Synthesis of 1T, 2H, and 6R Germanane Polytypes - Chemistry of

Article Cite This: Chem. Mater. 2018, 30, 1335−1343

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Synthesis of 1T, 2H, and 6R Germanane Polytypes Nicholas D. Cultrara,† Yaxian Wang,‡ Maxx Q. Arguilla,† Michael R. Scudder,† Shishi Jiang,† Wolfgang Windl,‡ Svilen Bobev,§ and Joshua E. Goldberger*,† †

Department of Chemistry and Biochemistry, The Ohio State University, Columbus, Ohio 43210, United States Department of Materials Science and Engineering, The Ohio State University, Columbus, Ohio 43210, United States § Department of Chemistry and Biochemistry, University of Delaware, Newark, Delaware 19716, United States ‡

S Supporting Information *

ABSTRACT: Polytypism, or the ability for materials to crystallize with different stacking sequences, often leads to fundamentally different properties in families of two-dimensional materials. Here, we show that is possible to control the polytype of GeH, a representative two-dimensional material that is synthesized topotactically by first controlling the polytype sequence of the precursor Zintl phase. 1T, 2H, and 6R GeH can be prepared by the topotactic deintercalation of 1T EuGe2, 2H α-CaGe2, and 6R β-CaGe2, respectively. The 6R and 1T GeH polytypes exhibit remarkably similar properties and feature band gaps of 1.63 and 1.59 eV, respectively. However, the 2H CaGe2 precursor forms due to the incorporation of small amounts of In flux in the germanium lattice, which is retained when converted to GeH. Consequently, 2H GeH has a reduced band gap of 1.45 eV. Finally, temperature dependent diffraction of 6R GeH shows a negative coefficient of thermal expansion along the a-axis and a positive coefficient of thermal expansion along the out-of-plane c-axis.

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The 2H and 3R phases have Mo4+/W4+ in trigonal prismatic coordination and are semiconductors. In MoS2, for example, both the 2H and 3R phases have observed optical band gaps near ∼1.29 eV,14,15 and computational predictions indicate the 2H and 3R MoS2 polytypes have band gaps of 1.29 and 1.33 eV, respectively.16 The 2H and 3R phases also feature slightly different bond lengths, different Raman modes, exciton binding energies, and temperature dependent lattice constant changes.8,14,16,17 Two-dimensional germanium based materials have also been discovered recently. One such system is m-allo germanium, a layered Ge polymorph consisting of covalently bonded layers of [Ge12] derived by the topotactic deintercalation and oxidation of Li7Ge12.18−20 Another germanium system, hydrogenterminated germanane (GeH), is a 2D material which has attracted considerable interest.21−25 Germanane is a 2D germanium graphane analogue in which the germanium atoms arrange in a puckered honeycomb layer and are terminated with covalently bonded −H ligands, alternating above and below each Ge atom in the network. It has a highpredicted electron mobility and has been recently shown to be active element in field effect transistor devices, hydrogen evolution photocatalyst, and as a Li battery electrode.26,27 In

INTRODUCTION Recently, there has been a large research focus on the synthesis and properties of two-dimensional (2D) materials. It has been found that many materials whose crystal structures consist of 2D networks of atoms that are separated by van der Waals forces can have different electronic and thermal properties when exfoliated down to layers with precise layer numbers.1−6 It has been well-established that the band gap and transport properties of these materials can be dramatically influenced by their immediate surroundings.1−3,7−9 This has led to more recent efforts focused on stacking and understanding how to couple neighboring layers together to create exotic physical phenomena.10−13 This surface sensitivity arises partly because the orbitals comprising the conduction and valence band are often oriented toward and interact with their surroundings as well as due to differences in the local dielectric constant. Many layered solid-state materials can form different polytypes, in which each layer is virtually identical, but there are different stacking sequences in a single unit cell. These stacking sequences often give rise to differences in the band gap and electronic structure. This represents another manifestation of how the immediate surroundings influence the properties of 2D materials. For example, MoS2, MoSe2, and WS2 can crystallize into the 1T (one layer per trigonal unit cell), 2H (two layers per hexagonal unit cell), and 3R (three layers per rhombohedral unit cell) polytypes.14 The 1T phases are all metallic with Mo4+ and W4+ in octahedral coordination to S. © 2018 American Chemical Society

Received: November 29, 2017 Revised: January 29, 2018 Published: February 2, 2018 1335

DOI: 10.1021/acs.chemmater.7b04990 Chem. Mater. 2018, 30, 1335−1343

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Chemistry of Materials

Figure 1. Schematic representations of the structures of each precursor Zintl phase and the structures of the corresponding deintercalated phases. Intralayer covalent bonding between the germanium atoms (purple) is known for each phase and emphasized with cylinders drawn between closest neighbors, while the interactions between the calcium (yellow) or europium (magenta) atoms and the germanium atoms are mostly electrostatic, and no bonds are drawn. In the corresponding deintercalated phases, the hydrogen atoms terminating the layers are shown in pink.

general, GeH has a direct band gap of ∼1.6 eV, which can be tuned from 1.4 to 1.7 eV by substituting the H-terminating ligand with an organic moiety which, due to their sterics and electronics, strain the germanane framework.28 This ability to produce large variations in electronic structure makes it an intriguing 2D material for studying the influence of polytypism on its properties. However, GeH is a kinetically trapped phase that can be prepared only through the topotactic deintercalation of a precursor intermetallic Zintl phase containing germanium atoms in a structurally analogous framework. Zintl phases refer to compounds formed between the electropositive group 1 or 2 elements and the more electronegative group 13−15 elements,29 whose structure and bonding can be rationalized using the Zintl−Klemm concept.30 Therefore, the preparation of different germanane polytypes requires the ability to control the stacking arrangement of germanium atoms in the Zintl phase precursor. To these ends, the unit cells for EuGe2, α-CaGe2, and β-CaGe2 correspond to the 1T, 2H, and 6R (six layers per rhombohedral unit cell) stacking arrangements of germanium layers, respectively, which are separated by the divalent Eu2+ and Ca2+ cations.31−33 Previous studies have solely focused on studying GeH transformed from the 6R β-CaGe2 phase.21−24

Herein, we successfully prepared 1T, 2H, and 6R polytypes of GeH through the reaction of EuGe2, α-CaGe2, and β-CaGe2, respectively, with HCl. This shows there is retention of the stacking sequence through the topotactic deintercalation process. We elucidated that the 2H α-CaGe2 phase, synthesized using indium flux, has about 2−3% indium substitution on the germanium layers. After HCl treatment, these indium substitutions remain on the germanane framework and become terminated with −OH, which is reflected in the Raman and infrared spectra, the interlayer spacings, and a reduction in the optical band gap compared to the 1T and 6R phases. In contrast, the vibrational and electronic properties of the 1T and 6R GeH phases are very similar. Finally, we characterized the temperature dependent changes in lattice constants for 6R GeH, the most prevalent polymorph, and show that there is a negative coefficient of thermal expansion in the in-plane direction and a positive coefficient in the cross-plane direction.

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EXPERIMENTAL METHODS

Synthesis. All reactions were carried out in evacuated fused silica tubes which were loaded in an argon filled glovebox using methods adopted previously.21,31,32 In the growth of 6R CaGe2, stoichiometric amounts of calcium and germanium are loaded into fused silica tubes and sealed while under evacuation to pressures 50 h. The tubes were then inverted and centrifuged, separating excess In from the crystalline CaGe2. 1T EuGe2 was obtained by loading stoichiometric amounts of europium and germanium in an alumina crucible and sealed in a fused silica tube. The sample was heated at 1050 °C for 24 h followed by cooling to room temperature for 24 h. After synthesis, all crystals were collected and placed in concentrated hydrochloric acid at −40 °C from 5 to 40 days until complete reaction of the precursor phase. Appreciable efforts increase the crystallinity of the 2H and 1T phases by varying acid concentration and temperature of deintercalation. Following the deintercalation process, samples were washed using H2O then methanol before collection using centrifugation. Following centrifugation, samples were dried in vacuum using a Schlenk line. Characterization. Powder X-ray diffraction (XRD) patterns were collected for all Zintl phase precursors and deintercalated germanane phases using a combination of in-house and synchrotron techniques. The Zintl phases were measured in flat plate mode using an in-house Bruker AXS D8 diffractometer employing Cu Kα1 radiation with λ = 1.5406 Å. Powder diffraction pattern for the 1T GeH sample was taken after sealing the sample in a capillary and measured in-house while the 6R and 2H GeH samples were sealed in capillaries, and powder diffraction patterns were collected at beamline 11-BM at Argonne National Laboratory using wavelengths of 0.459255 and 0.4141660 Å, respectively. Rietveld refinements for 6R and 2H GeH and Le Bail refinement of the 1T GeH phases were performed using GSAS, while all Zintl phases were characterized via Rietveld refinement on TOPAS. Raman spectra were collected using a Renishaw InVia Raman microscope exciting with a 785 nm diode laser at 162 kW/cm2 and

50 s exposure, equipped with a CCD detector at room temperature. For temperature dependent Raman spectra, a 633 (He−Ne) laser source was used. The relative elemental composition was measured using X-ray fluorescence using an Olympus X-5000 Mobile XRF System. Fourier transform infrared (FTIR) spectra were collected with a PerkinElmer Frontier Dual-Range FIR/MidIR spectrometer that was loaded in an N2-filled glovebox and collected in transmission mode after forming a mixed KBr-GeH pellet. Diffuse reflectance absorption measurements were collected using a PerkinElmer Lambda 950 UV/ vis NIR spectrophotometer with a combination of a silicon photodiode and an InGaAs detector. Electronic Structure Calculations. Band gaps of GeH with different polytypes were confirmed by density functional theory (DFT) calculations using the Vienna Ab Initio Simulation Package (VASP).34,35 For structures where relaxations were necessary, we used PBE36 projector augmented-wave (PAW) potentials37 and Grimme’s DFT-D2 method38 to describe the van der Waals interaction between layers. A cutoff energy of 600 eV was necessary for satisfactory convergence of the structural optimization. The same settings were used for band structure calculations, except that the Heyd−Scuseria− Ernzerhof (HSE06) hybrid functional39,40 was employed for accurate description of the band structures.

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RESULTS AND DISCUSSION First, polycrystalline 6R and 2H CaGe2 and 1T EuGe2 (Figure 1) were synthesized for their subsequent deintercalation into germanane. Figure 2a shows the powder X-ray diffraction pattern for each of these precursor Zintl phases. These samples were highly crystalline, and the observed impurity phases include Ge in β-CaGe2, residual In flux in α-CaGe2, and Ge and trace Eu3Ge5 in EuGe2. Rietveld analysis of these patterns 1337


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Chemistry of Materials Table 1. Structural Parameters from Rietveld (for 6R and 2H) and Le Bail (1T) Refinement of Deintercalated GeH space group a (Å) c (Å) Ge(1) Ge(2) U11U22 (Å2) U33 (Å2) wRp/Rp

6R GeH

2H GeH

R3m ̅ 3.97142(5) 33.033(5) x = 0, y = 0, z = 0.1772(1) x = 0, y = 0, z = 0.3418(1) 0.03372(2) 0.583(6) 0.1217/0.0978

P63mc 3.9535(2) 11.64(9) x = 0, y = 0, z = 0.00(1) x = 1/3, y = 2/3, z = 0.561(1) 0.0413(4) 0.790(6) 0.0737/0.0584

(Figures S1−3) indicated a phase purity of 91% for β-CaGe2, 89% for α-CaGe2, and 49% for EuGe2 (Tables S1−6). The Zintl phases were then reacted at −40 °C in concentrated HCl for 1−4 weeks to transform them into the GeH phases. Figures 2b−d show the XRD pattern of each of the GeH phases produced from the 6R β-CaGe2, 2H α-CaGe2, and 1T EuGe2 phase. The topotactic deintercalation of all three phases produce three unique unit cells that are structurally related to the Zintl phase precursor. Furthermore, residual Ge and Eu3Ge5 is observed in the 6R and 1T phases, while the residual indium flux in the 2H precursor is completely dissolved by the concentrated HCl. The resulting GeH phases can be indexed to 6R (R3̅m), 2H (P63mc), and 1T (P3̅m1) unit cells, respectively, with the unit cell parameters given in Table 1. Each of these phases has approximately the same primitive unit cell a-axis lattice constant, but can feature one-, two-, or sixlayer unit cells along the c-axis (Figure 1). Consequently, each phase can be identified using distinct reflections in the diffraction pattern. For instance, the 1T phase has the 011 reflection for which the analogous 6R reflection 016 is forbidden. Additionally, the 011 reflection in 2H GeH would not have an equivalent reflection in 1T GeH. Similarly, the 6R structure features the 012 and 014̅ reflections, which do not exist in the other unit cells. The GeH reflections in all phases also show significant broadening compared to the precursor Zintl phases, which is indicative of smaller crystalline domain sizes. Due to the broadness and presence of overlapping reflections, the structure of the originally deintercalated germanane phase was misidentified as the 2H GeH structure. The use of high-resolution synchrotron powder diffraction data is essential to now unambiguously distinguish between the 6R and 2H GeH phases.33 The structures of the 2H and 6R phases were confirmed using Rietveld analysis. Due to the small numbers and broadness of reflections in the 1T phase, Le Bail refinement was used to confirm the structure (Table 1, Figures S4−6, and Tables S7−9). Excellent refinements can be achieved only when the GeH space group is the same as the parent Zintl phase. Furthermore, anisotropic thermal parameters greatly improve the refinements of the 2H and 6R phases; however, they result in U33 values that are unrealistically large (0.6−0.8 Å2). This is commonly observed in all germanane refinements and can be attributed to the distribution of interlayer distances to these topotactically deintercalated phases. The GeH phases all have c-axes that are expanded and a-axes that are contracted from to the original Zintl phases. Specifically, the a-axis of β-CaGe2 shrinks from 3.9872 to 3.97142 Å in 6R GeH, and the c-axis increases from 30.583 to 33.033 Å, which corresponds to a thickness of about 5.51 Å per layer. This increase in the c-axis is expected due to the replacement of the Ca2+ ion with 2 Ge−H bonds between each layer. Next, the a-axis of α-CaGe2 shrinks from 3.9966 to

1T GeH (Le Bail) P3m ̅ 1 3.9499(5) 5.776(11)

0.0215/0.0178

3.9543 Å in 2H GeH, and the c-axis increases from 10.211 to 11.64 Å, which corresponds to a thickness of about 5.82 Å per layer. Interestingly, the 2H GeH has a c-axis much larger than that of 6R GeH, which will be subsequently explained by the presence of residual indium in the framework, resulting in In− OH bonds. Finally, the a-axis of EuGe2 shrinks from 4.1035 to 3.9499 Å in 2H GeH, and the c-axis increases from 4.9972 to 5.776 Å, which corresponds to ∼5.78 Å per layer. This c-axis is smaller than what is observed in the 2H phase. Finally, it is important to note that the broadness of the 2H GeH reflections makes it possible that there could be some local 6R ordering in this phase; however, very reasonable Rietveld fits could be obtained assuming 2H GeH only. It is interesting that the polytype is retained after topotactic transformation. This can be understood by examining the differences in the 6R, 2H, and 1T polymorph crystal structures (Figure 1). In the 1T phase, each puckered honeycomb germanium layer is stacked perfectly on top of each other from one unit cell to the next. In the 2H phase, there are two different puckered honeycomb layers. These layers are rotated by 30° from each another. The 6R polymorph consists of six different layers having a stacking sequence that we denote as AA′BB′CC′. There is a 30° rotation between each layer, which is emphasized using the prime notation. Furthermore, there is a 1/3 a and 2/3 b translation between every other layer, for example, between A to B and B to C. The fact that the 1T phase contains no rotation between layers, whereas the 2H and 6R do, suggests that there is a large energy cost for transforming between these different polytypes. Infrared spectroscopy was used to further confirm hydrogen termination for the three GeH polytypes. All three frameworks exhibit an intense Ge−H stretching mode at ∼2000 cm−1 as well as multiple Ge−H wagging modes at 475, 507, and 570 cm−1. In 2H GeH, the Ge−H stretching mode is centered at 1980 cm−1, which is red-shifted by ∼20 cm−1 in comparison to the 6R and 1T phases (Figure 3). Additionally, there are weak vibrational modes at 770 and 826 cm−1 that have been previously attributed to Ge−H2 bending modes that can appear due to Ge vacancies or on the edges. The 2H GeH phase additionally features a mode at 3650 cm−1, indicative of an O−H stretching peak, as well as a vibrational mode centered at 650 cm−1. The new vibrational modes and the red-shifting of the 2H Ge−H stretch can be readily explained by the incorporation of a small amount of indium substituted onto the germanium framework in 2H CaGe2, which is retained upon topotactic deintercalation and is terminated with −OH. The new vibrational modes centered at 650 cm−1 arise from In−OH termination, as we previously demonstrated that hydroxide terminating the Ge−OH framework appears at around 850 cm−1, and Sn−OH is centered at 560 cm−1.41,42 1338


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Figure 3. FTIR spectra of 6R (black), 2H (red), and 1T (blue) germanane products.

XRF measurements confirm that after deintercalation, the 2H GeH phase has 0.027 mol of indium per mole of germanium in the framework (Figure 4b). It is important to

Figure 5. (a) Raman and (b) DRA spectra of 6R (black), 2H (red), and 1T (blue) germanane products. Figure 4. (a) Shift in Ge−H stretch using FTIR in 2H (red) compared to 1T (blue) and 6R (black) germanane, associated with heavier In atom substituted on the framework. (b) Standardized analysis of 2H germanane XRF spectrum (inset) confirms 2.6% indium is retained by germanane. Standards are depicted as black circles

densities of states. Of the three polytypes, 6R GeH has the largest band gap at 1.63 eV, followed by the 1T at ∼1.59 eV. The 2H GeH has a much lower band gap, occurring at ∼1.45 eV. The reduction of the band gap for 2H GeH is likely due to a combination of the heavier In atom on the germanane framework itself along with the electron-withdrawing −OH termination.28,41 The small change in band gaps of GeH depending on polytype was further confirmed using DFT simulations. As described above, Le Bail and Rietveld analyses found layer spacings of 5.78, 5.84, and 5.51 Å for 1T, 2H, and 6R, respectively. While calculations for 6R with the experimental structural data resulted in band gaps in the same range as measured, the large interlayer spacing in 1T and 2H resulted in band gaps that were too large by more than 0.5 eV. Because of the weak interlayer interactions in the van der Waals solid, the simulated lattice constants are extremely sensitive to the specific DFT method utilized. The values can change significantly depending on the dispersive interaction model. Furthermore, as we have shown previously, there is a large sensitivity of the band gap to the Ge−Ge and Ge−H bond lengths as well as the Ge−Ge−Ge and Ge−Ge−H bond angles.28,43 Holding the experimental lattice constant but altering the Ge framework to have different degrees of buckling and Ge−Ge bond lengths can change the band gap by ±0.5 eV, and one could design GeH frameworks that have the same experimental lattice constants but different degrees of buckling that lead to the observed band gaps. Therefore, to be scientifically rigorous, we chose to simulate the band gap for the fully relaxed cells for the 1T and 2H phases. To investigate this discrepancy, we repeated band structure calculations for those structures for fully relaxed cells, finding a = 4.05 Å and an interlayer spacing of 5.41 Å for both 1T and 2H. These lattice constants were used to calculate the final

point out that no In peaks were observed in the XRD (Figure 2c), and that after washing the sample with HCl multiple times, the molar percentage of indium in the XRF spectrum did not change. Consequently, we hypothesize that in the indium flux synthesis conditions of α-CaGe2, indium substitutes with germanium onto the germanane framework and is retained through the deintercalation process but terminated with hydroxide at these indium sites. Indeed, the broadness and red-shifting of the 2H Ge−H infrared stretching frequency relative to the 6R and 1T phases (Figure 4a) is indicative of a heavier atom on the germanane framework. Such changes have been previously observed when Sn is substituted onto the germanium lattice.42 The Raman spectroscopy of isolated GeH flakes show subtle differences between the 1T and 6R GeH phase and a much larger change in the 2H phase. As expected, the incorporation of the heavier In atom onto the germanane framework causes the Raman modes to shift to lower wavenumbers. In the Raman spectra (Figure 5a), the intense in-plane Ge−Ge E2 modes for 1T and 6R GeH occur at 301.6 and 301.8 cm−1, respectively, while the 2H phase shifts to 300.2 cm−1. Furthermore, the outof-plane A1 mode for the 1T and 6R GeH phases both occur at 227.7 cm−1, whereas it occurs at 225.4 cm−1 for the 2H phase. Diffuse reflectance absorbance measurements were collected to elucidate how the band gap changes with polytype (Figure 5b). It has been previously established that the linear fittings of the Kubelka−Munk functions provide an excellent approximation of the relative band gaps for germanane materials, as Tauc− Davis Mott models are often ambiguous due to the reduced 1339


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and the shift in the Ge−H stretching vibration in the FTIR spectra. Furthermore, it also explains why the 2H α-CaGe2 forms while using In flux. The 2H CaGe2 polymorph was also observed when small percentages of Sn (∼3%) were substituted onto the CaGe2 the framework.41 Thus, the substitution of small amounts of Ge in CaGe2 with a larger 5p element promotes crystallization into the 2H polytype and not the 6R polytype. Our previous studies of GeH materials derived from the deintercalation 6R CaGe2 originally assigned these phases as 2H GeH; however, the broadness and the few number of reflections in lab X-ray data led to an incorrect assignment.21,31,32 In this work, the use of higher resolution synchrotron XRD enables the observation of many more diffraction reflections, which allows us to more definitively assign these phases as 6R GeH. This is in agreement with simulations from Lou et al.33 Thermal Parameters of the 6R Phase. For any material, determining the changes in structure as a function of temperature is essential for understanding their thermal transport phenomena as well as evaluating their possible application in mechanical and thermoelectric devices. Thus far, almost all studies of thermal expansion on 2D materials have been theoretical in nature. Here, we evaluated the changes in lattice constants for 6R GeH, the most prevalent polymorph, as a function of temperature via X-ray diffraction with synchrotron radiation. Capillary mode powder diffraction patterns of 6R GeH, with an internal Ge standard, were collected at 100, 120, 140, 160, 180, 220, 260, and 295 K, as shown in Figure 7a. The lattice parameters of Ge and 6R GeH were determined via refining the XRD patterns using a Le Bail method. The changes in lattice constants for the internal Ge standard are in excellent agreement with previously reported measurements44 (Figure S7). 6R GeH exhibits an expansion of the interplane spacing as the temperature increases from 100 to 295 K because the c lattice parameter increases from 32.939(4) to 33.016(4) Å, an increase of ∼0.3% (Figure 7b). We estimated by drawing a linear trend line through the data that the 6R structure expands in the c-axis with a thermal expansion coefficient of 1.1 × 10−5 K−1. Conversely, as the temperature is increased, the in-plane lattice parameter contracts from 3.9761(5) to 3.97203(5) Å with a thermal expansion coefficient of −5.0 × 10−6 K−1. The volume of the GeH phase expands modestly, by 0.1%, as temperature increases. Other layered van der Waals materials such as arsenic and graphite for bulk materials, also exhibit a negative thermal expansion along the in-plane direction and a positive thermal expansion coefficient along the out-of-plane direction at low temperatures.45 Additionally, the Raman spectra were collected at each temperature for which diffraction was obtained (Figure 7c). The out-of-plane A1 vibration located at 227−229 cm−1 and the in-plane E2 vibration located at 301−304 cm−1 both decrease in wavenumber as the temperature of the system is increased (Figure 7d). This relationship is directly correlated to the increase in Ge−Ge and Ge−H bond lengths as the temperature increases. Similar trends are widely observed in other 2D materials such as MoS2 and phosphorene.46−48

band structures. The simulated band gaps for 1T, 2H, and 6R are 1.56, 1.61, and 1.59 eV (Figure 6). Taking into account that

Figure 6. DFT calculations of (a) 1T, (b) 2H, and (c) 6R GeH. The band gaps at Γ and A are shown.

2H GeH is likely to contain In atoms, which results in a smaller band gap of ∼1.45 eV, good consistency is found between DFT calculations and diffuse reflectance absorbance measurements. Both 2H and 6R have a direct band gap at the Γ point of the Brillouin zone, while 1T has a direct band gap at the A point. For 6R, the energy gaps at Γ and A are only slightly different. The presence of In on the 2H GeH framework terminated with −OH explains why it has the lowest observed band gap, largest interlayer distance, lowest wavenumber Raman mode,

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CONCLUSION Here, we demonstrated that each of the three precursor Zintl phase materials can be synthesized and subsequently deintercalated to prepare the 1T, 2H, and 6R polytypes of germanane. X-ray diffraction showed that the layer stacking 1340


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Figure 7. (a) Synchrotron XRD and (c) Raman spectroscopy data at different temperatures. (b) Temperature dependent a- (triangles) and c(circles) lattice parameters. (d) Raman shifts of the A1 (triangle) and E2 (circles) modes in the 6R germanane phase. Measurements were carried out at 100 (black), 120 (blue), 140 (purple), 160 (dark green), 180 (light green), 220 (yellow), 260 (orange), and 295 (red) K.

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ACKNOWLEDGMENTS We acknowledge the Analytical Spectroscopy Laboratory and the Surface Analysis Laboratory (NSF DMR-0114098) of The Ohio State University Department of Chemistry and Biochemistry and The Ohio State University Nanosystems Laboratory (NSL). Use of the Advanced Photon Source at Argonne National Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract DE-AC02-06CH11357. Primary funding for this research was provided by NSF EFRI1433467. Partial Funding for this research was provided by the Center for Emergent Materials: an NSF MRSEC under Award DMR-1420451. Work at the University of Delaware received financial support from the National Science Foundation, Grants DMR-0743916 (CAREER) and DMR1709813. J.E.G. acknowledges the Camille and Henry Dreyfus Foundation for partial support.

sequence is retained through the deintercalation process. Each structure shows the hallmark Ge−H stretches in FTIR and similar Raman vibrations, indicating the germanane is in similar chemical environments despite the changes in the stacking sequences. Slight differences in the 2H phase properties in comparison to the 6R and 1T structures are due to the small amount of In retained on the Ge lattice, which is required to access the 2H phase via flux synthesis. Lastly, the thermal expansion of the 6R germanane phase was studied along with the temperature dependent Raman spectra. The in-plane lattice constants exhibit a negative thermal expansion, while the outof-plane lattice constants have a positive thermal expansion. The wavenumber of the Raman modes decrease as a function of temperature, common in layered materials.

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ASSOCIATED CONTENT

S Supporting Information *

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b04990. Rietveld and Le Bail refinements for 6R, 2H, 1T GeH, EuGe2, α-CaGe2, and β-CaGe2; temperature dependent lattice constants of Ge standard (PDF)

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REFERENCES

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Svilen Bobev: 0000-0002-0780-4787 Joshua E. Goldberger: 0000-0003-4284-604X Author Contributions

The manuscript was written primarily by N.D.C. and J.E.G. with contributions from all authors. Notes

The authors declare no competing financial interest. 1341


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