−1500 μM 0.1−100 U L−1 0.5−7.5 μM 0.2 μM−2.5 mM 1 - 3030 μM 0.2 μM to 12 mM

human IgA

Glassy carbon Glassy carbon Glassy carbon Glassy Carbon Cu-rod Glassy carbon Glassy carbon Glassy carbon Glassy carbon Glassy carbon Glassy carbon Glassy carbon Glassy carbon Glassy carbon Au electrode

1.81−90.5 ng mL−1

50 nM 10 μM 5 μM 0.1 pM 0.1 μM 0.27 μM 0.02 U L− 0.5−7.5 μM 45 nM 0.2 μM 0.2 μM 10 nM 5 μM 20 nM 1.81 ng mL−1

α-salmonella

Graphite

1 × 103−5 × 105 cells mL−1

400 cells mL−1

Detection scheme

Ref

Amperometric amperometric amperometric amperometric amperometric amperometric Amperometric Amperometric Amperometric Amperometric Amperometric Amperometric Amperometric Amperometric potentiostatic-step to measure capacitance square-wave anodic stripping voltammetry

794 794 2039 2036 2035 2033 2034 2031 2032 2030 2029 2029 2029 2029 2037 2038

nanotube (NT)-modified electrodes, which showed strong electrocatalytic activity toward the oxidation of glucose and exhibited a good linear dependence and high sensitivity to glucose concentration changes.2031 Significantly higher detection limits were obtained using a similar approach, in which Cu electrodes that were decorated with in situ grown CuS NTs were used as a working electrode for the electrochemical detection of glucose.2032 The higher sensitivity in this report was mainly attributed to facile electron transfer of the in situ grown CuS NTs and the Cu electrode, and to the high surface area of the CuS NTs. An alternative working electrode modification for the electrochemical detection of H2O2 was performed by using CuS reduced graphene oxide nanocomposites (CuS/RGO).2033 In this system, the use of RGO with oxygen-containing groups provided large amounts of anchoring sites for the CuS NCs, which rendered more accessible surface sites for the H2O2 reduction process.2033 CuS NC-decorated graphene (GR) nanocomposites were also used as a catalytic amplification platform for the electrochemical detection of alkaline phosphatase.2034 The surface exposed Cu2+/Cu+ states of the CuS NCs presented a strong oxidizing ability for the phenolic-type products of alkaline phosphatase, which could produce electrochemical signals at the electrode surfaces. Cu2S nanoroses/Cu rod integrated electrodes were fabricated in a separate report and were used for amperometric detection of cholesterol.2035 Electrochemical sensors based on CuS-carbon derivates nanocomposites have also been used for the detection of DNA, in which a magnification of the signal was observed by adding Au NCs into the CuS/GR composite.2036 Another approach used for the detection of protein molecules involved designing potentiometric-based biosensors, which consisted of the measurement of potentials at the working electrode with respect to the reference electrode.2026 An insulating ultrathin CuS film, prepared by using a sol-gel technique, was used to detect immunoglobulin, in which the thinness and insulating behavior of the sol-gel CuS films allowed the capacitance changes to be measured via a potentiostatic-step.2037 Less conventional electrochemical sensors using CuS NCs have also been developed. In particular, a highly amplified, NC-based, bio-barcoded electrochemical

nanotubes (MW-CNTs), which were decorated with Cu2S NCs. The authors fabricated amperometric glucose sensors using the Cu2S-MWCNT hybrid nanocomposites and noted that the Cu2S-MWCNT nanocomposites responded more sensitively than those using Cu2S NCs or MWCNTs alone. The enhanced performance of this biosensor was explained by the synergetic effect of the catalytic activity of the Cu2S NCs and the electrical network formed through their direct binding with the MWCNTs, which allowed for excellent electrochemical communication.794 Based on this initial work, the combination of CNTs with Cu chalcogenide NCs has been further investigated, to explore the different sensing properties offered by these nanocomposite combinations. A summary of all the composites used and their sensing performance is outlined in Table 35. For example, Myung et al. fabricated highly sensitive nonenzymatic amperometric glucose biosensors using platinum (Pt) NCs, Cu2S NCs, and SnO2 NCs to form hybrid nanostructures with CNTs, where the NCs were grown in situ on the CNTs in a solvothermal method.2029 The authors observed that the relative sensitivity of the NCs follows the order Pt > Cu2S > SnO2, where the Pt-CNT hybrid nanocomposites showed the highest sensitivity for glucose detection. The unexpected high electrocatalytic activity of the Cu2S NCs toward the target analytes (especially H2O2) was identified as the promoter of the enhanced sensitivity, although their binding interaction with the graphitic layers was weaker than that of the other NCs. In a separate report, Cu2S NCs were incorporated inside the pores of ordered mesoporous carbons (Cu2S/OMCs), in which they exhibited improved electroreduction activity toward H2O2, compared to pristine OMC modified electrodes.2030 The improved electrocatalytic activity was attributed to a series of factors, such as the high dispersion of the Cu2S NCs and their high specific surface area on OMCs, and the accessible ordered mesopores of the Cu2S/OMCs, which provided sufficient room for H2O2 to transport to the electrode surface. In addition, the Cu2S/OMCs nanocomposites provided a favorable environment, which allowed electrons to shuttle between H2O2 and the working electrode, thus enhancing the electron transfer reaction of H2O2.2030 Liu et al. presented a route to form a nonenzymatic amperometric glucose biosensor by using CuS 6035


Chemical Reviews

Review

Table 36. Summary of Cu Chalcogenides NCs That Have Been Used in Other Sensing Methods System

Analyte

CuInZnxS1+x/ZnS CIS CuTe Hollow Cu2−xTe

C-reactive (CRP) H2O2 Nile Red/Nile Blue CO gas

Detection range

LOD

1 ng mL−1 0.1−20 μM

0.3 μM

5 to 500 ppm

5 ppm

Detection scheme

Ref

lateral flow immunoassays system Color variation SERS Electric resistance

2040 2041 255 143

biological environment through their surfaces,2042 thus the NC surface should be carefully designed to obtain the proper NC functionality, such as targeting to specific epitopes, invisibility to the environments (to avoid their clearance by the immune system), or high drug loading capability. Many of the synthetic methods developed to produce Cu chalcogenide NCs render them hydrophobic, which makes them bio-incompatible. Modification of the surface physicochemical properties of NCs can influence particle uptake, biological responses, pharmacokinetics, and biodistribution. Several surface functionalization routes have been explored in order to provide biocompatibility and reduce toxicity in vitro and in vivo. The most common ligand used to provide biocompatible NCs is polyethylene glycol (PEG). PEG is a nontoxic, highly hydrophilic polymer with the lowest level of protein or cellular adsorption, with the advantage that many PEGylated therapeutics have been FDA-approved.2043 In addition, the biological properties of Cu chalcogenide NCs have been investigated using an extensive range of other molecules, such as bovine serum albumin, 2044 cysteine, 2045 phospholipids,597,2046,2047 folic acid,2048 lauric acid,2049 amphiphilic polymer,596,2050 amphiphilic protein−polymer bioconjugate2051 trimethyl(tetradecyl)ammonium bromide (TTAB),2052 polyvinylpyrrolidone (PVP),2053 pluronic F127 block copolymer micelles,896 zwitterionic polymer ligands,2054 glutathione, and mercaptopropionic acid.2055

immunosensor for the simultaneous multiplexed detection of the food pathogens was developed by the immobilization of antibodies onto MWCNT-PAH/SPE.2038 This biosensor was composed of three different NC tracers (CuS, PbS, and CdS), after the NC tracer was dissolved in 0.1 M nitric acid. In this report, the metal ions (Cu2+, Pb2+, and Cd2+) showed distinct, non-overlapping stripping curves by square-wave anodic stripping voltammetry. 12.4. Other Sensing Methods

The use of Cu chalcogenide NCs for sensing is still very much in its infancy. While the most common sensors have been designed, based on the mechanisms described above, there are only a few examples in the literature of sensors which are based on different phenomena (Table 36). Water-soluble CuInZnxS2+x/ZnS NCs (capped with amphiphilic-oligomer-polymaleic-acid dodecanol ester) were used as fluorescent labels in a lateral flow immunoassay (LFIA) system for the detection of C-reactive protein (CRP).2040 CIS NCs can be used to catalyze the oxidation of the peroxidase substrate, 3,3′,5,5′-tetramethylbenzidine (TMB), in the presence of hydrogen peroxide (H2O2), which exhibits a blue color change in aqueous solutions.2041 The color variation, which is dependent on the concentration of H2O2, indicates that a change in absorption can be used for the detection of H2O2. CuTe NCs can also be used as surface-enhanced Raman scattering (SERS) probes for the detection of a wide variety of molecules.255 In particular, CuTe nanoplates and nanocubes were used for the detection of Nile Blue and Nile Red. Hollow Cu2−xTe NCs have also shown gas-sensing properties toward the detection of carbon monoxide, in which chemisorbed oxygen species (O2−, O2−, and O−) were formed due to the absorption of oxygen molecules on the Cu2−xTe NC surface upon air exposure.143 The authors explained that these oxygen species form an electron depletion layer at the NC surface, where the hole density in the valence band increases when the adsorption process reaches equilibrium, which increases the electrical conductivity. However, when the Cu2−xTe NCs are exposed to carbon monoxide molecules, the NCs react with the adsorbed oxygen species and release the trapped electrons back to the conduction band, which decreases the hole density, and this reduces the electrical conductivity. Thus, these changes on the electrical conductivity during the gas adsorption and desorption processes allow these NCs to be used as CO sensors, with detection limits up to 5 ppm.143

13.1. Photothermal Therapy

Photothermal (PT) therapy is based on the hyperthermic ablation of cancer cells through the absorption of electromagnetic radiation in the NIR region.2056,2057 NIR light is required, due to its high transparency to skin, blood, and tissue, which allows for noninvasive and deep tissue penetration.2058 NCs with a strong absorption peak in the NIR region (700−1100 nm) are perfect candidates as thermal coupling agents to enhance the efficiency of PT therapy (PTT). A good PT agent should satisfy the following requirements: (i) strong absorbance in the NIR region; (ii) efficient conversion of the laser energy into heat; (iii) high biocompatibility; (iv) good photostability; (v) small relative size (590

980 nm, 0.72 W/cm2, HeLa 5 min (PTT/DD)

980 nm, 0.72 W/cm2, Hep3B 10 min (PTT) MCF-7 940 nm, 2.0 or 4.0 W/cm2, 5 min (PTT)

Irradiation conditions

DOX 0.12 mg/mg

DOX 9.5%

DOX loading capacity 15.53% DOX Loading content 49.3%

Drug (conc)

Ref

PDT/PTT

PTT/CT

Proposed: PTT/PET

RT/PTT/CT

16.3

2088

576

2186

2117

2171

μ-PET/CT/RT/PTT

2184

610 2005 485 2051 601 2146 2161

2055 2049 2041

606 892

484

2105

2073

2045

2048

2183

2092

2182

2185 7.4 MBq/ mouse 50 μCi of 131I per mouse 33

27.4

70

38.0

PT conversion efficiency (%)

anticancer drugs, antibacterial reagents PTT/DD

MRI, FI ACP detection, FI CLSI CLSI MRI, CLSI FI, UI CT/PTT/PAT

CLSI CLSI FI, CLSI

MRI, FI FI

PET/CRET imaging

PTT/DD

PTT

PTT/DD

PTT/DD

Detection/PTT

PTT

Application



FA

PVP

Cit

PEG−lipid hybrid/PEG-Cu2S-FolDOX conjugates

DSPE-mPEG2000 and DSPEPEG2000-NH2 ACA

PVP

TGA-stabilized

[64Cu]CuS

[64Cu]CuS

Au@Cu7S4

Cu2S

Cu2-xS NDs

CuS NPls

MSN-DNA-CuS

6040

BSA

FA-PEG

PSIOAm

PSIOAm

Y2O3:Yb/Er-CuxS

Cu7S4

Cu7S4

20

14

PVP

rGO-CuS

CuS

12

PEG-silane

10

CuS@mSiO2

PEG

[64Cu]CuS@MSN

15

8

DSPE-PEG2000-NH2

Cu2−xS NPs

40

5

l: 59.4; t: 23.8

12−20

Compound Copper Chalcogenide Nanocrystals - ACS Publications

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