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Structural and Electrical Properties of ([SnSe]1+#)m(NbSe2)1 Compounds: Single NbSe2 layers separated by increasing thickness of SnSe Matti B Alemayehu, Matthias Falmbigl, Kim Ta, Corinna Grosse, Richard Don Westover, Sage R Bauers, Saskia F. Fischer, and David C. Johnson Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/cm5039864 • Publication Date (Web): 06 Jan 2015 Downloaded from http://pubs.acs.org on January 8, 2015

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Structural and Electrical Properties of ([SnSe]1+δ)m(NbSe2)1 Compounds: Single NbSe2 layers separated by increasing thickness of SnSe Matti B. Alemayehu a,*, Matthias Falmbigl a, Kim Ta a, Corinna Grosse b, Richard D. Westover a , Sage R. Bauers a, Saskia F. Fischer b and David C. Johnson a,* a

Department of Chemistry and Materials Science Institute, University of Oregon, Eugene, Oregon 97403, United States b Novel Materials, Humboldt-Universität zu Berlin, 10099 Berlin, Germany

Ferecrystals, charge transfer, transition metal dichalcogenides, thin films

ABSTRACT: The compounds ([SnSe]1+δ)m(NbSe2)1,where 1 ≤ m ≤ 10, were prepared from a series of designed precursors. The c-axis lattice parameter systematically increases by 0.577(5) nm as the value of m is increased, which indicates that an additional bilayer of rock salt structured SnSe is inserted for each unit of m. The in-plane structure of both constituents systematically changes as the thickness of SnSe increases. Both X-ray diffraction and electron microscopy studies show the presence of turbostratic disorder between the different constituent layers. The electrical resistivity and Hall coefficient increase systematically as a function of m stronger than would be expected for noninteracting metallic NbSe2 and semiconducting SnSe layers, suggesting the presence of charge transfer between the layers. The temperature dependence of the resistivity changes from metallic behavior for m < 4 to weakly increasing, for higher m, as temperature decreases. Compounds with m > 3 show an upturn in the resistivity below 50 K and a corresponding increase in the Hall coefficient, which both become more pronounced as m increases. This suggests localization of carriers, which is expected in two-dimensional crystals. The extent of charge transfer in ([SnSe]1+δ)m(NbSe2)1 can be tuned as a function of SnSe thickness and spans over the same range as reported in the literature for various NbX2 based intercalated and misfit layer compounds.

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INTRODUCTION Graphene is the most explored example of a twodimensional atomic crystal whose properties are significant1 ly different from the three-dimensional crystal (graphite). Its unprecedented properties, promising for applications ranging from electronics to composite materials, have generated significant excitement and exploration into other materials with 2D structures that can be cleaved in the search for 2 additional novel phenomena. There is already mention in the literature of creating stacks of different layers with different properties (insulating, conducting, superconducting, magnetic C) to create a new world of materials on demand, combining unusual characteristics together in one material 3 with precise control of the nanoarchitecture. Challenges in achieving this vision include understanding the interaction 4 between layers, determining the structure of each layer 5 and the orientation between layers, and, of course, prepar6 ing these new materials in the first place. The self-assembly of designed precursors has been used to construct intergrowths of different compounds such as two

distinct transition metal diselenides, ([TSe2]1+δ)m(T’Se2)n , 8 9 Bi2Te3 or PbTe with TiTe2, and various rock salt structured monoselenides, ([MSe]1+δ)m(TSe2)n with transition metal diselenides, where M = Sn, Pb, Bi, Ce, T = Ti, V, Nb, Ta, Mo, W, n and m are integers less than 30 and δ describes the difference in the area per cation of the different structures. These latter compounds, referred to as ferecrystals, are similar to misfit layer compounds, where semiconducting monochalcogenide double layers are effectively intercalated between the layers of transition metal dichalcogenide compounds. Unlike ferecrystals, in misfit layer compounds, m and n are typically restricted to 1, although several ex10 amples with n = 2 and 3 have been reported. There is a significant structural difference between ferecrystals and misfit layer compounds. Misfit layer compounds form threedimensional crystals where the constituent structures are distorted in one or both in-plane directions to form a commensurate structure. Ferecrystals have rotational disorder between the different structures and also within the dichalcogenide constituent, and as a consequence, the constituents maintain independent in-plane lattice parameters. The



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independent structure of the subunits in ferecrystals provides opportunities to systematically explore changes in properties by varying n, m or the identity of M and/or T. When n = 1, a single sheet of dichalcogenide is isolated between structurally distinct rock salt layers with an incommensurate interface between them. Recently, there have been many reports of the properties of transition metal dichalcogenides changing as a function of thickness as researchers begin to explore two-dimensional multi- and single layer sheets of these compounds. Twodimensional transition-metal dichalcogenide single layers have promising properties, including moderate carrier mobil2,11 6 12–14 ity, flexibility, direct band gaps, metallic and semi15,16 metallic behavior. Of particular relevance to the research reported herein, a single NbSe2 nanosheet was re16 ported to be semi-metallic and conducts through electrons while its bulk counterpart is metallic and switches carrier 17 type during a charge density wave transition at ~ 35 K. The carrier concentration found in 2D NbSe2 is two orders of magnitude smaller than the carrier concentration per monolayer in 3D NbSe2. This indicates significant change in the electronic structure of NbSe2 from a normal metal in 3D to a 16 semimetal in 2D. Both bulk and 2D NbSe2 have low carrier 16,17 mobility at room temperature. Here we explore the effect of spacing single NbSe2 layers with intervening layers of SnSe ranging from 0.577 to 5.77 nm in thickness. The electrical properties systematically change with the thickness of the SnSe layer, which donates more charge to the NbSe2 layer as the SnSe thickness increases. As the spacing between the NbSe2 layers becomes larger, single NbSe2 layer properties such as localization become more apparent. EXPERIMENTAL SECTION Precursors were designed using a custom-built high vacuum vapor deposition system evacuated to a base pressure -7 18 of 1 × 10 Torr. Sn and Nb were evaporated using Thermionics 3 kW electron beam guns at the rate of 0.04 nm/s and 0.02 nm/s respectively while Se was evaporated at the rate of 0.05 nm/s with a Knudsen effusion cell. The films were deposited onto (100) oriented silicon wafers. Pneumatic shutters located above each source were used to control the amount of material deposited. A sequential LabVIEW program is used to control the deposition scheme. For the formation of the superlattice, the amount of each element required must be calibrated precisely. The calibration procedure of this particular compound is described 19 elsewhere. High angle and low angle X-ray measurements were performed using a Bruker AXS D8 X-ray diffractometer with Cu

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Kα radiation (λ = 0.15418 nm), Göbel mirror and BraggBrentano geometry. Locked coupled scans with (θ-2θ) ranges of 0-10° and 10-65° in 2θ were used for X-ray reflectivity and X-ray diffraction measurements respectively. Total thicknesses and unit cell parameters were determined via X-ray reflectivity and diffraction as reported previously. Inplane diffraction for samples with m < 6 was performed at the advanced photon source at BM-33-C. The ab- plane dimensions for samples with higher m values were measured on a Rigaku Smartlab X-ray diffractometer with Cu Kα radiation (λ = 0.15418 nm). Atomic composition of the precursors as well as the selfassembled compounds were measured with a Cameca SX100 electron probe microanalyzer via a technique described 20 elsewhere. Pieces of the precursors and the annealed samples on silicon were glued onto an aluminum block for analysis. Elemental standards of the constituent elements were used to identify the characteristic radiations coming from the specific elements. A wavelength dispersive spectrometer (WDS) was used to quantify the amount of each element in the compounds. Scanning transmission electron microscopy (STEM) samples were prepared using FEI Helios dual beam using 21 methods developed by Shaffer et al. High angle annular dark field (HAADF) STEM images were acquired on an FEI Titan 80-300. The sample was aligned in the microscope along the nearest silicon substrate zone axis. Selected-area electron diffraction (SAED) was performed at a TEM/STEM JEOL JEM2200FS operated at 200 keV. A SAED aperture of diameter 130 nm was used. The samples were prepared by a conventional cross-section preparation technique with final ion-milling using Ar ions with energies of 5 kV down to 1.4 kV. Temperature-dependent resistivity measurements were conducted via a standard Van der Pauw technique. Indium contacts are made to the four corners of the cross arms. Sheet resistance is determined by sourcing current through two of the cross arms with a programmable current source while controlling the temperature between 20-295 K. The potential across the remaining two cross arms was measured by a nanovoltmeter. Total film thickness from X-ray reflectivity measurements was used to determine the total resistivity of the film from the sheet resistance. Hall measurements were conducted using a constant current value of 0.100 A. The Hall voltage was obtained from the slope of magnetic field vs. voltage plot between 20 and 295 K. Both, the temperature-dependent resistivity and Hall measurements were independently confirmed for m = 1, 4, 5, 6 at the Humboldt-Universität (HU) labs.


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a) b) Figure 1. a) The (009) reflection of ([SnSe]1+δ)3(NbSe2)1 as a function of annealing temperature b) FWHM vs. temperature of the (009) peak of ([SnSe]1+δ)3(NbSe2)1 with a minimum observed at 400 °C for 20 minutes RESULTS AND DISCUSSION A detailed account of the synthesis procedure is described 19 in a previous paper on the 1:1 compound. Briefly, a sequential deposition scheme was used to prepare precursors that mimic the composition profiles of the final products. Sn and Se layers were calibrated to form a Sn:Se bilayer with a 1:1 composition ratio of the elements and a thickness that resulted in the formation of two (001) planes of rock salt structured SnSe per bilayer deposited. Nb and Se layers were calibrated to form a Nb:Se bilayer with a 1:2 composition ratio of the elements and a thickness that resulted in the formation of a single Se-Nb-Se trilayer. A designed number of Sn:Se bilayers (1 to 10 depending on m in the desired product) were evaporated onto a Si (100) wafer followed by one bilayer of Nb:Se. This sequence was repeated until a total thickness of 500 Å was reached. Atomic compositions were determined via electron probe microanalysis (EPMA). In order to see if the targeted compounds could be made and to define the optimum formation conditions for larger m compounds, an annealing study was performed on a precursor designed to form ([SnSe]1+δ)3(NbSe2)1. Six pieces of the same (3,1) sample were annealed at temperatures ranging between 100 °C – 600 °C for 20 minutes. Specular diffraction patterns collected on these annealed pieces contained an increasing number and intensity of 00l reflections as the precursor evolved into crystallographically aligned ferecrystal. The scan taken after the 600 °C annealing re-

vealed that the ferecrystal had begun to decompose. Figure 1 shows the change in the (009) reflection as a function of annealing temperature and a graph of its intensity and full width at half maximum (FWHM) as a function of annealing temperature. The 400 °C scan has reflections with the narrowest line width and highest intensity, indicating optimal formation conditions are achieved at this temperature. The shift in peak position to higher angle with increasing temperature corresponds to the loss of excess selenium from the sample (~ 4%). The samples discussed later in this manuscript were all annealed at 400 °C for 20 minutes. The diffraction patterns of all the ten compounds (Figure 2) show only 00l Bragg reflections arising from the repeat structure due to preferred alignment of the structure on the substrate. For each precursor, a specific compound with a single c-axis lattice parameter was formed. The c-axis lattice parameters increase linearly with increasing number SnSe layers (see Table 1). The slope of the linear fit of the total repeat thickness vs. m (0.577(5)) is slightly smaller than the smallest value of a SnSe thickness in 1:1 misfit layer compounds, which range between 0.5785-0.5990 22 nm. The smaller value results from the difference in the spacing between consecutive rock salt layers as opposed to the dichalcogenide-rock salt distance in the 1:1 compounds. The thickness of the NbSe2 trilayer, 0.648(8) nm, is larger than the previously reported range of thicknesses in differ23 ent polytypes of bulk NbSe2: 0.628-0.638 nm.



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Figure 2. 00l diffraction patterns of all the ten compounds synthesized for the ([SnSe]1+δ)m(NbSe2)1 series. All Bragg peaks can be indexed to their respective repeat unit along the c-axis. Note: index labeling is placed to the left of the corresponding Bragg reflection. Table 1. c-lattice parameter of the ten compounds synthesized with the line width for the peak at 14° in 2θ (Cu Kα radiation). The FWHM slightly varies without any dependency on m. Number of SnSe Layers (m) 1 2 3 4 5 6 7 8 9 10

c-lattice parameter (nm)

FWHM (°) at 14° in 2θ

1.223(1) 1.803(1) 2.378(1) 2.957(1) 3.533(1) 4.106(1) 4.686(1) 5.265(1) 5.841(1) 6.418(3)

0.242(1) 0.216(1) 0.207(1) 0.201(1) 0.242(1) 0.263(1) 0.250(1) 0.273(1) 0.210(1) 0.261(1)

To gain insight to the structure resulting in the diffraction patterns in Figure 2, HAADF-STEM images were collected and that for a (3,1) compound is displayed in Figure 3a. Consecutive units of three double SnSe layers followed by a Se-Nb-Se trilayer of NbSe2 are observed in the examined region. The brighter layers represent the SnSe constituents and the darker ones represent NbSe2. The presence of nanometer scale coherence length localized within the subu24 nits indicates the presence of turbostratic disorder. However, within a unit of three SnSe double layers the same orientation is observed. The trigonal prismatic coordination of the niobium metals is clearly observed from the chevron shapes corresponding to a (100) orientation. On scanning the STEM images collected on several of these compounds with m ≥ 3, we find evidence for regions that have trigonal prismatic as well as octahedrally coordinated niobium atoms (Figure 3b and 3c). In general, the coordination of transition metals in dichalcogenide compounds is governed by the

electronic configuration, atomic radius and preparation con25 26 ditions. NbSe2 is known to have different polytypes. The 1T- polytype has octahedral coordination of the Nb atoms, the 2H- polytype has trigonal prismatic coordination of the Nb atoms, and the 4Hb-NbSe2 polytype has a 50:50 mixture of octahedrally and trigonal prismatically coordinated niobium. The observation of octahedrally coordinated niobium in compounds with larger m values suggests that as m increases there is an increasing charge transfer from the SnSe layers to the NbSe2 layers. A similar structural change was reported for Li intercalated MoS2 where the Mo atom changes coordination from trigonal prismatic to octahedral 27 due to charge transfer from the Li.


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Figure 3. HAADF-STEM image of ([SnSe]1+δ)3(NbSe2)1 compound and an atomic schematic of the structure with green, red and blue spheres representing tin, selenium and niobium respectively. (Oh)-NbSe2 represents octahedrally coordinated Nb while (TP)-NbSe2 represents trigonal prismatic coordination. A schematic that represents the different orientations of the SnSe and NbSe2 within the intergrowth is also displayed. Grazing incidence in-plane diffraction data was collected to determine the in-plane structure of the title compounds. Representative diffraction patterns are displayed in Figure 4a. All peaks can be indexed as hk0 reflections originating from the independent crystal lattices of the SnSe and the NbSe2 constituents (see Figure 4); hence, these reflections contain information about the projection of the structure onto the ab-plane. As m in ([SnSe]1+δ)m(NbSe2)1 increases, there is a systematic increase in the intensities of reflections from the SnSe constituent relative to reflections from the NbSe2 constituent. The peaks corresponding to the hk0 reflections of the NbSe2 constituent can be indexed using a hexagonal in-plane unit cell, similar to 1T-NbSe2 or 2HNbSe2, and the resulting in-plane lattice parameters are listed in table 2. For bulk NbSe2, different ranges for the aaxis lattice parameter were reported depending on the polytype: 1T- (0.353 nm), 2H- (0.3449-0.3460 nm) and 4H28–30 (0.3433-0.3444 nm). The a-axis lattice parameters for NbSe2 in ([SnSe]1+δ)m(NbSe2)1 (0.346-0.347 nm) fall between the 2H- and the 1T- polytypes of NbSe2. The SnSe reflections can be indexed using a square in-plane unit cell, but a slight broadening of hk0 peaks where h≠k (see inset of Figure 4 for the (310) and (130) reflections as an example) suggests a symmetry reduction to a rectangular in31,32 plane unit cell. Table 2 contains the rectangular in-plane unit cell lattice parameters of the SnSe constituent for all of the samples measured.

Figure 4. In-plane diffraction patterns of ([SnSe]1+δ )m(NbSe2)1 for m = 1, 2 and 4. The inset shows the distortion of the SnSe basal plane as m increases from 1 to 4. As summarized in Table 2, the in-plane lattice parameters of both constituents systematically change as m is increased. The difference between the a- and b- in-plane lattice parameters of SnSe increases as a function of m. The distortion of SnSe in the ab-plane in ([SnSe]1+δ)4(NbSe2)1 is 6 times larger than the distortion observed in 33 ([SnSe])4(MoSe2)4. The bigger distortion might be the result of charge transfer from SnSe to NbSe2 as discussed further later in this manuscript. The in-plane area also systematically increases by 2.3% as m is increased (See Figure 5). Based on the logarithmic dependence of the area expansion on m, it is predicted that the area of the SnSe 2 would converge to the bulk area (0.185 nm ) at m = 16. However, the difference between the a- and b-axis lattice parameters in the m = 10 compound is still 3.3% smaller than what is observed in bulk SnSe The a-lattice parameter of the NbSe2 constituent increases slightly as m increases. The resulting change in the in-plane area of the NbSe2 is 0.6% as m increases from 1 to 6, which is a factor of 4 smaller than the observed change in the in-plane area of the SnSe constituent. A similar increase in the a-lattice parameter of VSe2 with increasing SnSe thickness was re33 ported for ([SnSe]1.15)m(VSe2)1 with m = 1, 2, and 3. The net result of the changes in the in-plane lattice parameters of both constituents is a factor of 0.02 for the misfit parameter, δ, range of 0.16 (m = 1) to 0.14 (m = 10) as summarized in Table 2.



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Table 2: In-plane lattice parameters of the SnSe and NbSe2 constituents, and the resulting misfit parameter, δ, for ([SnSe]1+δ)m(NbSe2)1. SnSe shows a distortion into a rectangular basal plane and lattice parameters were extracted from a least squares fit to the indexed hk0-reflections using the space group Pcmn. The lattice parameter of the rectangular plane (ar) can be converted into the setting for the square plane (as) via: √2 .

(1,1)

(2,1)

(4,1)

a-lattice SnSe (nm)

0.4256(4)

0.4268(3)

0.4294(3)

0.4332(2)

0.4335(1)

0.4336(1)

0.4347(2)

0.4354(3)

b-lattice SnSe (nm)

0.4215(4)

0.4221(3)

0.4231(2)

0.4227(2)

0.4228(1)

0.4225(1)

0.4226(1)

0.4217(1)

a-lattice NbSe2 (nm)

0.3462(1)

0.3465(2)

0.3470(1)

0.3472(1)

0.3473(2)

0.3472(2)

0.3471(1)

0.3470(3)

δ

0.157

0.154

0.148

Figure 5. Dependence of the in-plane lattice parameters of SnSe on the number of rock salt layers. The curves are least squares fits to a logarithmic function. Note: error bars are smaller than the size of the symbols. In the selected-area electron diffraction (SAED) patterns in Figure 6 a streaking of the lateral reflections is visible, confirming the turbostratic disorder in the ferecrystals. The streaking is more pronounced for the sample m = 1, because more repeat units, with more layer interfaces, are contained in the selected-area aperture than for m = 6. The SnSe lateral reflections are indexed in Figure 6 using the notation of the crystal structure of binary SnSe at room 34 temperature with space group Pmcn and with lattice parameters found by in-plane X-ray diffraction of the ([SnSe]1+δ)1(NbSe2)1 ferecrystal. The reflections of the 35 NbSe2 layers are labeled according to binary 2H-NbSe2. Due to the limited resolution in the SAED patterns, the significant change of the in-plane lattice parameters with increasing m is not observed. The SAED patterns also show 00l reflections of the ferecrystals, which result from the stacking sequences of the samples. The c-lattice parameters deduced from the SAED patterns agree with those calculated from X-ray diffraction in Table 1. ELECTICAL PROPERTIES The literature on the electrical properties of misfit layer

(6,1)

0.140

(7,1)

(8,1)

0.140

0.140

(9,1)

0.136

(10,1)

0.136

compounds containing NbSe2 discusses the charge transport as being dominated by the conductivity of the NbSe2 constituent. In a first approximation, this suggests that the resistivity of the ([SnSe]1+δ)m(NbSe2)1 compounds should increase linearly with a decreasing volume fraction of NbSe2. The temperature-dependent resistivity data for all of the compounds investigated here are shown in Figure 7. The m = 1 compound was independently prepared 3 times, and the resistivity measured for the different preparations is 19 metal-like and very similar to the one shown in Figure 7. The temperature dependence of the resistivity systematically changes as m increases. For m = 1 to 3, the temperature dependence is very similar to that reported for single crystalline NbSe2. For m > 3, an upturn of the resistivity with decreasing temperature becomes apparent below 40 K suggesting localization of the charge carriers. For m = 6, the resistivity is almost temperature independent until a more pronounced upturn occurs below 50 K. In compounds with m > 6, the electrical resistivity increases with decreasing temperature and the ratio ρ /ρ 295 K rises continuously until it reaches a value of 2 for the m = 10 compound (see inset of Figure 7). The change in the temperature dependence of electrical properties as a function of m suggests a change in conduction mechanism. For m = 1 through 5, the temperature dependence of the electrical resistivity can be modeled using the Bloch-Grüneisen equation as expected for a metal, where ρ0 is the residual resistivity, the electron-phonon interaction constant, and the Debye temperature. .

!

1 11

Fits to this expression are shown in Figure 7 and fitting parameters are summarized in Table 3, which also contains analogous fits to the resistivity of previously reported isoelectronic compounds: ([PbSe]1.14)m(NbSe2)1 with 1 ≤ m ≤ 4. The normalized resistivity increases as a function of m due to a higher volume fraction of the semiconducting SnSe in the unit cell. The values as a function of m increase due to phonon stiffening.


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Figure 6. Selected-area electron diffraction patterns of ([SnSe]1+δ)m(NbSe2)1 for a) m = 1 and b) m = 6. The reflections assigned to the SnSe layers are labeled on the left-hand sides of the images and the reflections of the NbSe2 layers are labeled on the right-hand sides. The 00l reflections in the center of the images result from the stacking sequences of the ferecrystals. Parts of the right-hand side of the images were changed in brightness and contrast to enhance the reflections. This phenomenon was previously observed in intercalated compounds and was attributed to charge transfer from the 36 intercalant to the host material. The presence of increasingly thick SnSe layers between single sheets of NbSe2 also gives the compound increasingly three-dimensional elastic properties and stiffer phonons. Similar trends with a steeper slope and higher absolute values of the are observed in the PbSe containing compounds indicating the presence of higher charge transfer. For m = 6 through 10 where the electrical resistivity rises with decreasing temperature, the conduction cannot be described by the Bloch-Grüneisen equation.

6–10 with an inset showing the resistivity ratio (ρ/ ρ295 K) for m = 6-10. Table 3: Debye Temperature and residual resistivity for ([SnSe]1+δ)m(NbSe2)1 and ([PbSe]1.14)m(NbSe2)1 with m = 05 extracted from Bloch-Grüneisen fits Number of SnSe Layers (m)

θD [K]

ρ0 [µΩm]

Number of PbSe Layers (m)

θD [K]

ρ0 [µΩm]

0* 1 2 3 4 5

119(5) 227(5) 233(5) 213(5) 251(5) 269(5)

0.06 2.18 4.53 5.30 10.03 13.17

0* 1 2 3 4 5

269(5) 316(5) 331(5) 341(5) -

3.41 5.39 11.54 18.62 -

Figure 8 compares the change in room temperature resistivity as a function of m with an extrapolation of the (1,1) behavior that assumes no charge transfer. The systematic increase in room temperature resistivity with increasing thickness of the SnSe constituent (m) occurs at a much higher rate than would be expected for adding additional non-conducting SnSe layers to a metallic NbSe2 layer. Similar behavior was reported previously for ([PbSe]1.14)m(NbSe2)1; however, a higher value of m is required in the SnSe compounds before the resistivity dramatically increases. The higher resistivity than expected from a sum of the constituents suggests that there is a systematic decrease in either the carrier concentration, the mobility, or a combination of both.

Figure 7. Temperature dependent resistivity for m = 1-5 with 17 Bloch-Grüneisen fit with single crystal NbSe2 plotted for comparison and temperature-dependent resistivity for m =



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Figure 8. Change in the room temperature resistivity as a function of m for the ([SnSe]1+δ)m(NbSe2)1 compounds studied in this paper compared with those measured for 37 ([PbSe]1.14)m(NbSe2)1. The straight dashed lines show the expected behavior based on no interactions between a metallic NbSe2 layer and a non-conducting MSe layer. The unanticipated changes in the resistivity as a function of m led us to obtain additional information on the charge transport in the ([SnSe]1+δ)m(NbSe2)1 compounds from Hall effect measurements, which are summarized in Figure 9. For all compounds, the Hall coefficient is positive, indicating that holes are the predominant carriers. Bulk NbSe2 has been reported to have a positive Hall coefficient where SnSe has negative or positive Hall coefficient depending on 17,38 doping levels. There is a systematic increase in the Hall coefficient as m increases, suggesting a corresponding drop in carrier concentration within a single-band model. For m > 1, there is a decrease in the Hall coefficient as temperature is lowered from room temperature and an upturn at temperatures below 50 K. The high temperature slope and the magnitude of the increase in the Hall coefficient at low temperatures both increase in magnitude as m increases. The decline in Hall coefficient as temperature is lowered suggests that the carrier concentration increases with decreasing temperature. The upturn at lower temperatures suggests the increased importance of minority charge carriers and/or localization of carriers. Data from the literature for a single crystal of NbSe2, which experiences a change in the dominant carrier type (from p to n) at 35 K, is also graphed in 17 Figure 9. This change in carrier type is correlated with a charge density wave transition at this temperature. The ([SnSe]1+δ)m(NbSe2)1 compounds all have higher Hall coefficients than single crystal NbSe2, indicating a lower carrier density within the single-band model.

Figure 9. Temperature dependent Hall coefficients measured for ([SnSe]1+δ)m(NbSe2)1 compounds. The Hall coeffi17 cient of a NbSe2 single crystal is plotted as a comparison. For crystalline misfit layer compounds, prior reports have estimated a carrier concentration from the Hall coefficient assuming a single band model. This approximation is rationalized by the observation that the electrical behavior of the misfit compound is dominated by the electrical properties of the dichalcogenide constituent. Figure 10 contains the temperature dependence of the carrier concentration for the compounds discussed in this paper calculated using this same approximation. The magnitude of the room temperature carrier concentration systematically declines as m is increased, as expected from the systematic increase in resistivity. The decay in carrier concentration with increasing m is stronger than expected, based on the volume fraction of the NbSe2 in the unit cell. This suggests that there is charge (electron) transfer from SnSe to NbSe2 and that this charge transfer increases with m. For all samples except m = 1, we see an increase in carrier concentration as temperature is lowered from room temperature. This is not usual for semiconductors or metals and could be a consequence of assuming a single-band model, a possible energy-dependence of scattering rates caused by setting the Hall scattering factor r = 1 in p = r/(RH × e), or a change in the charge transfer between SnSe and NbSe2 as temperature decreases. As m increases, we see a more pronounced decline in carrier concentration at temperatures below 50 K, within the single-band model approximation, suggesting an increased localization of charge carriers as


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the NbSe2 layers are separated further. With increasing m, the effectiveness of the barrier material SnSe increases; hence, the transmission of holes through SnSe is strongly reduced. The transmission through a tunneling barrier depends exponentially on the barrier thickness. For less transmitting barriers, interface scattering is enhanced in the NbSe2 layer and additionally localization of charge carriers becomes feasible. Overall, the change in temperature dependence of the resistivity with m is therefore not simply identified as originating from the change in the temperature dependence of the carrier concentration within the singleband model. However, as shown below, a distinct change in the temperature-dependence of the mobility in the singleband model with m correlates with the resistivity trends and temperature dependence.

tive slope to a positive slope as m is increased. Thus, the change in the temperature dependence of the resistivity appears to be strongly correlated to the change in the temperature dependence of the mobility. However, the room temperature mobility values of all the ten compounds are 2 -1 -1 within the range of ≈ 3-6 cm V s , whereas the absolute change in carrier concentration is significantly larger and accounts for the change in magnitude observed in the resistivity values.

Figure 11. The double logarithmic plot of the temperature dependence of the mobility for the ([SnSe]1+δ)m(NbSe2)1 compounds calculated from resistivity and the Hall coefficient data assuming a single band model for m = 1-10.

Figure 10. The temperature dependence of the carrier concentration calculated from the Hall coefficients assuming a single band model for the ([SnSe]1+δ)m(NbSe2)1 compounds. Within the single-band model, mobility values can be calculated from the measured resistivity and Hall coefficients. The temperature dependencies of such mobility values are shown in Figure 11. The absolute values for the mobility are on average higher than that reported for both single crystals 17 16 and single layers of NbSe2. This suggests that the crystal quality of the compounds reported herein, with respect to carrier scattering, are equivalent to single crystals prepared 39 by near equilibrium vapor phase transport techniques. The most striking part of the mobility values is the systematic change in the temperature dependence, going from a nega-

Similar observations were made in the isovalent 37 ([PbSe]1.14)m(NbSe2)1 compounds. However, comparing the room temperature total carrier concentration of both, ([SnSe]1+δ)m(NbSe2)1 and ([PbSe]1.14)m(NbSe2)1, as a function of rock salt layers reveals that the SnSe containing compounds have higher carrier concentration at each m value (see Figure 11). In both compounds, a similar decline in carrier concentration as a function of m is observed. However, analogous to the resistivity trend in Figure 8, the effect of adding six layers of PbSe is equivalent to adding ten layers of SnSe on the magnitude of the charge carrier density. As summarized in Table 3, the higher Debye temperature and residual resistivity values in ([PbSe]1.14)m(NbSe2)1 are consistent with the presence of more charge transfer in these compounds. On average, there are 20% more charge carriers, at each m, in ([SnSe]1+δ)m(NbSe2)1 as opposed to ([PbSe]1.14)m(NbSe2)1. Similar charge transfer was previously reported for intercalated and misfit layer compounds, with each intercalant and rock salt layer donating more charge as a function of intercalant concentration and valence state of rock salt cation 39 respectively.



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Figure 12. Room temperature total carrier concentration of 37 ([SnSe]1+δ)m(NbSe2)1 and ([PbSe]1.14)m(NbSe2)1 as a function of m with bulk NbSe2 carrier concentration represented by a dashed line. In order to compare the carrier concentration of the ([SnSe]1+δ)m(NbSe2)1 to intercalated bulk NbX2 compounds (X = S and Se) and NbX2 containing misfit layer compounds, we normalized the total carrier concentrations to the volume fraction of NbSe2 in the ([SnSe]1+δ)m(NbSe2)1 unit cell, assuming the carrier concentration in the SnSe constituent is zero. The calculated carriers per Nb atom (h/Nb) at room temperature show a systematic decline as a function of m, from 0.84 h/Nb for m =1 to 0.08 h/Nb for m =10, as shown in Figure 13. The decline in carriers is attributed to electron transfer from SnSe to the originally halffilled d band of NbSe2. The extent of charge transfer reported for various NbX2 based misfit layer and intercalated compounds spans over a similar range (0.10 – 0.87 h/Nb) (Figure 13) with the amount of charge transfer controlled by the valence of elements incorporated into the rock salt lay39 er. In the case of intercalation compounds, h/Nb can be varied by changing the amount of the intercalated atom, restricted only by the solubility limit of the intercalated atoms. In the compounds synthesized here, we are able to tune the electronic structure of the NbSe2 within the same h/Nb range as obtained for various misfit layer and intercalated compounds, by varying the thickness of the SnSe layer. In order to compare the extent of charge transfer in ([SnSe]1+δ)m(NbSe2)1 and intercalated and misfit layer compounds, a least squares fit to an exponential function was applied to h/Nb vs. m of ([SnSe]1+δ)m(NbSe2)1. Based on the exponential fit, the equivalent m values that result in the exact h/Nb values of intercalated and misfit layer compounds were calculated (see Figure 13). The same extent of charge transfer as observed for trivalent rock salt cations (Tb, La) in misfit layer compounds can be achieved by increasing the thickness of the SnSe layer in ([SnSe]1+δ)m(NbSe2)1 from seven to ten. Our ability to finetune the electrical properties of these compounds by systematically varying the constituent thickness in the intergrowth provides a unique opportunity to design materials with targeted properties.

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Figure 13. h/Nb for ([SnSe]1+δ)m(NbSe2)1 ferecrystal with comparative NbX2 based misfit layer compounds (MLCs) and intercalated compounds fitted to an exponential function (dashed line) that gives the equivalent m values accord39 ing to their h/Nb value as reported by Wiegers. CONCLUSION The compounds ([SnSe]1+δ)m(NbSe2)1, where 1 ≤ m ≤ 10, were synthesized by systematically changing the SnSe layer thickness. A corresponding increase in the c-lattice parameter of 0.577(5) nm was observed as a function of SnSe layer. The in-plane unit cell parameters of the SnSe constituent increase as well as the degree of mismatch between the SnSe and NbSe2 as a function of m. HAADF-STEM revealed mixed coordination of niobium in NbSe2: trigonal prismatic and octahedral. The temperature dependent resistivities for m < 5 compounds exhibit metallic behavior while a semiconductor-like temperature dependence is observed for m > 5 compounds due to localization of carriers at low temperatures. The electrical properties systematically vary as a function of the thickness of the SnSe layer but at a much higher rate than would be expected for noninteracting layers. A corresponding change in carrier concentration and mobility clearly reveals the presence of charge transfer between SnSe and NbSe2. However, the extent of charge transfer is less than in previously reported isovalent ([PbSe]1.14)m(NbSe2)1 compounds. The ability to systematically change the thickness of the SnSe constituent allows a precise tuning of the electronic properties of the NbSe2 equivalent to that of intercalated compounds with different intercalant concentration and misfit layer compounds with rock salt metals of different valence states. Within this investigation, we prove that the precise control of the layer thickness of individual layers offers a unique opportunity to achieve the goals of rational materials design.

AUTHOR INFORMATION Corresponding Authors * [email protected] and [email protected]

ACKNOWLEDGMENT The authors thank Josh Razink and Robert Fischer in CAMCOR for assistance in preparing TEM samples and collecting STEM images. Grant MRI 0923577 provided


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funding for the dual beam FIB used to make TEM cross sections. The authors acknowledge support from the National Science Foundation under grant DMR-1266217. Coauthor MF acknowledges support from the National Science Foundation through CCI grant number CHE-1102637. The authors thank Jenia Karapetrova for assistance at Beamline 33-C at the Advanced Photon Source (APS) in Argonne National Laboratories. The use of the APS was supported by the U.S. Department of Energy, Office of Science, and the Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357.

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Table of Content


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Figure 1. a) The (009) reflection of ([SnSe]1+δ)3(NbSe2)1 as a function of annealing temperature b) FWHM vs. temperature of the (009) peak of ([SnSe]1+δ)3(NbSe2)1 with a minimum observed at 400 °C for 20 minutes 308x108mm (300 x 300 DPI)



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Figure 2. 00l diffraction patterns of all the ten compounds synthesized for the ([SnSe]1+δ)m(NbSe2)1 series. All Bragg peaks can be indexed to their respective repeat unit along the c-axis. Note: index labeling is placed to the left of the corresponding Bragg reflection. 254x117mm (300 x 300 DPI)


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Figure 3. HAADF-STEM image of ([SnSe]1+δ)3(NbSe2)1 compound and an atomic schematic of the structure with green, red and blue spheres representing tin, selenium and niobium respectively. (Oh)NbSe2 represents octahe-drally coordinated Nb while (TP)-NbSe2 represents trigonal prismatic coordination. A schematic that represents the different orientations of the SnSe and NbSe2 within the intergrowth is also displayed. 173x127mm (150 x 150 DPI)



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Figure 4. In-plane diffraction patterns of ([SnSe]1+δ )m(NbSe2)1 for m = 1, 2 and 4. The inset shows the distor-tion of the SnSe basal plane as m increases from 1 to 4. 84x77mm (150 x 150 DPI)


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Figure 5. Dependence of the in-plane lattice parameters of SnSe on the number of rock salt layers. The curves are least squares fits to a logarithmic function. Note: error bars are smaller than the size of the symbols. 203x127mm (125 x 132 DPI)



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Figure 6. Selected-area electron diffraction patterns of ([SnSe]1+δ)m(NbSe2)1 for a) m = 1 and b) m = 6. The reflections as-signed to the SnSe layers are labeled on the left-hand sides of the images and the reflections of the NbSe2 layers are la-beled on the right-hand sides. The 00l reflections in the center of the images result from the stacking sequences of the ferecrystals. Parts of the right-hand side of the images were changed in brightness and contrast to enhance the reflections. 163x62mm (78 x 77 DPI)


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Figure 7. Temperature dependent resistivity for m = 1-5 with Bloch-Grüneisen fit with single crystal NbSe217 plotted for comparison and temperature-dependent resistivity for m = 6–10 with an inset showing the resistivity ratio (ρ/ ρ295 K) for m = 6-10. 111x123mm (150 x 150 DPI)



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Figure 8. Change in the room temperature resistivity as a function of m for the ([SnSe]1+δ)m(NbSe2)1 compounds studied in this paper compared with those measured for ([PbSe]1.14)m(NbSe2)1.37 The straight dashed lines show the expected behavior based on no interactions between a metallic NbSe2 layer and a non-conducting MSe layer. 147x114mm (86 x 86 DPI)


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Figure 9. Temperature dependent Hall coefficients meas-ured for ([SnSe]1+δ)m(NbSe2)1 compounds. The Hall coeffi-cient of a NbSe2 single crystal 17 is plotted as a compari-son. 95x127mm (150 x 150 DPI)



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Figure 10. The temperature dependence of the carrier concentration calculated from the Hall coefficients assuming a single band model for the ([SnSe]1+δ)m(NbSe2)1 com-pounds.

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Figure 11. The double logarithmic plot of the temperature dependence of the mobility for the ([SnSe]1+δ)m(NbSe2)1 compounds calculated from resistivity and the Hall coefficient data assuming a single band model for m = 1-10. 175x127mm (72 x 73 DPI)



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Figure 12. Room temperature total carrier concentration of ([SnSe]1+δ)m(NbSe2)1 and ([PbSe]1.14)m(NbSe2)137 as a function of m with bulk NbSe2 carrier concentration represented by a dashed line. 174x127mm (73 x 73 DPI)


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Figure 13. h/Nb for ([SnSe]1+δ)m(NbSe2)1 ferecrystal with comparative NbX2 based misfit layer and intercalated compounds fitted to an exponential function (dashed line) that gives the equivalent m values according to their h/Nb value as reported by Wiegers.39 164x127mm (77 x 77 DPI)


Structural and Electrical Properties of ([SnSe]1+Î´) - ACS Publications

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