PbTe

Subscriber access provided by Universitaetsbibliothek | Johann Christian Senckenberg

The Kirkendall Effect: Main Growth Mechanism for a New SnTe/PbTe/SnO Nano-Heterostructure 2

Arthur Shapiro, Youngjin Jang, Faris Horani, Yaron Kauffmann, and Efrat Lifshitz Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b01455 • Publication Date (Web): 16 Apr 2018 Downloaded from http://pubs.acs.org on April 16, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 10 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

The Kirkendall Effect: Main Growth Mechanism for a New SnTe/PbTe/SnO2 Nano-Heterostructure Arthur Shapiro,†,‡ Youngjin Jang,†,‡ Faris Horani,† Yaron Kauffmann,§ and Efrat Lifshitz*,† †Schulich Faculty of Chemistry, Solid State Institute, Russell Berrie Nanotechnology Institute, Nancy and Stephen Grand Technion Energy Program, and §Department of Materials Science and Engineering, Technion–Israel Institute of Technology, Haifa 3200003, Israel

ABSTRACT: Attention to semiconductor nanostructures with a narrow bandgap energy and low production cost has increased in recent years, due to practical demands for use in various opto-electronics and communication devices. Colloidal nanostructures from the IV-VI semiconductors, such as lead and tin chalcogenides, seem to be the most suitable materials platform; however, their poor chemical and spectral stability has impeded practical applications. The present work explored the mechanism for formation of new nanostructures, SnTe/PbTe/SnO2, with a core/shell/shell heterostructure architecture. The preparation involved a single-step post treatment for the pre-prepared SnTe cores, which simultaneously generated two different consecutive shells. The process followed a remarkable Kirkendall effect, where Sn ions diffused to the exterior surface from a region below the surface and left a ringlike vacancy area. Then, Pb ions diffused inward and created a PbTe shell, filling the Sn-deficient region. Finally, the ejected Snions at the exterior surface underwent oxidation and formed a disordered SnO2 layer. These intriguing processes were corroborated by a theoretical estimation of the relative diffusion length of the individual elements at the reaction temperature. The nanostructures which were comprised of low-toxic elements were endowed with optical tunability, and chemical stability, which lasted more than one month at ambient conditions.

Introduction Colloidal nanostructures have attracted substantial attention due to their intriguing optical and electronic properties derived from the quantum confinement effect.1-3 Nanostructures have been implemented in a wide variety of applications such as solar cells,4-9 biological tagging,10-12 thermoelectrics13-16 and light-emitting diodes (LEDs).17-22 In particular, there has been increasing interest in materials with optical activity in the near- and short-wave infrared for various opto-electronic and communication technologies. Colloidal semiconductors with a narrow bandgap could supply a solution, with the surplus benefit of low production costs. However, colloidal nanostructures often suffer from chemical instability, related to the high reactivity of their exterior surfaces to oxidation or exchange with foreign atoms, each of which causes considerable degradation of the nanostructures' properties.23-25 Therefore, achieving optimum performance and sustainable properties via control of nanostructures' surfaces is an issue of a paramount importance.26 Preparation of heterostructures (e.g., core/shell) should offer interface control, providing surface passivation27, 28 as well as spectral stability.29-32 Bulk SnTe is a narrow bandgap IV-VI semiconductor which possesses a direct bandgap of 0.18 eV at room temperature and a cubic rock-salt crystal structure.33-35 This material has interesting characteristics such as small electron and hole effective masses (meh≈0.025 m0),35 high dielectric constant (ε∞≈45)35 and an extremely large exciton Bohr radius (95 nm),35 providing strong quantum confinement effect. In addition, SnTe is a good alternative material for practical applications because of the low toxicity of Sn to the environment.35-37

Pioneer work in the field of colloidal synthesis of SnTe nanostructures was reported by Talapin and co-workers.33 A few other successful attempts for synthesizing colloidal SnTe nanostructures have been reported based on different Sn precursors and ligands.34, 35, 38 Recently, the Klimov group reported the use of diisobutylphosphine for preparing SnTe nanorods and nanowires.35 The Lifshitz group developed a procedure by introducing activating agent (1,2-hexadecanediol) to boost the low reactivity of the inexpensive Sn precursor (SnCl2).38 Reiss and co-workers reported a strong tendency of SnTe for surface oxidation.39 Additional studies developed Sn-based core/shell heterostructures for the reduction of the SnTe surface activity.38, 39 The Lifshitz group reported a prototype core/shell nanostructure, SnTe/CdTe, prepared by a cation exchange, exhibiting enhanced chemical stability.38 Other alternatives for shell formation can be considered. PbTe is an effective shell constituent, when its rock-salt crystal structure is compatible with that of SnTe with a close lattice matching. Crystallographic constants of SnTe and PbTe are 6.300 Å and 6.462 Å,40 respectively. In addition, the incorporation of PbTe and SnTe as a core/shell heterostructure induces additional tuning of the electronic structure beyond the size confinement41-44 due to the type-II alignment (See Figure S1 in the Supporting Information (SI)). SnTe is known as a ptype semiconductor owing to the high concentration of cation vacancies,41, 43 while PbTe can behave as p-type or n-type semiconductor according to its stoichiometry or surface properties. So, SnTe/PbTe should enable tunability of the electrical properties. This paper presents a unique procedure for the formation of a SnTe/PbTe core/shell nano-heterostructure with

ACS Paragon Plus Environment

Chemistry of Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

surplus coating of glassy-like SnO2 layer at the exterior surfaces. To the best of our knowledge, colloidal nanoheterostructures of that kind are reported here for the first time. The preparation of SnTe/PbTe/SnO2 nanostructures occurred via cation diffusion of Sn outward and Pb inward across an interface. The process is governed by the Kirkendall effect, taking place when diffusion rates of the exchanged elements are different (Sn and Pb in this case). The reaction stages were followed by applying energy-dispersive X-ray spectroscopy (EDX). A theoretical calculation verified a small diffusion distance of a Pb cation, which explains the formation of a PbTe shell as an intermediate layer; however, the larger diffusion of Sn toward the exterior surface ended in the formation of a SnO2 glassy exterior coating. Overall, the work illustrated a new procedure for the formation of SnTe-based colloidal nano-heterostructures, with a controlled surface stability and tunability of optical properties. Experimental section Chemicals and Materials. Tin chloride (SnCl2; 98%), lead(II) bromide (PbBr2; 99.999%), lead(II) oxide (PbO; 99.999%), tellurium (Te; 99.99%), 1,2-dodecanediol (DDD; 90%), 1-octadecene (ODE; technical grade, 90%), oleylamine (OLA; technical grade, 70%), oleic acid (OA; 90%) and nhexadecane (HDC; 99%) were purchased from Sigma-Aldrich. Trioctylphosphine (TOP; 97%) was purchased from Strem. Tetrachloroethylene (TCE; spectroscopic grade) was purchased from Merck. Toluene (analytical), hexane (analytical) and ethyl acetate (absolute) were purchased from Bio-Lab Ltd. Acetone (absolute) was purchased from Gadot. These chemicals were used without further purification. SnTe nanostructures synthesis. The SnTe nanostructures were synthesized by slight modification of the published procedure.38 SnCl2 (0.4 mmol), DDD (2 mmol) and OLA (8 mL) were degassed in a three-neck flask under vacuum for 1 hour. 1 mL stock solution prepared by mixing 0.4 mL of 1 M TOPTe and 0.6 mL of ODE was swiftly injected into the reaction mixture at 180 oC under nitrogen. After a specified time, the reaction vessel was cooled down and the mixture of toluene and OA (volume ratio of 4:1) was added into the crude solution to replace weakly bound OLA, and then the nanostructures were washed twice by acetone. Afterward, the nanostructures were re-dissolved in toluene for further use. SnTe/PbTe/SnO2 nano-heterostructures synthesis. PbBr2 (0.188 mmol) and ODE (5 mL) were dried in a three-neck flask under vacuum for 1 hour at 120 °C. OLA (0.5 mL) and OA (0.5 mL) were added at 120 oC under nitrogen atmosphere. After complete dissolution of PbBr2 at 150 °C, the vessel was cooled down to room temperature and SnTe nanostructure (24 mg) in toluene solution was added. After the solvent was evaporated under vacuum for a few minutes, the temperature was raised to 50 °C and the reaction vessel was kept at this temperature for a specified time. Afterward, the reaction was quenched and the product was isolated by centrifugation with addition of ethyl acetate. The resulting precipitate was redispersed in hexane or toluene for further use and characterization. Annealing treatment. The annealing of SnTe/PbTe/SnO2 nanostructures was performed in a glove box. A certain amount of nanostructure solution in hexane was dissolved in HDC after drying hexane. The solution was heated to 120-

Page 2 of 10

140 °C and maintained for 20 minutes. After an annealing process, the solution was isolated by centrifugation. The products were dissolved in hexane or toluene for further characterization. Characterization. Transmission electron microscopy (TEM) micrographs, high resolution TEM (HRTEM) micrographs, high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) micrographs, energy-dispersive X-ray spectroscopy (EDX) spectra and EDX maps were obtained using a Thermo Fisher (FEI) Double Cs-corrected Titan Themis G2 60-300 operated at 300 keV and equipped with a Bruker DualX detector. X-ray powder diffraction (XRD) patterns were obtained by Rigaku SmartLab diffractometer with Cu Kα radiation (λ=1.5418 Å). UV-VIS absorption spectra of the nanostructures were obtained using a JASCO V-570 UV-VIS-NIR spectrometer. Fourier Transform Infrared (FTIR) spectra were obtained using a Bruker FTIR spectrometer with a resolution of 4 cm-1. Raman spectra were obtained using a micro-Raman spectrometer (Horiba Jobin Yvon LabRam HR), equipped with Nd:YAG laser for excitation at 532 nm. Results and discussion The core SnTe nanostructures were synthesized using a modified procedure from that given in ref 38. The procedure supplied SnTe cores capped with oleic acid ligands. Then these SnTe cores were injected into a new solution containing PbBr2, oleic acid and oleylamine at a low temperature for the formation of a shell coating. A detailed synthetic procedure is given at the Experimental section. Representative transmission electron microscopy (TEM) images of SnTe cores before and after a shell coating are presented in Figure 1A and 1B, respectively, revealing a cube-like shape of the nanostructures, and size variation from 12.7 ± 1.2 nm to 14.3 ± 1.9 nm after coating. Figure 1C and 1D show high resolution TEM (HRTEM) images of the produced nanostructures, exhibiting a single crystal pattern with the dominant (200) lattice planes of SnTe structure. It is important to note that a fuzzy covering is seen around the core/shell structures in the TEM images, suggested below to be associated with formation of SnO2 coating. The fast Fourier transform (FFT) image of the nanostructures is displayed in Figure S2 in the SI, showing typical rings, assigned to the SnTe crystal planes. Figure 2 illustrates a high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) analysis with EDX measurement, revealing the elemental distribution across a nanostructure. Figure 2A displays a HAADF-STEM image, consisting of three different contrast regions (apart from the darkest contrast in the background which is the amorphous carbon supporting film); the darkest part is observed in the external shell, while the brightest region is located at an intermediate layer. The EDX mapping images for Sn and Pb elements are presented in Figure 2B and 2C. The map in Figure 2B suggests that the Sn element occupies nearly the entire core, despite a pronounced vacancy in the shape of a ring below the surface boundary. In addition, the Sn mapping indicates the existence of a fuzzy region in exterior periphery, wrapping the vacant ring. The map in panel 2C shows that Pb occupies only a ring region, but it is absent elsewhere. Panel 2D displays the Te mapping, covering nearly the entire structure except across the fuzzy region. Figure 2E presents an overlap between the maps of Sn and Pb elements, distinctively exhibiting a match between the Sn-deficient layer and the Pb


Page 3 of 10 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


ring, proving the occurrence of cation exchange of Sn2+ with Pb2+. Figure 2F represents the overlay of Pb and Te distributions, revealing the formation of PbTe at an inner ring below the surface. The similarity in crystallographic parameters between bulk SnTe and PbTe compounds, makes an exchange between Sn2+ and Pb2+ feasible. Overall, the images shown in Figure 2 validate the formation of the SnTe core covered by a PbTe ring shell. Figure S3 in the SI displays EDX mapping image of oxygen element, revealing its existence at the fuzzy region, and suggesting the formation of a SnO2 glassy layer. The oxygen atoms presumably originate from oleic acid molecules which are acting as surface ligands. The external SnO2 layer's thickness is in close agreement with the nanostructure's size increase after a shell coating. It should be noted that a presence of Br element was excluded by EDX mapping. The observations finally reveal the formation of SnTe/PbTe/SnO2 core/shell/shell nanostructures. The mechanism for the formation of the structure is discussed further on. The EDX mapping images were acquired after keeping the nanostructures in ambient conditions for more than one month, as shown in Figure S4 in the SI. The images in Figure S4 suggest that the elemental distribution is nearly maintained in comparison to that shown in Figure 2, reflecting the achievement of chemical stability.

Figure 1. TEM images of (A) SnTe and (B) SnTe/PbTe/SnO2 nanostructures, as well as (C, D) HRTEM images of SnTe/PbTe/SnO2 nano-heterostructures.

Figure 2. (A) HAADF-STEM image and (B-F) corresponding EDX mapping images of the following elements in a SnTe/PbTe/SnO2 nano-heterostructure: (B) Sn, (C) Pb, (D) Te, (E) Sn and Pb and (F) Pb and Te.

An EDX spectrum of the nano-heterostructures is presented in Figure 3A, illustrating dominancy of Sn constituent, together with the presence of Pb atoms. The atomic ratios of Sn, Pb and Te elements were estimated to be 44.3%, 8.6% and 47.1%, respectively, indicating incorporation of small amounts of Pb in the nano-heterostructures and low toxic Sn element's occupation of ~84% in the cations' positions. Only a small remnant of Br atoms was also observed, which may be related to their presence at the surface surrounding, although it was not pronounced in the STEM mapping. X-ray powder diffraction (XRD) pattern of SnTe/PbTe/SnO2 nano-heterostructures is shown in Figure 3B, elucidating the cubic rock-salt SnTe of the core structure. The main diffraction peaks are indexed as the (200), (220), (222), (400), (420) and (422) planes of the rock-salt SnTe, without obvious distinction for the SnTe/PbTe boundary. A few other weak peaks are assigned to SnO2 and Sn, associated with their low crystallinity. The XRD spectrum lacks evidence of the existence of PbBr2, PbO or TeO2 compounds. Figure 3C represents the Raman spectrum of the core/shell/shell nano-heterostructures. Two main peaks at 100150 cm-1 in the spectrum correspond to the longitudinal optical (LO) and transversal optical (TO) modes of SnTe.45, 46 Since the LO and the TO modes of PbTe are very close to those of SnTe, they cannot be resolved. The weak peaks at 277 and 495 cm-1 seem to correspond to Raman modes of SnO2.47-49 The Raman spectrum lacks evidence for TeO2 modes (which according to the literature were supposed to appear at 390 and 650 cm-1),50 in agreement with the XRD measurement. Absorption spectra of two SnTe cores of different sizes and their corresponding SnTe/PbTe/SnO2 nano-heterostructures, are presented in Figure 3D. The spectrum of the pristine cores is characterized by an absorption edge and additional week features above 1500 nm. The few sharp peaks at 1700 nm and 2300 nm are related to vibrational modes of the organic ligands. Figure S5 at the SI displays an absorbance curve of pristine oleic acid solution (without nano-heterostructures), corroborating the molecules' vibration modes. The absorption edge of the core/shell/shell nano-heterostructures is slightly shifted to the red side with respect to that of the pristine cores, a shift that can be associated with the quasi-type-II alignment at the core/shell interface. Figure 3E displays Fourier transform infrared (FTIR) spectra of the samples shown in Figure



3C. The spectra are composed of broad envelopes and sharp lines. The sharp lines were identified as ligands' vibration modes. The broad envelope may be related to a Plasmon or to an intra-band transition in semiconductors with a low band gap energy. The exact identification of those envelopes should be further examined in the future. However, it seems from Figure 3E that the envelope of the core/shell/shell sample is red-shifted from envelope of the pristine cores, and presumably related to the change of electronic structure from a type-I to a type-II when moving from the core to the core/shell structure. A stability test was carried out to examine the chemical solidity of the nanostructures. Surprisingly, the absorption spectrum of the SnTe/PbTe/SnO2 core/shell/shell nanoheterostructures remained nearly intact after air exposure for up to 60 days, as shown in Figure 3F.

Figure 3. (A) EDX, (B) XRD and (C) Raman spectra of SnTe/PbTe/SnO2 nano-heterostructures. (D) Absorption spectra of two SnTe cores of different sizes (core 1: 8.1 ± 0.7 nm and core 2: 10.5 ± 0.9 nm) and their corresponding SnTe/PbTe/SnO2 nanoheterostructures. (E) FTIR spectra of SnTe/PbTe/SnO2 nanoheterostructures. (F) Absorption spectra of the nanoheterostructures after storage time as indicated in the legend.

A few control experiments were carried out for learning the influence of synthesis conditions on the final core/shell/shell product. Figure 4A displays a set of EDX images related to nanostructures formed under different reaction times, elucidating a slight increase of the PbTe ring's thickness, but a thin out of the SnO2 layer, with the increase of the time duration at a fixed temperature of 50 °C. The diminution of the SnO2 shell may be related to a gradual evaporation of the oleic acid ligands, thereby a reduction of an oxygen source. A control reac-

Page 4 of 10

tion prepared at room temperature (Figure 4B and Figure S6 in the SI) led to the formation of an extremely thin PbTe layer. Hence, temperature and reaction duration can regulate the inward penetration of Pb2+ ions. The influence of precursor type was examined by replacing PbBr2 precursor with PbO (Figure 4C and Figure S7 in the SI). Examination of the relevant EDX images suggests achievement of similar results with both precursors. The independent behavior on the starting materials further supports the lack of involvement of Br ions in the chemical reaction.

Figure 4. EDX mapping images of SnTe/PbTe/SnO2 nanoheterostructure obtained at under different conditions: (A) a reaction time, (B) the reaction temperature and (C) the type of Pb precursor.

Next we examined the influence of a post annealing treatment of the SnTe/PbTe/SnO2 nano-heterostructures. According to previous studies, it was anticipated that annealing might accelerate cation diffusion processes,51 or/and a formation of alloying regions.51-53 In the present case, annealing was performed under an inert atmosphere at temperature range of 120140 °C (see Experimental section). Representative HAADFSTEM image and EDX mapping images of the nano-


Page 5 of 10 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


heterostructures after annealing are presented in Figure 5. Figure 5A displays a HAADF-STEM image with an obvious modification compared to that shown in Figure 2A, suggesting re-arrangement of elemental distribution. Figure 5B to 5F demonstrate the EDX mapping of the individual elements or their combination. The observations shown in Figure 5B and 5C reveal the occurrence of extended penetration of Pb2+ ions along a selective crystallographic axis with respect to the nonannealed samples, and a spread of Sn deficiency region. As expected, the Te mapping hardly changed after the annealing treatment (Figure 5D), due to the extremely low diffusion constant of this element even at elevated temperatures, and with respect to the fast diffusion of the cations (see a table and a discussion below).38 The overlay images (Figure 5E and 5F) further confirm the exchange of occupation between Pb and Sn, where Sn diffused outward from the central part of the core to the exterior region and Pb occupied an inner deficient region. It is worth mentioning that most particles of the annealed samples showed such re-arrangement and an additional example is supplied in Figure S8 at the SI. The TEM image of the nano-heterostructures after an annealing stage is displayed at the SI, Figure S9A, while the corresponding Fourier filtering of Figure S9A is shown in Figure S9B. The mentioned observations imply that diffusion process at higher temperature is relatively fast at a specific facet,54-57 in contradiction with a diffusion along other directions.

Figure 5. (A) HAADF-STEM image and (B-F) corresponding EDX mapping images of the following elements in SnTe/PbTe/SnO2 nano-heterostructure and after annealing at 140 °C: (B) Sn, (C) Pb, (D) Te, (E) Sn and Pb and (F) Pb and Te.

The diffusion coefficient (D) and diffusion length (L) at two reaction temperatures (i.e., 50 °C and 140 °C) were estimated by employing an Arrhenius equation (eq. 1) and a root-meansquare distance (eq. 2), respectively.58-60 The diffusion coefficient and length's values are given in Table 1.

(1)

√2 (2) D0, E, k, T and t are the pre-exponential factor, the activation energy, the Boltzmann constant, temperature and time, respectively. Self-diffusion constants (D0 and E) were adopted from reference 58, and are supplied in Table S1 at the SI. Since these values are typically reported in the literature at high temperature ranges, extrapolation was performed for the working temperatures in the present case.59 Careful view of Table 1 indicates that the diffusion rates and distances of the Sn2+ and Pb2+ ions at 50 °C differ only by a

factor of ~3 with the same order of magnitude. Therefore, very mild exchange between SnTe and PbTe at 50 °C was observed under extension of the reaction duration (see Figure 4A), pronounced as a minor grows of the PbTe thickness, however, by a gradual reduction of the SnO2 thickness. But, when the reaction temperature increased, some voids in the core were observed (see Figure 5B) due to faster diffusion process of both Sn and Pb (see Table 1). On the other hand, diffusion rate of Te was much slower than those of the two cations due to large ionic radius of Te even at the elevated temperatures, obviously indicating a topotactic reaction, where the large anions preserve the crystallographic frame.55, 61-64 Worth noting that a self-diffusion rate of Te may vary from one host kind to another, however, here we found a compatible diffusion rate to that found in other similar matrixes (e.g., ~6.1×10-13 nm2/s in CdTe65 and ~1.38×10-8 nm2/s in PbTe66). Table 1. Diffusion rates and lengths of the Sn, Te and Pb elements at 50 °C and 140 °C 50 °C

Sn 2

Te -3

D [nm /s]

3.6 × 10

Pb -11

1.1 × 10-3

-4

8.1 × 10

L for 10 min [nm]

3.6

3.1 × 10

2.0

140 °C

Sn

Te

Pb

2

-5

D [nm /s]

21.3

1.3 × 10

4.2

L for 10 min [nm]

160

0.1

71

The inward motion of Pb2+ is kinetically controlled and depends predominantly on the hopping mechanism from one void to another (substitution or interstitial sites). The kinetic process described is most efficient along a selective (111) direction (an axis normal to the {111} planes). The efficiency along the selective direction can be explained with aid of the schemes shown in Figure 6. The rock salt SnTe exhibits AB (A=B) arrangement of Sn cations along (111) direction, with high density,67, 68 as shown on the right scheme in Figure 6. The left illustration of Figure 6 displays a unit cell, in which the Sn positions are marked by the ascending numbers (1-6) with one vacancy at position 6. In general, self-diffusion of Sn-ion can take place among the 6 positions of one plane A, and even toward the adjacent plane B. This self-diffusion process acts as a barrier stage to a fast outward migration. However, a thermal activation process at 140 °C, overcomes the barriers and permit outward motion from inner planes toward the surface. This assumption goes along with the calculated diffusion length of ~160 nm (see Table 1). A careful view of the HAADF-STEM images (see Figure 5) indicates existence of voids along the (111) direction. In parallel, the incoming Pb-cations having compatible ionic radius with Sn and extended diffusion length at 140 °C (see Table 1), can statistically occupy Sn-vacancies in the (111) plane and block the Sn ions' self-diffusion. At this point, creation of a PbTe layer onto SnTe is more favorable at a certain plane with closer crystallographic matching54-56 (according to eq. 3) and with a low formation energy of a PbTe coating. m

-

.

(3)

where, m is a commensurate lattice mismatch at an interface with two different lattice spacings (d1 and d2). The calculation confirms that the lattice mismatch along the {111} facets between SnTe and PbTe (1.62%) is smaller than that along the



{100} and {110} facets (2.19% and 1.77%, respectively). Furthermore, in the rock-salt structures, {111} surfaces typically present higher surface energy than that of {100} facets,69, 70 offering high reactivity. The arguments discussed above and the illustration given in Figure 6 imply that the cation exchange is mainly driven by a kinetic process (self-diffusion) at interfaces of the {111} as well as combined with a thermodynamic control by the low PbTe formation energy.

Page 6 of 10

pm,89 respectively) are very similar without a preference for occupation of interstitial sites. In current nanostructures, a PbTe layer is located between SnTe in the core and SnO2 in the exterior periphery, thus a SnO2 shell layer will not interrupt the functionality of the interior components, e.g., core/shell type-II characteristic of the SnTe/PbTe pair. However, the large barrier created by SnO2 (see a band diagram of Figure S1 in the SI), could in fact be an obstacle for charge extraction for a solar cell, but, on the other hand, could be an excellent carriers' confining layer, with benefits in LEDs, lasers and photo-detectors with activity in the IR spectral regime. Scheme 1. Schematic illustration of the formation mechanism of SnTe/PbTe/SnO2 nano-heterostructures.

Figure 6. Schematic illustrations showing: Left: a diffusion process of Pb ions (red spheres) into a vacancy of Sn ion (green spheres) at the (111) plane; Right: close-packing of Sn or Pb ions into AB planes (including one exemplary vacancy site).

Following the discussion above, the mechanism which takes place at low temperatures is illustrated schematically in Scheme 1. A diffusion process across an interface of elements with different diffusion rates is known as a Kirkendall effect.71-74 This effect has already been discussed in a few cases which presented hollowed cages in nanostructures,60, 71, 75-79 but most recently has also been mentioned in correlation with the formation of inverted core/shell nano-heterostructures.38, 80 On a sub-micron level, the Kirkendall effect may follow one of the following diffusion mechanisms: (a) Direct vacancy mechanism, where vacancies are displaced from one crystal side to another; (b) kick-out mechanism,81-83 in which foreign impurities occupy interstitial positions, moving from one to another and eventually displacing a lattice atom. Consequently, the dislodged host atoms may become self-interstitial while the impurity atoms become immobile; (c) Frank-Turnbull mechanism,83, 84 involving free motion of impurities from one interstitial side to another until they get trapped at vacancy positions, whereupon they are almost immobile. The mechanism that supports the formation of SnTe/PbTe/SnO2 nanoheterostructures encompasses a combination of diffusion processes: In one scenario (i) Sn atoms diffuse outward from a distance of no more than 3.6 nm below the surface due to heating at 50 °C (a direct vacancy mechanism), leaving behind a deficiency layer. This is followed by penetration of Pb atoms via the Frank-Turnbull mechanism into the existing vacancies and becoming immobilized. It is important to note that SnTe is a p-type semiconductor, associated with the creation of Sn vacancy with a negative formation energy.41 This inherent characteristic enabled in the past the incorporation of foreign cations (e.g., Mn and Cd) into vacancy sites, and accordingly showed a pronounced increase of a Seebeck coefficient with a benefit for thermoelectric applications.42, 85-88 In a second scenario (ii), Pb atoms diffuse into interstitial sites and after random motion through such sites, they finally displace Sn atoms from a crystal site. Scenario (ii) can probably be excluded due to the fact that ionic radii of Pb2+ and Sn2+ (~119 and 118

Conclusion In summary, the present work demonstrates the preparation of SnTe/PbTe/SnO2 core/shell/shell nano-heterostructures by mixing Pb precursors with an isolated SnTe core at a low temperature. Careful characterization revealed that the growth process of PbTe layer is based on nanoscale cation diffusion, called the Kirkendall effect. The outward diffusion of Sn atoms generates a deficient region, while the inward diffusion of Pb induces the formation of a PbTe layer. The ejected Sn atoms at the surface finally form a glassy SnO2 layer owing to surface oxidation. The theoretical calculation confirmed unequal diffusion rate and distances between Sn and Pb, leading to limited penetration of Pb into the core. The examination after 60 days of air exposure verified improved chemical stability of the nano-heterostructures. This work demonstrates a facile route for preparing infra-red active colloidal nanostructures and provides a new insight into interface chemistry for achieving tunable electronic structures, with chemical and spectral stability.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publications website at DOI:xx.xx Band structure, FFT image, TEM images, HAADF-STEM images and EDX mapping images of SnTe/PbTe/SnO2 nanostructure, absorption spectrum and table of diffusion variables.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]

Author Contributions ‡A.S. and Y.J. contributed equally.

Notes The authors declare no competing ﬁnancial interest.

ACKNOWLEDGMENT ACS Paragon Plus Environment

Page 7 of 10 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


This work was supported by the Israel Council for Higher Education-Focal Area Technology (Project No. 872967), the Volkswagen Stiftung (Project No. 88116), the Israel Ministry of Defense (Project No. 4440827018), the Israel Ministry of Trade (Maymad Project No. 54662), the Israel Science Foundation Bikura (Project No. 1508/14), the Israel Science Foundation (Project No. 985/11 and 914/15), the Niedersachsen-Deutsche Technion Gesellschaft E.V. (Project No. ZN2916) and the European Commission via the Marie-Sklodowska Curie action Phonsi (Project No. H2020-MSCAITN-642656).

REFERENCES (1) Kagan, C. R.; Lifshitz, E.; Sargent, E. H.; Talapin, D. V., Building Devices from Colloidal Quantum Dots. Science 2016, 353, aac5523. (2) Smith, A. M.; Nie, S. M., Semiconductor Nanocrystals: Structure, Properties, and Band Gap Engineering. Acc. Chem. Res. 2010, 43, 190-200. (3) Talapin, D. V.; Lee, J.-S.; Kovalenko, M. V.; Shevchenko, E. V., Prospects of Colloidal Nanocrystals for Electronic and Optoelectronic Applications. Chem. Rev. 2010, 110, 389-458. (4) McDonald, S. A.; Konstantatos, G.; Zhang, S.; Cyr, P. W.; Klem, E. J.; Levina, L.; Sargent, E. H., Solution-Processed PbS Quantum Dot Infrared Photodetectors and Photovoltaics. Nat. Mater. 2005, 4, 138-142. (5) Nozik, A. J.; Beard, M. C.; Luther, J. M.; Law, M.; Ellingson, R. J.; Johnson, J. C., Semiconductor Quantum Dots and Quantum Dot Arrays and Applications of Multiple Exciton Generation to Third-Generation Photovoltaic Solar Cells. Chem. Rev. 2010, 110, 6873-6890. (6) Tang, J.; Sargent, E. H., Infrared Colloidal Quantum Dots for Photovoltaics: Fundamentals and Recent Progress. Adv. Mater. 2011, 23, 12-29. (7) Chuang, C. H.; Brown, P. R.; Bulovic, V.; Bawendi, M. G., Improved Performance and Stability in Quantum Dot Solar Cells through Band Alignment Engineering. Nat. Mater. 2014, 13, 796-801. (8) Zhang, Z.; Yang, J.; Wen, X.; Yuan, L.; Shrestha, S.; Stride, J. A.; Conibeer, G. J.; Patterson, R. J.; Huang, S., Effect of Halide Treatments on PbSe Quantum Dot Thin Films: Stability, Hot Carrier Lifetime, and Application to Photovoltaics. J. Phys. Chem. C 2015, 119, 24149-24155. (9) Kim, S.; Marshall, A. R.; Kroupa, D. M.; Miller, E. M.; Luther, J. M.; Jeong, S.; Beard, M. C., Air-Stable and Efficient PbSe Quantum-Dot Solar Cells Based upon ZnSe to PbSe CationExchanged Quantum Dots. ACS Nano 2015, 9, 8157-8164. (10) Bruchez, M.; Moronne, M.; Gin, P.; Weiss, S.; Alivisatos, A. P., Semiconductor Nanocrystals as Fluorescent Biological Labels. Science 1998, 281, 2013-2016. (11) Chan, W. C. W.; Nie, S., Quantum Dot Bioconjugates for Ultrasensitive Nonisotopic Detection. Science 1998, 281, 2016-2018. (12) Han, M. Y.; Gao, X. H.; Su, J. Z.; Nie, S., Quantum-DotTagged Microbeads for Multiplexed Optical Coding of Biomolecules. Nat. Biotechnol. 2001, 19, 631-635. (13) Harman, T. C.; Taylor, P. J.; Walsh, M. P.; LaForge, B. E., Quantum Dot Superlattice Thermoelectric Materials and Devices. Science 2002, 297, 2229-2232. (14) Heremans, J. P.; Thrush, C. M.; Morelli, D. T., Thermopower Enhancement in Lead Telluride Nanostructures. Phys. Rev. B 2004, 70, 115334. (15) Kovalenko, M. V.; Spokoyny, B.; Lee, J. S.; Scheele, M.; Weber, A.; Perera, S.; Landry, D.; Talapin, D. V., Semiconductor Nanocrystals Functionalized with Antimony Telluride Zintl Ions for Nanostructured Thermoelectrics. J. Am. Chem. Soc. 2010, 132, 66866695. (16) Scheele, M.; Oeschler, N.; Veremchuk, I.; Peters, S.-O.; Littig, A.; Kornowski, A.; Klinke, C.; Weller, H., Thermoelectric Properties of Lead Chalcogenide Core–Shell Nanostructures. ACS Nano 2011, 5, 8541-8551.

(17) Colvin, V. L.; Schlamp, M. C.; Alivisatos, A. P., LightEmitting Diodes Made from Cadmium Selenide Nanocrystals and a Semiconducting Polymer. Nature 1994, 370, 354-357. (18) Schlamp, M. C.; Peng, X. G.; Alivisatos, A. P., Improved Efficiencies in Light Emitting Diodes Made with CdSe(CdS) Core/Shell Type Nanocrystals and a Semiconducting Polymer. J. Appl. Phys. 1997, 82, 5837-5842. (19) Coe, S.; Woo, W. K.; Bawendi, M.; Bulovic, V., Electroluminescence from Single Monolayers of Nanocrystals in Molecular Organic Devices. Nature 2002, 420, 800-803. (20) Achermann, M.; Petruska, M. A.; Koleske, D. D.; Crawford, M. H.; Klimov, V. I., Nanocrystal-Based Light-Emitting Diodes Utilizing High-Efficiency Nonradiative Energy Transfer for Color Conversion. Nano Lett. 2006, 6, 1396-1400. (21) Mueller, A. H.; Petruska, M. A.; Achermann, M.; Werder, D. J.; Akhadov, E. A.; Koleske, D. D.; Hoffbauer, M. A.; Klimov, V. I., Multicolor Light-Emitting Diodes Based on Semiconductor Nanocrystals Encapsulated in GaN Charge Injection Layers. Nano Lett. 2005, 5, 1039-1044. (22) Kim, S.; Kim, T.; Kang, M.; Kwak, S. K.; Yoo, T. W.; Park, L. S.; Yang, I.; Hwang, S.; Lee, J. E.; Kim, S. K.; Kim, S. W., Highly Luminescent InP/GaP/ZnS Nanocrystals and Their Application to White Light-Emitting Diodes. J. Am. Chem. Soc. 2012, 134, 3804-3809. (23) Hines, D. A.; Kamat, P. V., Recent Advances in Quantum Dot Surface Chemistry. ACS Appl. Mater. Interfaces 2014, 6, 30413057. (24) Krause, M. M.; Kambhampati, P., Linking Surface Chemistry to Optical Properties of Semiconductor Nanocrystals. Phys. Chem. Chem. Phys. 2015, 17, 18882-18894. (25) Boles, M. A.; Ling, D.; Hyeon, T.; Talapin, D. V., The Surface Science of Nanocrystals. Nat. Mater. 2016, 15, 141-153. (26) Jang, Y.; Shapiro, A.; Isarov, M.; Rubin-Brusilovski, A.; Safran, A.; Budniak, A. K.; Horani, F.; Dehnel, J.; Sashchiuk, A.; Lifshitz, E., Interface Control of Electronic and Optical Properties in IV-VI and II-VI Core/Shell Colloidal Quantum Dots: a Review. Chem. Commun. 2017, 53, 1002-1024. (27) Cao; Banin, U., Growth and Properties of Semiconductor Core/Shell Nanocrystals with InAs Cores. J. Am. Chem. Soc. 2000, 122, 9692-9702. (28) Xie, R. G.; Kolb, U.; Li, J. X.; Basche, T.; Mews, A., Synthesis and Characterization of Highly Luminescent CdSe-Core CdS/Zn0.5Cd0.5S/ZnS Multishell Nanocrystals. J. Am. Chem. Soc. 2005, 127, 7480-7488. (29) Peng, X. G.; Schlamp, M. C.; Kadavanich, A. V.; Alivisatos, A. P., Epitaxial Growth of Highly Luminescent CdSe/CdS Core/Shell Nanocrystals with Photostability and Electronic Accessibility. J. Am. Chem. Soc. 1997, 119, 7019-7029. (30) Chin, P. T.; Donega Cde, M.; van Bavel, S. S.; Meskers, S. C.; Sommerdijk, N. A.; Janssen, R. A., Highly Luminescent CdTe/CdSe Colloidal Heteronanocrystals with TemperatureDependent Emission Color. J. Am. Chem. Soc. 2007, 129, 1488014886. (31) Dorfs, D.; Franzl, T.; Osovsky, R.; Brumer, M.; Lifshitz, E.; Klar, T. A.; Eychmüller, A., Type-I and Type-II Nanoscale Heterostructures Based on CdTe Nanocrystals: A Comparative Study. Small 2008, 4, 1148-1152. (32) Chen, O.; Zhao, J.; Chauhan, V. P.; Cui, J.; Wong, C.; Harris, D. K.; Wei, H.; Han, H. S.; Fukumura, D.; Jain, R. K.; Bawendi, M. G., Compact High-Quality CdSe-CdS Core-Shell Nanocrystals with Narrow Emission Linewidths and Suppressed Blinking. Nat. Mater. 2013, 12, 445-451. (33) Kovalenko, M. V.; Heiss, W.; Shevchenko, E. V.; Lee, J.S.; Schwinghammer, H.; Alivisatos, A. P.; Talapin, D. V., SnTe Nanocrystals: A New Example of Narrow-Gap Semiconductor Quantum Dots. J. Am. Chem. Soc. 2007, 129, 11354-11355. (34) Xu, Y.; Al-Salim, N.; Hodgkiss, J. M.; Tilley, R. D., Solution Synthesis and Optical Properties of SnTe Nanocrystals. Cryst. Growth Des. 2011, 11, 2721-2723. (35) Guo, S.; Fidler, A. F.; He, K.; Su, D.; Chen, G.; Lin, Q.; Pietryga, J. M.; Klimov, V. I., Shape-Controlled Narrow-Gap SnTe



Nanostructures: From Nanocubes to Nanorods and Nanowires. J. Am. Chem. Soc. 2015, 137, 15074-15077. (36) Reiss, P.; Carriere, M.; Lincheneau, C.; Vaure, L.; Tamang, S., Synthesis of Semiconductor Nanocrystals, Focusing on Nontoxic and Earth-Abundant Materials. Chem. Rev. 2016, 116, 10731-10819. (37) Smith, P. J., Chemistry of Tin. Springer Netherlands: 2012. (38) Jang, Y.; Yanover, D.; Capek, R. K.; Shapiro, A.; Grumbach, N.; Kauffmann, Y.; Sashchiuk, A.; Lifshitz, E., Cation Exchange Combined with Kirkendall Effect in the Preparation of SnTe/CdTe and CdTe/SnTe Core/Shell Nanocrystals. J. Phys. Chem. Lett. 2016, 7, 2602-2609. (39) de Kergommeaux, A.; Faure-Vincent, J.; Pron, A.; de Bettignies, R.; Malaman, B.; Reiss, P., Surface Oxidation of Tin Chalcogenide Nanocrystals Revealed by 119Sn-Mössbauer Spectroscopy. J. Am. Chem. Soc. 2012, 134, 11659-11666. (40) Reiss, P.; Protière, M.; Li, L., Core/Shell Semiconductor Nanocrystals. Small 2009, 5, 154-168. (41) Wang, N.; West, D.; Liu, J.; Li, J.; Yan, Q.; Gu, B.-L.; Zhang, S. B.; Duan, W., Microscopic Origin of The p-type Conductivity of the Topological Crystalline Insulator SnTe and the Effect of Pb Alloying. Phys. Rev. B 2014, 89, 045142. (42) Tan, G.; Shi, F.; Hao, S.; Chi, H.; Zhao, L. D.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G., Codoping in SnTe: Enhancement of Thermoelectric Performance through Synergy of Resonance Levels and Band Convergence. J. Am. Chem. Soc. 2015, 137, 5100-5112. (43) Rogacheva, E. I.; Grigorov, S. N.; Nashchekina, O. N.; Tavrina, T. V.; Lyubchenko, S. G.; Sipatov, A. Y.; Volobuev, V. V.; Fedorov, A. G.; Dresselhaus, M. S., Growth Mechanism and Thermoelectric Properties of PbTe/SnTe/PbTe Heterostructures. Thin Solid Films 2005, 493, 41-48. (44) Rogacheva, E. I.; Nashchekina, O. N.; Meriuts, A. V.; Lyubchenko, S. G.; Dresselhaus, M. S.; Dresselhaus, G., Quantum Size Effects in n-PbTe⁄p-SnTe⁄n-PbTe Heterostructures. Appl. Phys. Lett. 2005, 86, 063103. (45) Sugai, S.; Murase, K.; Kawamura, H., Observation of Soft TO-Phonon in SnTe by Raman Scattering. Solid State Commun. 1977, 23, 127-129. (46) Shimada, T.; Kobayashi, K. L. I.; Katayama, Y.; Komatsubara, K. F., Soft-Phonon-Induced Raman Scattering in IV-VI Compounds. Phys. Rev. Lett. 1977, 39, 143-146. (47) Katiyar, R. S.; Dawson, P.; Hargreave, M. M.; Wilkinson, G. R., Dynamics of the rutile structure. III. Lattice dynamics, infrared and Raman spectra of SnO2. J. Phys. C: Solid State Phys. 1971, 4, 2421-2431. (48) Diéguez, A.; Romano-Rodrı́guez, A.; Vilà, A.; Morante, J. R., The complete Raman spectrum of nanometric SnO2 particles. J. Appl. Phys. 2001, 90, 1550-1557. (49) Peng, X. S.; Zhang, L. D.; Meng, G. W.; Tian, Y. T.; Lin, Y.; Geng, B. Y.; Sun, S. H., Micro-Raman and Infrared Properties of SnO2 Nanobelts Synthesized from Sn and SiO2 Powders. J. Appl. Phys. 2003, 93, 1760-1763. (50) Tschirner, N.; Lange, H.; Lambert, K.; Hens, Z.; Thomsen, C., Raman Spectroscopy of PbTe/CdTe Nanocrystals. Phys Status Solidi B 2011, 248, 2748-2750. (51) Isarov, M.; Grumbach, N.; Maikov, G. I.; Tilchin, J.; Jang, Y.; Sashchiuk, A.; Lifshitz, E., The Effect of Low Temperature Coating and Annealing on Structural and Optical Properties of CdSe/CdS Core/Shell QDs. Lith. J. Phys. 2015, 55, 297-304. (52) Grumbach, N.; Capek, R. K.; Tilchin, E.; RubinBrusilovski, A.; Yang, J.; Ein-Eli, Y.; Lifshitz, E., Comprehensive Route to the Formation of Alloy Interface in Core/Shell Colloidal Quantum Dots. J. Phys. Chem. C 2015, 119, 12749-12756. (53) Rubin-Brusilovski, A.; Jang, Y.; Shapiro, A.; Safran, A.; Sashchiuk, A.; Lifshitz, E., Influence of Interfacial Strain on Optical Properties of PbSe/PbS Colloidal Quantum Dots. Chem. Mater. 2016, 28, 9056-9063. (54) Sadtler, B.; Demchenko, D. O.; Zheng, H.; Hughes, S. M.; Merkle, M. G.; Dahmen, U.; Wang, L.-W.; Alivisatos, A. P., Selective Facet Reactivity during Cation Exchange in Cadmium Sulfide Nanorods. J. Am. Chem. Soc. 2009, 131, 5285-5293.

Page 8 of 10

(55) Miszta, K.; Dorfs, D.; Genovese, A.; Kim, M. R.; Manna, L., Cation Exchange Reactions in Colloidal Branched Nanocrystals. ACS Nano 2011, 5, 7176-7183. (56) Casavola, M.; van Huis, M. A.; Bals, S.; Lambert, K.; Hens, Z.; Vanmaekelbergh, D., Anisotropic Cation Exchange in PbSe/CdSe Core/Shell Nanocrystals of Different Geometry. Chem. Mater. 2012, 24, 294-302. (57) Zaiats, G.; Shapiro, A.; Yanover, D.; Kauffmann, Y.; Sashchiuk, A.; Lifshitz, E., Optical and Electronic Properties of Nonconcentric PbSe/CdSe Colloidal Quantum Dots. J. Phys. Chem. Lett. 2015, 6, 2444-2448. (58) Mehrer, H.; Stolica, N.; Stolwijk, N., Diffusion in Solid Metals and Alloys. Springer Berlin Heidelberg: 1990. (59) Mokari, T.; Aharoni, A.; Popov, I.; Banin, U., Diffusion of Gold into InAs Nanocrystals. Angew. Chem., Int. Ed. 2006, 45, 80018005. (60) Cabot, A.; Smith, R. K.; Yin, Y.; Zheng, H.; Reinhard, B. M.; Liu, H.; Alivisatos, A. P., Sulfidation of Cadmium at the Nanoscale. ACS Nano 2008, 2, 1452-1458. (61) Son, D. H.; Hughes, S. M.; Yin, Y.; Alivisatos, A. P., Cation Exchange Reactions in Ionic Nanocrystals. Science 2004, 306, 1009-1012. (62) Luther, J. M.; Zheng, H.; Sadtler, B.; Alivisatos, A. P., Synthesis of PbS Nanorods and Other Ionic Nanocrystals of Complex Morphology by Sequential Cation Exchange Reactions. J. Am. Chem. Soc. 2009, 131, 16851-16857. (63) Jain, P. K.; Amirav, L.; Aloni, S.; Alivisatos, A. P., Nanoheterostructure Cation Exchange: Anionic Framework Conservation. J. Am. Chem. Soc. 2010, 132, 9997-9999. (64) Li, H.; Zanella, M.; Genovese, A.; Povia, M.; Falqui, A.; Giannini, C.; Manna, L., Sequential Cation Exchange in Nanocrystals: Preservation of Crystal Phase and Formation of Metastable Phases. Nano Lett. 2011, 11, 4964-4970. (65) Borsenberger, P. M.; Stevenson, D. A., Self-Diffusion of Cadmium and Tellurium in Cadmium Telluride. J. Phys. Chem. Solids 1968, 29, 1277-1286. (66) Gomez, M. P.; Stevenson, D. A.; Huggins, R. A., SelfDiffusion of Pb and Te in Lead Telluride. J. Phys. Chem. Solids 1971, 32, 335-344. (67) Cho, K.-S.; Talapin, D. V.; Gaschler, W.; Murray, C. B., Designing PbSe Nanowires and Nanorings through Oriented Attachment of Nanoparticles. J. Am. Chem. Soc. 2005, 127, 7140-7147. (68) Woo, J. Y.; Ko, J.-H.; Song, J. H.; Kim, K.; Choi, H.; Kim, Y.-H.; Lee, D. C.; Jeong, S., Ultrastable PbSe Nanocrystal Quantum Dots via in Situ Formation of Atomically Thin Halide Adlayers on PbSe(100). J. Am. Chem. Soc. 2014, 136, 8883-8886. (69) Lee, S. M.; Cho, S. N.; Cheon, J., Anisotropic Shape Control of Colloidal Inorganic Nanocrystals. Adv. Mater. 2003, 15, 441444. (70) Rubin-Brusilovski, A.; Maikov, G.; Kolan, D.; Vaxenburg, R.; Tilchin, J.; Kauffmann, Y.; Sashchiuk, A.; Lifshitz, E., Influence of Alloying on the Optical Properties of IV–VI Nanorods. J. Phys. Chem. C 2012, 116, 18983-18989. (71) Yin, Y.; Rioux, R. M.; Erdonmez, C. K.; Hughes, S.; Somorjai, G. A.; Alivisatos, A. P., Formation of Hollow Nanocrystals Through the Nanoscale Kirkendall Effect. Science 2004, 304, 711714. (72) Wang, W.; Dahl, M.; Yin, Y., Hollow Nanocrystals through the Nanoscale Kirkendall Effect. Chem. Mater. 2013, 25, 1179-1189. (73) Anderson, B. D.; Tracy, J. B., Nanoparticle Conversion Chemistry: Kirkendall Effect, Galvanic Exchange, and Anion Exchange. Nanoscale 2014, 6, 12195-12216. (74) Yang, Z.; Yang, N.; Pileni, M.-P., Nano Kirkendall Effect Related to Nanocrystallinity of Metal Nanocrystals: Influence of the Outward and Inward Atomic Diffusion on the Final Nanoparticle Structure. J. Phys. Chem. C 2015, 119, 22249-22260. (75) Railsback, J. G.; Johnston-Peck, A. C.; Wang, J.; Tracy, J. B., Size-Dependent Nanoscale Kirkendall Effect During the Oxidation of Nickel Nanoparticles. ACS Nano 2010, 4, 1913-1920.


Page 9 of 10 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(76) Hung, L. I.; Tsung, C. K.; Huang, W.; Yang, P., RoomTemperature Formation of Hollow Cu2O Nanoparticles. Adv. Mater. 2010, 22, 1910-1914. (77) Dong, C.; Yao, D.; Feng, J.; Huang, T.; Hu, X.; Wu, Z.; Liu, Y.; Yang, B.; Zhang, H., Facile Synthesis of Cu2GeS3 and Cu2MGeS4 (M = Zn, Mn, Fe, Co, and Ni) Hollow Nanoparticles, Based on the Nanoscale Kirkendall Effect. Chem. Mater. 2016, 28, 9139-9149. (78) Jin Fan, H.; Knez, M.; Scholz, R.; Nielsch, K.; Pippel, E.; Hesse, D.; Zacharias, M.; Gosele, U., Monocrystalline Spinel Nanotube Fabrication Based on the Kirkendall Effect. Nat. Mater. 2006, 5, 627-631. (79) Niu, K. Y.; Park, J.; Zheng, H.; Alivisatos, A. P., Revealing Bismuth Oxide Hollow Nanoparticle Formation by the Kirkendall Effect. Nano Lett. 2013, 13, 5715-9. (80) Liu, J.; Amit, Y.; Li, Y.; Plonka, A. M.; Ghose, S.; Zhang, L.; Stach, E. A.; Banin, U.; Frenkel, A. I., Reversed Nanoscale Kirkendall Effect in Au–InAs Hybrid Nanoparticles. Chem. Mater. 2016, 28, 8032-8043. (81) Gösele, U.; Frank, W.; Seeger, A., Mechanism and Kinetics of the Diffusion of Gold in Silicon. Appl. Phys. 1980, 23, 361-368. (82) Gösele, U.; Morehead, F.; Frank, W.; Seeger, A., Diffusion of Gold in Silicon: A New Model. Appl. Phys. Lett. 1981, 38, 157159. (83) Gösele, U. M., Fast Diffusion in Semiconductors. Annu. Rev. Mater. Sci. 1988, 18, 257-282. (84) Frank, F. C.; Turnbull, D., Mechanism of Diffusion of Copper in Germanium. Phys. Rev. 1956, 104, 617-618. (85) Tan, G.; Shi, F.; Hao, S.; Chi, H.; Bailey, T. P.; Zhao, L. D.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G., Valence Band Modification and High Thermoelectric Performance in SnTe Heavily Alloyed with MnTe. J. Am. Chem. Soc. 2015, 137, 11507-11516. (86) Tan, G.; Zhao, L. D.; Shi, F.; Doak, J. W.; Lo, S. H.; Sun, H.; Wolverton, C.; Dravid, V. P.; Uher, C.; Kanatzidis, M. G., High Thermoelectric Performance of p-type SnTe via a Synergistic Band Engineering and Nanostructuring Approach. J. Am. Chem. Soc. 2014, 136, 7006-7017. (87) Zheng, L.; Li, W.; Lin, S.; Li, J.; Chen, Z.; Pei, Y., Interstitial Defects Improving Thermoelectric SnTe in Addition to Band Convergence. ACS Energy Lett. 2017, 2, 563-568. (88) Li, W.; Zheng, L.; Ge, B.; Lin, S.; Zhang, X.; Chen, Z.; Chang, Y.; Pei, Y., Promoting SnTe as an Eco-Friendly Solution for p-PbTe Thermoelectric via Band Convergence and Interstitial Defects. Adv. Mater. 2017, 29, 1605887. (89) Dean, J. A.; Lange, N. A., Handbook of Chemistry. McGraw-Hill: 1999.



Page 10 of 10

Insert Table of Contents artwork here


10

PbTe

Recommend Documents