Utilizing Diselenide Precursors Towards the Rationally Controlled

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Cite This: Chem. Mater. 2018, 30, 5704−5713

Utilizing Diselenide Precursors toward Rationally Controlled Synthesis of Metastable CuInSe2 Nanocrystals Bryce A. Tappan, Gözde Barim, Jacky C. Kwok, and Richard L. Brutchey* Department of Chemistry, University of Southern California, Los Angeles, California 90089, United States

Chem. Mater. 2018.30:5704-5713. Downloaded from pubs.acs.org by KAOHSIUNG MEDICAL UNIV on 08/29/18. For personal use only.

S Supporting Information *

ABSTRACT: Within the past decade, there has been an emergence of reports regarding the synthetic isolation of multinary metal chalcogenide nanocrystals that persist under ambient conditions with metastable crystal structures; however, many of the direct syntheses remain largely serendipitous with respect to the conditions needed to achieve the metastable product. Toward the development of more rational design principles that enable the predictable isolation of metastable nanocrystals, we demonstrate a molecular programming approach for the synthesis of CuInSe2 nanocrystals utilizing diorganyl diselenide precursors of the structure R-Se-Se-R. Specifically, we show that the kinetics of diselenide precursor conversion are dependent upon C−Se and Se−Se bond dissociation energies and that the strength of the C−Se bond is the phase-directing variable. When dibenzyl and dimethyl diselenide precursors with relatively weaker C−Se bonds are employed, the resulting nanocrystals form in the thermodynamically stable chalcopyrite phase of CuInSe2. However, precursors like diphenyl diselenide that possess stronger C− Se bonds alter the reaction kinetics so as to steer the reaction toward formation of the metastable wurtzite-like phase. These two phases form via distinct copper selenide intermediates, with chalcopyrite forming through Cu2−xSe and the wurtzite-like phase forming through Cu3Se2 intermediates, and it was found that the ultimate wurtzite-like phase displays remarkable resistance to relaxation to the chalcopyrite phase. This molecular programming approach should be applicable toward the isolation of other metastable phases of metal chalcogenide nanocrystals.

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INTRODUCTION Chemical systems under thermodynamic equilibrium will spontaneously undergo a minimization of their free energy by converging to the lowest energy state or structure possible for a given set of thermodynamic parameters (e.g., temperature, pressure). Higher-energy species may be kinetically stabilized if the activation barrier associated with conversion to the ground state is significant enough to prevent the system from achieving its thermodynamic minimum. In such cases, the kinetically trapped species are referred to as metastable, as they persist despite their thermodynamic instability. Examples of kinetic stabilization of metastable species pervade the natural world; the diamond allotrope of carbon is a classic example.1 In crystalline solids, metastability arises for materials that exhibit energetic differences between polymorphs, compositions, and crystallite size.2,3 As is often the case, metastable materials can display drastically different properties than their more thermodynamically stable counterparts;2,4−6 however, traditional high-temperature solid-state techniques do not provide sufficient synthetic control to predictably yield metastable products. Such methods typically rely on supplying the reaction vessel with an excess of heat in order to overcome the kinetic barriers of solid-state diffusion, which usually drives the formation of the most thermodynamically stable products.7,8 Solution-phase colloidal nanocrystal syntheses, on the other hand, can afford more © 2018 American Chemical Society

kinetic control, providing new pathways to explore metastable materials. Metastability can be inherent to colloidal nanocrystals since these materials possess higher surface-to-volume ratios than their bulk counterparts.9 The increasing contribution from surface energy to the overall lattice energy with decreasing particle size sometimes allows phases that typically form only at high temperatures to be stabilized at significantly lower temperatures at the nanoscale.10,11 Cation-exchange reactions represent a common kinetically controlled technique employed for the synthesis of metastable nanocrystals. For example, syntheses of metastable CoS,12 MnS,13 ZnSe,14 and Cu2−xSe15 nanocrystals via cation exchange have been reported. While cation exchange has proven to be a versatile method to isolate metastable nanocrystals, it may not always be feasible to perform postsynthetic modifications of a parent nanocrystal platform to yield a given metastable product. There also exist numerous examples of syntheses of metastable nanocrystals via direct solution-phase chemistry.16−18 We, and others, have reported the syntheses of metastable semiconductor nanocrystals such as CuInSe2,19,20 Received: May 24, 2018 Revised: July 30, 2018 Published: July 30, 2018 5704

DOI: 10.1021/acs.chemmater.8b02205 Chem. Mater. 2018, 30, 5704−5713

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Cu2SnSe3,21 CuInS2,22,23 CuInxGa1−xS2,24 and Cu2ZnSnS4;25,26 however, these reports have been mainly happenstantial with respect to the isolation of metastable phases of the resulting nanocrystal products. While these discoveries nicely illustrate the potential of using solution-phase chemistry to directly access metastable phases, they do not provide robust synthetic design principles that enable predictable syntheses of metastable phases. Thus, there remains a need for rational methodologies that may be employed toward the fabrication of metastable colloidal nanocrystals. In this vein, chemists have turned to a molecular programming approach by which the properties of nanomaterials can be rationally controlled through careful selection of molecular precursors that follow predictable trends in reactivity.27−31 Diorganyl dichalcogenides (R-E-E-R, where E = S, Se, or Te and R = alkyl, allyl, benzyl, or phenyl) have emerged over the past decade as versatile precursors for the controlled synthesis of metal chalcogenide nanocrystals.32 In part, the utility of this family of precursors stems from their predictable reactivity as a function of the C−E and E−E bond strengths, as first proposed by Vela and co-workers.33 Most notably, they posited that the C−E bond strength is highly dependent upon the identity of the organic functional group of the dichalcogenide, which they used, in turn, to control the morphology of CdS and CdSe nanocrystals. Recent work supports the importance of considering the C−S and S−S bond strengths when diorganyl disulfides are used in the solutionphase synthesis of iron sulfide, which demonstrated that the resulting composition of iron sulfide is tunable from sulfur-rich pyrite (FeS2) to more iron-rich compositions, such as greigite (Fe3S4) and pyrrhotite (Fe7S8), as a function of the C−S bond strength.34 Herein, we show how diorganyl dichalcogenides can be used to rationally control the phase of ternary CuInSe2 semiconductor nanocrystals via a molecular programming approach. Specifically, the judicious choice of R2Se2 [where R = benzyl (Bn), methyl (Me), or phenyl (Ph)] precursor enables kinetic control over nanocrystal formation such that either the thermodynamic chalcopyrite phase or a metastable, hexagonal wurtzite-like phase may be selectively synthesized. Among the different classes of semiconductor nanocrystals, the ternary I-IIIVI2 family, with an A+B3+E2−2 composition, has garnered a great deal of attention for use in photovoltaics, light-emitting diodes, and nonlinear optical devices.24 The most common crystal structure of I-III-VI2 semiconductors is chalcopyrite; the tetrahedral structure of chalcopyrites can be considered a superlattice structure of zinc blende in which the A+ and B3+ ions are ordered in the cation sublattice sites. This is the thermodynamically preferred structure for bulk CuInSe2. A random distribution of the cations leads to the zinc blende phase.24 We reported the first example of a wurtzite-like phase for CuInSe2 nanocrystals, which was a previously unknown phase for the bulk material.19 Subsequent ab initio calculations on this wurtzite-like phase of CuInSe2 revealed it to possess advantageous electronic and optical properties over the chalcopyrite phase, such as increased optical density under near-infrared and visible light.35 The molecular programming approach demonstrated herein can be used to rationally enact selectivity within the CuInSe2 phase space for colloidal nanocrystals.

Article

EXPERIMENTAL SECTION

Materials and General Procedures. Copper(II) dichloride dihydrate (CuCl2·2H2O, Sigma−Aldrich), sodium oleate (>97%, TCI America), indium(III) trichloride (InCl3, 98%, Sigma−Aldrich), diphenyl diselenide (Ph2Se2, 98%, Sigma−Aldrich), dibenzyl diselenide (Bn2Se2, 98%, Alfa Aesar), dimethyl diselenide (Me2Se2, 96%, Sigma− Aldrich), benzeneselenol (PhSeH, 97%, Sigma−Aldrich), and oleylamine (70%, Sigma−Aldrich) were obtained as indicated. Oleylamine was degassed under vacuum at 90 °C for 4 h and then overnight at room temperature prior to use. Reactions were conducted under a nitrogen atmosphere by using standard Schlenk techniques. All reactions employed J-KEM temperature controllers with in situ thermocouples in order to control and monitor the temperature of the reaction vessel. Synthesis of Copper(II) Oleate. An adapted literature approach was used.20 Sodium oleate (3.0 g, 9.85 mmol) and CuCl2·2H2O (0.84 g, 4.93 mmol) were placed in a round-bottom flask. A solution containing 10 mL of ethanol, 8 mL of water, and 17 mL of hexanes was added to the flask, and the mixture was heated to 70 °C. After 25 min, an additional 10 mL of hexanes was added to the solution, and the flask was kept at 70 °C for 4 h. The resulting product collected in the hexanes layer and was washed three times with 30 mL of water in a separatory funnel. The hexanes layer was collected and all volatiles were removed to yield the blue-green Cu(oleate)2 product. Synthesis of CuInSe2 Nanocrystals. We adapted the synthesis of CuInSe2 nanocrystals from Wang et al.,20 as their synthetic method produces nanocrystals with more uniform morphologies. Cu(oleate)2 (0.16 g, 0.25 mmol) and R2Se2 (0.25 mmol; R = Bn or Ph) were loaded into a three-neck round-bottom flask under air. Anhydrous InCl3 (0.055 g, 0.25 mmol when Ph2Se2 was used; 0.066 g, 0.30 mmol when Bn2Se2 was used) was loaded into a two-neck flask under a nitrogen atmosphere. Oleylamine was injected into the two-neck (4 mL) and three-neck (8 mL) flasks under a nitrogen atmosphere. Both flasks were then brought to 70 °C and degassed for 30 min under vacuum. Then the temperature of the flasks was raised to 120 and 140 °C for the two- and three-neck flasks, respectively, and they were degassed for an additional 30 min. The temperature of the three-neck flask was then set to 255 °C. Nanocrystal nucleation was evident prior to injection of the InCl3 precursor as the reaction mixture turned black. The oleylamine solution of InCl3 was injected into the three-neck flask once the reaction mixture reached 230 °C. The reaction was then permitted to heat to 255 °C, at which point it was left to react for 1.5 h. The reaction was thermally quenched by placing it in a room-temperature water bath. The resulting nanocrystal suspension was split equally between two centrifuge tubes that were filled to 40 mL with ethanol, sonicated for 10 min, and centrifuged at 6000 rpm for 2 min. This washing procedure was repeated twice with 10 mL of hexanes used to redisperse the nanocrystals and 30 mL of ethanol as the antisolvent. The nanocrystals were then either dispersed in toluene, for spectroscopic analyses, or dispersed in hexanes and dried down to a powder, for X-ray diffraction and thermal analysis. Synthesis of CuInSe2 Nanocrystals by Use of Me2Se2. For syntheses involving Me2Se2 (bp 155−157 °C), the InCl3 and Cu(oleate)2 precursors were heated in 8 mL of oleylamine, and 4 mL of oleylamine containing an equimolar amount of Me2Se2 was injected at 230 °C. Room-temperature water was used for the condenser attached to the three-neck flask in order to prevent Me2Se2 precursor volatilization. Synthesis of Copper Selenide Nanocrystals. Cu(oleate)2 (0.16 g, 0.25 mmol) and R2Se2 (0.25 mmol, R = Bn or Ph) were placed in a three-neck round-bottom flask and dissolved in 12 mL of oleylamine under nitrogen. The flask was then heated to 70 °C and degassed for 30 min under vacuum. Then the temperature was raised to 140 °C and the flask was degassed for an additional 30 min. The three-neck flask was ramped to 220 °C under nitrogen at 5−6 °C·min−1 and held at that temperature for the desired reaction duration. The reaction was then thermally quenched by placing it in a room-temperature water bath and worked up as previously described. Characterization. Powder X-ray diffraction (XRD) measurements were made on a Rigaku Ultima IV powder X-ray diffractometer using 5705

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Figure 1. (a) Powder X-ray diffraction patterns of CuInSe2 nanocrystals derived from the corresponding diselenide precursors shown at right. The C− Se bond dissociation energies calculated by DFT are given below each precursor. (b) Transmission electron micrographs of the nanocrystals that result from syntheses using the four different selenium precursors. Cu Kα radiation (λ = 1.54 Å). Powder samples were prepared on a zerodiffraction silicon substrate. For the aliquot studies of Figure 2, diffraction patterns were taken on a Rigaku Miniflex 600 powder X-ray diffractometer. Aliquots of the nanocrystal suspensions in hexanes were prepared for XRD by drop-casting the suspension onto a glass substrate. Utraviolet−visible−near-infrared (NIR) absorption spectroscopy was performed on nanocrystal suspensions in toluene in a 150 mm integrating sphere by use of a PerkinElmer Lambda 950 UV−vis−NIR spectrometer and a 1 cm path length quartz cuvette. Transmission electron microscopy (TEM) was performed on dropcast samples supported on holey carbon-coated copper TEM grids (Ted Pella, Inc.). The grids were placed in a vacuum oven overnight at 60 °C for removal of volatile organics. TEM imaging was performed on a JEOL JEM-2100 microscope at an operating voltage of 200 kV, equipped with a Gatan Orius charge-coupled device (CCD) camera. Thermogravimetric analysis (TGA) was performed on a Netzsch STA449c instrument with a heating rate of 5 °C·min−1 with an approximate sample size of 10 mg in an alumina crucible. Fourier transform infrared spectroscopy (FT-IR) was performed on a Bruker Vertex 80 FT-IR spectrometer; solid samples were prepared within a matrix of KBr. Density Functional Theory. Bond dissociation energies were calculated as enthalpies of homolytic cleavage for the bonds of interest. All bond dissociation energy calculations were carried out at the density functional theory (DFT) level using Q-Chem36 through the USC Center for High-Performance Computing. First, geometry optimizations were performed with the 6-31G(d) basis set and the Boese− Martin for kinetics (BMK) functional. Then, single-point energy calculations of the optimized structures were performed with the 6311G(d,p) extended basis set and the BMK functional.

CuInSe2 nanocrystals resulting from reaction with Bn2Se2 was 62% ± 1.5% and 62% ± 1.6% with Ph2Se2, as assessed by organic content-corrected thermogravimetric analysis. With a starting diselenide/Cu molar ratio of 1:1, this result suggests that both selenium atoms from the diselenide precursor are available for nanocrystal growth and the diselenide, therefore, does not act as a limiting reagent. Upon workup, the as-prepared nanocrystals were dispersible in toluene and hexanes, and they maintained colloidal stability for >6 months. When the mechanism of nanocrystal formation is considered for syntheses employing diorganyl diselenides, it is clear that both the C−Se bonds and the Se−Se bond of the diselenide precursor must be cleaved in order to liberate Se2− and incorporate it into the nanocrystal. Consequently, the rate of nanocrystal formation must depend upon the kinetics of dichalcogenide precursor conversion. It follows that the kinetics of nanocrystal formation may then be tunable by modulating the bond strengths (or relative reactivities) of these precursors as a function of the R group substituent. The results of these colloidal nanocrystal syntheses with Ph2Se2 and Bn2Se2, as characterized by powder XRD and TEM, are provided in Figure 1. Consistent with our previous results (although under significantly different synthetic conditions),19 the reaction with Ph2Se2 gives the hexagonal, wurtzite-like phase of CuInSe2 nanocrystals, with lattice parameters of a = 4.08 Å and c = 6.72 Å that are in agreement with the literature.19 Interestingly, the reaction with Bn2Se2 yields the thermodynamic chalcopyrite phase of CuInSe2 nanocrystals, with lattice parameters of a = 5.79 Å and c = 11.56 Å that are also consistent with the literature.39−41 To confirm that this trend could be explained by invoking differences in precursor conversion kinetics, we sought to compare the relative C−Se and Se−Se bond dissociation energies (BDEs) between the diselenide precursors. Density functional theory (DFT) was employed to calculate the enthalpies of homolytic bond cleavage of the C−Se and Se−Se bonds for each precursor. It was found that the C−Se bonds vary more significantly in strength than do the Se−Se bonds for these diselenide precursors, with a net difference of 22 kcal·mol−1 separating the stronger (BDE = 65 kcal·mol−1) and weaker (BDE = 43 kcal·mol−1) C−Se bonds in Ph2Se2 and Bn2Se2, respectively. In comparison, the net difference in BDEs for the Se−Se bonds is 11 kcal·mol−1, with the Ph2Se2 precursor possessing a weaker Se−Se bond (BDE = 42 kcal·mol−1; see

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RESULTS AND DISCUSSION We posited that varying the R group of R2Se2 precursors should allow us to tune the kinetics of nanocrystal formation and possibly the resulting nanocrystal phase. To test this hypothesis, we employed a synthesis adapted from Wang et al.20 to synthesize colloidal CuInSe2 nanocrystals. In short, an oleylamine solution of InCl3 was injected into a hot oleylamine solution of Cu(oleate)2 and R2Se2 precursor, and then the reaction was heated to 255 °C for 1.5 h. Here, oleylamine serves a trifold purpose, acting as the solvent, a reducing agent for the Cu(oleate)2 precursor, and a ligand that coordinates to the surface of the resulting nanocrystals (Figures S1 and S2).37,38 These experimental conditions were held constant unless otherwise noted, with the exception of Me2Se2, which was hot-injected into an oleylamine solution of Cu(oleate)2 and InCl3 because of the volatility of Me2Se2. The ceramic yield of 5706

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Figure 2. (a, b) Phase progression of chalcopyrite CuInSe2 nanocrystals from Bn2Se2 precursor, as monitored by XRD. Times correspond to the time after injection of InCl3. A clear progression can be seen from these copper selenide intermediates to the final chalcopyrite CuInSe2 product. After 3 min, Cu2−xSe and CuSe intermediates are both observed. (c, d) Phase progression of the wurtzite-like CuInSe2 nanocrystals from Ph2Se2 precursor, as monitored by XRD. After 3 min, the umangite Cu3Se2 phase is the primary intermediate.

in which aliquots of the reaction mixture were removed at specific time intervals following the hot injection of the InCl3 solution. Phase progressions of the nanocrystalline products were obtained by taking powder XRD patterns of each aliquot. As seen in Figure 2, both phases of CuInSe2 are the result of reactions of copper selenide nanocrystal intermediates with In3+ ions in solution to yield either the chalcopyrite or wurtzite-like CuInSe2 nanocrystals. Copper selenide intermediates have been observed before in syntheses of CuInSe2 nanocrystals20,40 and have also been used as precursors in the preparation of CuInSe2 nanocrystals via cation-exchange reactions.42,43 In our syntheses, the nucleation of copper selenide nanocrystal intermediates is clearly evidenced by a change in the color of the reaction solution prior to InCl3 injection. These visual cues differ when employing either Ph2Se2 or Bn2Se2 as selenium precursors, which is indicative of their distinct influences on the reaction kinetics. In both cases, the starting oleylamine solution containing the diselenide precursor and Cu(oleate)2 is a vibrant blue at low temperatures. Upon heating to the desired reaction temperature with Ph2Se2, the solution changes from blue to a translucent yellow color at temperatures between 140 and 180 °C. Between 180 and 190 °C, the solution becomes opaque and assumes a light-brown hue. At 200 °C, the solution begins to turn black, indicating the nucleation of copper selenide nanocrystals. In contrast, upon heating a solution containing Bn2Se2 as the selenium precursor, the solution changes directly from blue to opaque black at temperatures just above 140 °C and yields crystalline copper selenide intermediates at temperatures as low as 160 °C (Figure S3), indicative of a lower energy barrier to nanocrystal nucleation for this precursor. Figure 2a shows the XRD pattern of an aliquot taken 3 min after the hot injection of InCl3 into a reaction containing Bn2Se2. This pattern can be indexed to several phases of copper selenide, including intermediates of the Cu2−xSe berzelianite structure

Table S1 for the full listing of Se−Se bond strengths). To better understand the effects of C−Se and Se−Se bonds in phase determination in the CuInSe2 nanocrystal synthesis, a third diselenide (i.e., Me2Se2) was investigated, which was also found to yield chalcopyrite nanocrystals (Figure 1). The Me2Se2 precursor possesses an intermediate calculated C−Se BDE = 57 kcal·mol−1 and a Se−Se BDE = 54 kcal·mol−1. When these BDEs are compared to those of Ph2Se2, which yields wurtzitelike nanocrystals, the difference in BDEs for the Se−Se bonds remains nearly constant at 12 kcal·mol−1, but the difference in BDEs for the C−Se bonds is reduced to 8 kcal·mol−1. These results suggest that employing diselenide precursors with weaker C−Se bonds than those of Ph2Se2 results in the formation of the thermodynamic, chalcopyrite phase. To confirm that kinetic control of CuInSe2 nanocrystal formation depends on the strength of the C−Se and not the Se−Se bond, two control experiments were conducted under otherwise identical synthetic conditions, in which benzeneselenol (PhSeH) and diphenyl selenide (Ph2Se) were employed as two selenium precursors lacking Se−Se bonds. As in the case of Ph2Se2, both of these precursors have strong calculated C−Se bonds (BDE = 78 kcal·mol−1 for PhSeH and 71 kcal·mol−1 for Ph2Se), meaning both could impart similar kinetic effects on phase determination. As demonstrated in Figure 1, the reaction with benzeneselenol indeed yields the wurtzite-like phase, as expected. In contrast, Ph2Se did not yield any isolable crystalline product, likely because two strong C−Se bonds must be broken to liberate each equivalent of selenium and this does not readily occur before the volatilization of the precursor (Ph2Se bp 115− 117 °C), even when Ph2Se is injected into a hot solution containing Cu(oleate)2 and InCl3. CuInSe2 Formation Pathways. To further probe the mechanisms of chalcopyrite and wurtzite-like phase determination for these CuInSe2 nanocrystals, studies were performed 5707

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band gap of ∼1.0 eV.53,54 Consistent with the assignment of Cu2−xSe as the intermediate that results in CuInSe2 formation, the chalcopyrite CuInSe2 nanocrystals derived from Me2Se2 also pass through this Cu2−xSe intermediate (Figure S9). In addition to the evidence provided by aliquot studies, the highly similar structures of antifluorite Cu2−xSe and chalcopyrite CuInSe2 suggest that this phase transformation should occur readily. The chalcopyrite structure can be described as a supercell of the zinc blende structure, comprised of a Se2− facecentered cubic (fcc) lattice in which 50% of the tetrahedral holes are filled by bands of Cu+ and In3+ that alternate along the [100] and [001] directions (Figure 3). While the Cu2−xSe structure

and the hexagonal klockmannite phase of CuSe. These two phases of copper selenide represent the most thermodynamically stable species at high temperatures under copper-rich (Cu2−xSe) or slightly more copper-deficient conditions (CuSe), as per the bulk Cu−Se binary phase diagram.44 To confirm that these phases correctly represent the reactive copper selenide intermediates, a control experiment in which Bn2Se2 and Cu(oleate)2 were heated in oleylamine in the absence of InCl3 led to the formation of berzelianite Cu2−xSe and klockmannite CuSe nanocrystals (Figure S4), which is consistent with the intermediates observed in Figure 2a,b. The early disappearance of the CuSe phase in Figure 2b indicates that CuSe is converted either to CuInSe2 directly or to the Cu2−xSe phase, which in turn yields chalcopyrite CuInSe2 upon diffusion of In3+ into the copper selenide crystal lattice. The direct synthesis of chalcopyrite CuInSe2 via the combination of CuSe and InSe precursors has been reported,45 but since no crystallographic evidence of InSe formation is observed in Figure 2a,b, this mechanism is ruled out. Because Cu2−xSe was observed to persist for longer than CuSe as the last remaining intermediate prior to the ultimate formation of CuInSe2, we hypothesized that the CuSe in solution converts to Cu2−xSe, which subsequently forms chalcopyrite CuInSe2. To test this hypothesis, a control experiment was performed in which aliquots were removed at specific time intervals from a solution in which Bn2Se2 and Cu(oleate)2 precursors were heated together in oleylamine under the same conditions as those employed in the aliquot study of Figure 2. However, this control revealed that after the initial nucleation of a mixture of Cu2−xSe and CuSe nanocrystals, CuSe persisted and did not convert to Cu2−xSe over time (Figure S5). This finding indicates that the consumption of Cu2−xSe in the formation of CuInSe2 with In3+ is required to drive the transformation of CuSe to Cu2−xSe in situ. Indeed, the diffusion of In3+ into Cu2−xSe to form CuInSe2 necessarily expels 1 − x Cu atoms from the crystal lattice into solution. These 1 − x Cu atoms may then recombine with CuSe nanocrystals to generate Cu2−xSe. This hypothesis is consistent with the fact that Cu2−xSe is more thermodynamically stable than klockmannite CuSe under Cu-rich conditions, and as observed in Figure 2b, the Cu2−xSe intermediate persists after the disappearance of CuSe. The aliquot studies of Figure 2 were performed at a lower temperature (220 °C) than the conditions employed to synthesize phase-pure CuInSe2 nanocrystals in order to better trace the transient intermediates that yield either chalcopyrite or wurtzite-like CuInSe2. Analogous aliquot studies at the reaction temperature of 255 °C were performed for both Bn2Se2 and Ph2Se2 to confirm that the intermediates identified from Figure 2 were the same as those present at higher temperatures (Figures S6 and S7). While Figure S6 shows that the klockmannite CuSe phase could not be observed due to fast conversion at 255 °C, the Cu2−xSe intermediate was still present at early reaction times. Cu2−xSe is a highly p-doped semiconductor and displays localized surface plasmon resonance (LSPR) signatures in the near-infrared region of the spectrum when x > 0.50−52 UV−vis− NIR spectra were taken for aliquots of this reaction employing Bn2Se2, and the spectra provide additional confirmation of the presence of Cu2−xSe intermediates by the broad LSPR feature centered between 1100 and 1600 nm (Figure S8). As the reaction proceeded to form chalcopyrite CuInSe2, the LSPR feature decreased in intensity until giving way to the absorption spectrum of chalcopyrite CuInSe2, which displays broad absorption through the visible region with a characteristic

Figure 3. Illustration of phase transformation of cubic antifluorite-type Cu2−xSe structure to tetragonal chalcopyrite CuInSe2 structure upon exchange of Cu for In cations (Cu = blue, In = pink, Se = green). The fcc Se sublattice is effectively unchanged in this transformation; the hightemperature lattice constant (433 K) for Cu2Se is 5.787 Å, which would require an expansion of only 0.078% to match the lattice constant of the chalcopyrite Se sublattice, 5.792 Å at 433 K.46,47 CIF files for the Cu2−xSe and chalcopyrite CuInSe2 phases were taken from refs 48 and 49.

depends upon copper stoichiometry at lower temperatures, the anionic sublattice of Se2− in Cu2−xSe is fcc at temperatures above 130 °C for all values of x.46 The high-temperature (433 K) lattice constant of antifluorite-type Cu2Se has been reported to be a = 5.787 Å.46 For comparison, the lattice parameter of the Se2− sublattice in chalcopyrite CuInSe2 is a = 5.792 Å at 433 K,47 which represents an expansion of the Cu2Se lattice by only 0.078%. Previous work has shown that the soft Cu+ ions from the relatively copper-rich Cu2−xSe phase has made this phase a useful starting material for nanoscale cation-exchange experiments, as Cu+ easily diffuses out of the structure to be replaced by harder Lewis acids.15,55 The near-exact match of the Se2− sublattices between Cu2−xSe and CuInSe2, the mobility of Cu+ ions through the structure,56,57 and easy In3+ diffusion into the structure facilitate the transformation of Cu2−xSe to chalcopyrite CuInSe2. The wurtzite-like CuInSe2 nanocrystals from the reaction with Ph2Se2 form through a different intermediate, namely, the Cu3Se2 umangite phase of copper selenide (Figure 2c,d). As a control to verify that Cu3Se2 is the reactive intermediate that precedes wurtzite-like CuInSe2, Ph2Se2 was heated in oleylamine with Cu(oleate)2 to generate Cu3Se2, as illustrated in Figure S10. Cu3Se2 possesses a tetragonal crystal structure58,59 that contains two crystallographically unique copper sites (referred herein as first- and second-position copper atoms, labeled in blue and pink, respectively, in Figure 4d,e) and a single crystallographically distinct Se2− site. Although formally the stoichiometry of umangite would necessitate mixed oxidation states of 5708

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Figure 4. (a) Se sublattice of Cu3Se2. Each Se atom has six nearest neighbors that are nearly in plane. These planes of Se alternate in an ABAB fashion along the [010] direction. (b) Se sublattice of wurtzite-like CuInSe2. Close packing occurs along the [001] direction. By XRD, the experimental dspacing for the (002) planes in our wurtzite-like CuInSe2 nanocrystals is 3.343(2) Å, which requires a 4.44% expansion of the interplanar Se···Se distance within umangite as obtained from the crystal structure of Cu3Se2. (c) Wurtzite-like structure of CuInSe2; it was assumed that Cu and In cations are randomly distributed throughout the tetrahedral sites in the structure, so the blue tetrahedra represent both Cu and In positions. (d) Side view of the Cu3Se2 structure along the [001] direction. Note that the first Cu position (blue) shares four edges (highlighted in yellow) with second Cu position (pink). Each pink tetrahedron shares two edges (highlighted in red) with a blue tetrahedron and one edge (highlighted in teal) with an adjacent pink tetrahedron. (e) Top view of Cu3Se2 looking down the direction of close packing [010]. Pink tetrahedra are corner-sharing along the [001] direction, and there exist tetrahedral holes that prevent corner-sharing connections between pink tetrahedra along the [100] direction. These tetrahedral holes are demarcated by dotted red lines. In going to the wurtzite-like structure, Cu or In ions must fill the tetrahedral holes indicated in panel e. To maintain charge neutrality, two Cu ions must diffuse out of the Cu3Se2 intermediate per In3+ ion; one Cu must come from the first position of Cu (blue tetrahedra) and the other from the second position (pink tetrahedra). (f) Structure generated by removing these tetrahedra from the structure; the red dotted lines represent the tetrahedral holes that are filled in the transformation to the wurtzite phase, giving a structural motif similar to that shown in panel c. The CIF file for the Cu3Se2 phase was obtained from ref 59.

copper to balance two Se2− anions, X-ray photoelectron spectroscopy (XPS) measurements of umangite have revealed that copper is monovalent and the oxidation state of selenium is −3/2.60 However, XPS indicates that the oxidation states of Cu, In, and Se are +1, + 2, and −2, respectively, in chalcopyrite and wurtzite-like CuInSe2.20,61 In going from Cu3Se2 to CuInSe2, selenium must undergo a reduction to the −2 oxidation state with a concomitant decrease in the copper content as In3+ diffuses into the structure. To aid in thinking about this transformation, employing the formal Cu oxidation states of +1 and +2 in Cu3Se2 is a useful bookkeeping tool, even if these oxidation states do not accurately represent the electronic structure of Cu3Se2; to produce two Se2− anions that originated from the −3/2 oxidation state, one copper atom must be oxidized to Cu2+ for every two selenium atoms. In this way, umangite can be thought of as Cu+2Cu2+Se2−2. Upon diffusion of In3+, the equivalent of one Cu+ and one Cu2+ ion must diffuse out of the material to maintain charge neutrality and produce the correct Cu+In3+Se2−2 stoichiometry. The conversion of Cu3Se2 to wurtzite-like CuInSe2 nanocrystals represents a larger structural transformation from the intermediate to the final product than does the Cu2−xSe to chalcopyrite CuInSe2 conversion. The wurtzite-like CuInSe2 structure may be considered as a hexagonally close-packed framework of Se2− anions throughout which the Cu+ and In3+ cations fill 50% of the tetrahedral holes, creating a network of

corner-sharing selenium-terminated tetrahedra, as illustrated in Figure 4c. This hexagonal close packing mandates that each Se2− has six in-plane nearest neighbors (Figure 4b). To envision the structural similarities between the wurtzite-like phase of CuInSe2 and Cu3Se2, it is noted that the Se2− sublattice within the Cu3Se2 structure maintains a quasi-planar hexagonal framework of Se2− anions that stack in an alternating ABAB fashion (Figure 4a). The interplanar distance between these anion layers is 3.201 Å, which is close to the 3.343 Å d-spacing for the (002) planes of wurtzite-like CuInSe2 nanocrystals, as experimentally measured by XRD. Thus, to assume the structure of the wurtzite-like Se2− sublattice, the Cu3Se2 Se2− sublattice needs to expand by 4.44% in the interplanar direction while assuming a higher degree of planarity within each Se2− sheet. Knowing that the structure of the anionic sublattice of Cu3Se2 can be viewed as a slight distortion of the wurtzite-like anionic sublattice, we now turn to understand how the structure as a whole changes upon In3+ incorporation to give way to the wurtzite-like product. Within the umangite structure, the two distinct copper sites in Cu3Se2 form two types of Se-terminated tetrahedra: the first copper site (blue tetrahedra in Figure 4d,e) is a distorted tetrahedron that shares four edges with adjacent second-position copper tetrahedra (pink tetrahedra in Figure 4d,e). The second copper site can be thought of a copper atom at the center of a tetrahedron that shares an edge with the vicinal second-position copper tetrahedron and two edges with the 5709

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Figure 5. (a) TGA curves after heating of wurtzite-like CuInSe2 nanocrystals to 300 °C. No additional mass loss was sustained from the material after the first cycle of heating. (b) Powder XRD patterns taken after each cycle of heating to 300 °C show that the wurtzite-like phase is thermally stable up to this temperature. (c) TGA curve after heating of wurtzite-like CuInSe2 nanocrystals to 420 °C. The material lost ∼13% of its weight due to volatilization of surface ligands from the nanocrystals. (d) Powder XRD indicating that a nearly complete phase transformation to the thermodynamic chalcopyrite phase results after one cycle of heating to 420 °C.

adjacent first-position copper tetrahedron. These edge-sharing configurations are depicted best in Figure 4d; because the wurtzite-like structure is composed entirely of corner-sharing tetrahedra, diffusion of In3+ into Cu3Se2 must disrupt this edgesharing motif to yield a framework of only corner-sharing tetrahedra. The perspective illustrated in Figure 4e highlights the existence of tetrahedral holes in the umangite structure that, if filled, would yield the necessary corner-sharing structure within the (010) plane. As previously mentioned, the diffusion of every In3+ ion into the umangite intermediate causes the removal of two copper ions (one Cu+ and one Cu2+). Occupation of the tetrahedral holes highlighted in Figure 4e would cause the vicinal first-position copper sites to diffuse out of the structure to prevent the formation of an unstable face-sharing configuration. In addition, occupation of this tetrahedral hole should also expel one of the second-position copper atoms that would otherwise assume an unstable edge-sharing motif with the newly formed indium-centered tetrahedron.62 Removal of these two copper atoms per In3+ cation in the structure leads to the structure shown in Figure 4f, where the tetrahedral holes indicated by dashed lines would be filled either by the incoming In3+ ions or by Cu+ ions that shift to those sites, thus enabling In3+ incorporation at another corner-sharing tetrahedral site in the structure. In summary, we can understand this phase transformation by considering a distortion of the Cu3Se2 sublattice to the wurtzite-like Se2− sublattice and by assessing which holes are filled upon In3+ incorporation and how that affects each copper site in umangite. The distinct berzelianite and umangite intermediates, and the aforementioned discrepancy in the temperatures at which these copper selenides nucleate with Bn2Se2 and Ph2Se2 precursors, suggest that the stronger C−Se bond of Ph2Se2 contributes to an elevation of the activation

energy barrier associated with formation of Cu2−xSe intermediates and steers the reaction toward formation of wurtzitelike CuInSe2 via a metastable Cu3Se2 intermediate. Assessing the Persistence of the Wurtzite-Like Phase. Since the wurtzite-like phase is a metastable phase at room temperature with respect to the bulk material, we wanted to test its structural persistence and ascertain whether or not a phase transition to the chalcopyrite structure could be achieved by thermal activation. After synthesis and workup, the wurtzite-like CuInSe2 nanocrystals remain in this phase at room temperature for greater than 1 year. As a preliminary experiment, a powder sample of the wurtzite-like CuInSe2 nanocrystals was dispersed in 1-octadecene, a high-boiling, noncoordinating solvent, and heated to 300 °C for 2.5 h. Powder XRD measurements of the nanocrystals were taken before and after this treatment; the results show that heating the wurtzite-like nanocrystals in solution did not induce a phase change to chalcopyrite (Figure S11). The analogous experiment was then carried out with wurtzite-like nanocrystals in the solid state. The nanocrystals were heated under nitrogen to 300 °C at 5 °C·min−1 and experienced ∼3.5 wt % loss due to loss of residual organics, as demonstrated by the thermogravimetric analysis (TGA) trace provided in Figure 5a. After the nanocrystals were heated to 300 °C, held at that temperature for 10 min, and cooled back to room temperature, XRD revealed that no phase change occurred. This heating−cooling cycle was repeated two additional times, and XRD analysis of the nanocrystals shows that the wurtzite-like phase remains persistent upon heating and cooling, as shown in Figure 5b. Moreover, Scherrer analysis of the corresponding XRD patterns suggests that the grain size of the nanocrystals (i.e., 17 nm) does not change during these heating−cooling cycles. In contrast, upon heating the wurtzite-like nanocrystals to 5710

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420 °C, a greater fraction of the surface ligands was removed (representing ∼13 wt % loss by TGA; Figure 5c), which resulted in a phase transformation from the metastable wurtzite-like phase to predominantly the thermodynamically stable chalcopyrite phase, as illustrated in Figure 5d. Interestingly, while we observed that a postsynthetic transformation from wurtzite-like phase to chalcopyrite phase does not occur at 300 °C either in solution or in the solid state, we discovered that a mixture of these phases results when the nanocrystals are synthesized at 300 °C with Ph2Se2 (Figure S12). Since this mixture cannot be obtained from a direct wurtzite-like-to-chalcopyrite phase transformation at this temperature, it must be a result of the chemistry of copper selenide nanocrystal intermediates. Since Cu2−xSe is thermodynamically favored above 123 °C while Cu3Se2 is metastable at elevated temperatures,44 the mechanisms to produce wurtzitelike or chalcopyrite CuInSe2 must be competing at higher reaction temperatures, as dictated by the extent of a phase transformation from intermediate Cu3Se2 phase to Cu2−xSe phase. Such a transformation for copper selenide is known to occur at elevated temperatures,63 and a mixture of phases of CuInSe2 can be explained by an incomplete transformation of Cu3Se2 to Cu2−xSe, where the resultant Cu2−xSe yields chalcopyrite and the remaining Cu3Se2 forms wurtzite-like CuInSe2 nanocrystals. Several control experiments were performed that support this conclusion: (1) When Ph2Se2 and Cu(oleate)2 were heated in oleylamine at low temperatures for short times (3 min), pure Cu3Se2 resulted, as shown in Figure S10. In a typical synthesis of wurtzite-like CuInSe2, InCl3 is injected at 230 °C and the temperature is then increased to 255 °C for the duration of the reaction. These reaction conditions allow for diffusion of In3+ into Cu3Se2 before a phase transformation from umangite to Cu2−xSe can occur. (2) However, if Ph2Se2 and Cu(oleate)2 are heated together at low temperatures for longer times (15 min), then a partial phase transformation from Cu3Se2 to CuSe and Cu2−xSe phases results (Figure S13). This demonstrates that Cu3Se2 is a kinetic, metastable intermediate and that timely injection of InCl3 is necessary to obtain phase-pure wurtzite-like CuInSe2 from the Cu3Se2 intermediate nanocrystals before they convert to the more thermodynamically stable high-temperature berzelianite and klockmannite phases. The formation of a mixture of chalcopyrite and wurtzite-like CuInSe2 nanocrystals when Ph2Se2 is used at an elevated reaction temperature of 300 °C indicates that indium incorporation into the Cu3Se2 is a diffusion-limited process that does not give quantitative yields of wurtzite-like CuInSe2 within minutes of InCl3 injection, likely due to the complicated nature of the umangite-to-wurtzite-like phase transformation. Sustained heating of the reaction vessel is important to drive the formation of CuInSe2, and when this heating step is performed at higher temperatures, the rate of conversion of Cu3Se2 to CuSe or Cu2−xSe becomes kinetically competitive with the rate of phase transformation from Cu3Se2 to wurtzite-like CuInSe2. (3) As a final control experiment, it was found that when Ph2Se2 was reacted with Cu(oleate)2 at 300 °C, the thermodynamically stable Cu2−xSe phase resulted, as observed by XRD and UV−vis−NIR analysis (Figures S14 and 15). Thus, when isolating metastable materials, it is important to consider not only the persistence of the metastable state but also the various thermodynamic minima that may exist for intermediate compositions that give way to the final metastable or thermodynamic products.

CONCLUSIONS In conclusion, we have demonstrated a molecular programming approach using diorganyl diselenides that predictably allows for control over the resulting phase of colloidal CuInSe2 nanocrystals as a function of C−Se bond strength. It was found that the chalcopyrite phase forms from precursors with weaker C−Se bonds via a fast-nucleating Cu2−xSe nanocrystal intermediate, while the metastable wurtzite-like phase forms from precursors with stronger C−Se bonds through a slow-nucleating umangite Cu3Se2 phase. By comparing the Se2− sublattices to those of the final CuInSe2 phases and by considering the mobility/extraction of copper from these intermediate structures, we developed a structural rationalization for phase determination. The wurtzitelike phase of CuInSe2 nanocrystals was found to be kinetically persistent for many months at room temperature and up to at least 300 °C upon multiple heating and cooling cycles, but transitions to the chalcopyrite phase occurred upon decomposition and loss of the surface ligands at 420 °C. Therefore, even though the wurtzite-like phase is a high-temperature phase of CuIn(S,Se)2 in the bulk phase diagram, it can be kinetically accessed on the nanoscale at relatively low temperatures and remain persistent in its bulk metastable phase. There are multiple examples of diorganyl dichalcogenide precursors being used to serendipitously access metastable hexagonal phases of multinary chalcogenide nanocrystals, including CuInS2,64 Cu2ZnSnSe4,65 Cu2ZnSnS4−xSex,66,67 and Cu2SnSe3,21 instead of their thermodynamically preferred tetragonal or monoclinic phases. The molecular programming approach presented here offers the possibility of being able to rationally access these metastable hexagonal phases as a function of dichalcogenide C−Se or C−S bond strengths, moving beyond the traditional model of size-reducing bulk thermodynamic phases.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b02205. One table listing calculated bond dissociation energies; 20 figures showing control experiments, additional aliquot studies, XRD patterns, UV−vis−NIR spectra, and TEM images (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Twitter https://twitter.com/ BrutcheyGroup. ORCID

Richard L. Brutchey: 0000-0002-7781-5596 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Award DEFG02-11ER46826. Computation for the work described in this paper was supported by the University of Southern California’s Center for High-Performance Computing (hpc.usc.edu). We acknowledge Professor B. Melot for insightful discussions relevant to the structural chemistry in this paper. 5711

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DOI: 10.1021/acs.chemmater.8b02205 Chem. Mater. 2018, 30, 5704−5713

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DOI: 10.1021/acs.chemmater.8b02205 Chem. Mater. 2018, 30, 5704−5713