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

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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 Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b04990 • Publication Date (Web): 02 Feb 2018 Downloaded from http://pubs.acs.org on February 3, 2018

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Chemistry of Materials

Synthesis of 1T, 2H and 6R Germanane Polytypes Nicholas D. Cultrara1, Yaxian Wang,2 Maxx Q. Arguilla1, Michael R. Scudder1, Shishi Jiang1, Wolfgang Windl2, Svilen Bobev3, Joshua E. Goldberger1* 1

Address: Department of Chemistry and Biochemistry, The Ohio State University, Columbus, OH, 43210, USA 2 Department of Materials Science and Engineering, The Ohio State University, Columbus, OH, 43210, USA 3 Department of Chemistry and Biochemistry, University of Delaware, Newark, DE 19716, USA 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 outof-plane c-axis.

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 towards 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 (2-layers per hexagonal unit cell), and 3R (3 layers per rhombohedral unit cell) polytypes.14 The 1T

phases are all metallic with Mo4+ and W 4+ in octahedral coordination to S. The 2H and 3R phases have Mo4+/W 4+ 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 eV 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 mallo 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, hydrogen-terminated germanane (GeH), is a 2D material which has attracted considerable interest.21-25 Germanane is a 2dimensional germanium graphane analogue in which the germanium atoms arrange in a puckered honeycomb layer and are terminated with a covalently bonded –H ligands, alternatingly above and below each Ge atom in the network. It has a high-predicted 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 general GeH has a direct band-gap of ~1.6 eV, which can be tuned from 1.4-1.7 eV by substituting the H-terminating ligand with an organic moiety, which due to their sterics and electronics strain

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the germanane framework28. 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 only be prepared 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, αCaGe2, and β-CaGe2 correspond to the 1T, 2H, and 6R (6-layers per rhombohedral unit cell) stacking arrangements of germanium layers, respectively, which are separated by the divalent Eu2+ and Ca2+ cations.3133 Previous studies have solely focused on studying GeH transformed from the 6R β-CaGe2 phase. 21-24 Herein, we have successfully prepared 1T, 2H, and 6R polytypes of GeH through the reaction of EuGe2, αCaGe2, and β-CaGe2, respectively, with HCl. This shows there is retention of the stacking sequence through the topotactic deintercalation process. We have elucidated that the 2H α-CaGe2 phase, that is 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 have 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.

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 hr. The tubes were then inverted and centrifuged separating excess In from the crystalline CaGe2. 1T EuGe2 was obtained by loading stoichiometric amounts of

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europium and germanium in an alumina crucible and sealed in a fused silica tube. The sample is heated at 1050o C for 24 hours followed by cooling to room temperature for 24 hours. After synthesis, all crystals can be collected and placed in concentrated hydrochloric acid at -40o C from 5-40 days until complete reaction of the precursor phase. Appreciable efforts of increase the crystallinity of the 2H and 1T phases by varying acid concentration and temperature of deintercalation. Following the deintercalation process, samples are 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 patterns (XRD) 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 Å and 0.4141660 Å, respectively;. Rietveld refinements for 6R and 2H GeH and LeBail refinement of the 1T GeH phases were performed using GSAS 1 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 XRay Fluorescence using an Olympus X-5000 Mobile XRF System. Fourier Transform Infrared Spectra were collected with a Perkin-Elmer 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 Perkin-Elmer 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 are 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 augmentedwave (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


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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 .

-Ernzerhof (HSE06) hybrid functional39, 40 was employed for accurate description of the band structures.

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 (Figures S1-3) indicated a phase purity of 91% for βCaGe2, 89% for α-CaGe2, and 49% for EuGe2 (Table S16). The Zintl phases were then reacted at –40 oC in concentrated HCl for 1-4 weeks to transform them into the GeH phases. Figure 2b-d shows 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 (R 3 m), 2H (P63mc) and 1T (P 3 m1) unit cells, respectively. Each of these phases has the same primitive unit cell along the a-axis, but can feature 1-layer, 2-layer or 6-layer 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 (01 4 ) 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



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Figure 2. a) X-ray diffraction patterns of the precursor 6R β-CaGe2 (black), the 2H α-CaGe2 (red) and 1T EuGe2 phases (blue). X-ray diffraction patterns are shown for topotactically deintercalated b) 6R GeH, c) 2H GeH and d) 1T GeH, with the major reflections labeled. On account of the different X-ray wavelengths, the 2-theta ranges for b, c and d were chosen such that all three patterns range from 8.838 Å to 1.541 Å. Germanium, indium, and Eu3Ge5 impurity phases are denoted by (*), (+), and (ø), respectively. Dashed lines corresponding to the 2-theta positions of the 6R 012, 01 4 and 006 reflections are drawn in the 2H phase to accentuate their differences

data is essential to now unambiguously distinguishing 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, LeBail refinement was used to confirm the structure (Table 1, Figure S4-6, Table S7-9). Excellent refinements can only be achieved 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 Å to 3.97142 Å in 6R GeH, and the c-axis increases from 30.583 Å 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 Å to 3.9543 Å in 2H GeH, and the c-axis increases from 10.211 Å to 11.64 Å, which corresponds to a thickness of about 5.82 Å per layer. Interestingly, the 2H GeH has a much larger c-axis than 6R GeH, which will be subsequently explained by the presence of residual


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indium in the framework resulting in In–OH bonds. Finally, the a-axis of EuGe2 shrinks from 4.1035 Å to 3.9499 Å in 2H GeH, and the c-axis increases from 4.9972 Å 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.

energy cost for transforming between these different polytypes.

Table 1. Structural parameters from Rietveld (for 6R and 2H) and Le Bail (1T) refinement of deintercalated GeH.

6R GeH

2H GeH

1T GeH (LeBail)

Space Group

R 3m

P63mc

P 3 m1

a (Å)

3.97142(5)

3.9535(2)

3.9499(5)

c (Å)

33.033(5)

11.64(9)

5.776(11)

Ge(1)

x = 0, y = 0, z = 0.1772(1)

x = 0, y = 0, z = 0.00(1)

Ge(2)

x = 0, y = 0, z = 0.3418(1)

x = 1/3, y = 2/3, z = 0.561(1)

U11=U22 2 (Å )

0.03372(2)

0.0413(4)

U33 (Å )

0.583(6)

0.790(6)

wRp/Rp

0.1217/0.0978

0.0737/0.0584

2

0.0215/ 0.0178

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 30o from each another. The 6R polymorph consists of 6 different layers having a stacking sequence that we denote as AA'BB'CC'. There is 30° 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

Figure 3. FTIR spectrum of 6R (black), 2H (red) and 1T (blue) show expected results of germanane structure.

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 redshifted by ~20 cm–1 in comparison to the 6R and 1T

Figure 4. a) Shift in Ge-H stretch using FTIR in 2H (red) compared to 1T (blue) and 6R (black) associated with heavier atom substituted on structure b) XRF confirms 2.6% indium retained by germanane.

phases (Figure 3). Additionally, there are weak vibrational modes at 770 cm–1 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 would centered at 650 cm-1 would arise from In-OH termination, as we have previously demonstrated that hydroxide terminating the Ge-OH framework would appear at around 850 cm–1, and Sn-OH to be centered at 560 cm–1.41, 42



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Figure 5. a) Raman b) DRA and spectroscopy of 6R (black), 2H (red) and 1T (blue) shows expected results for the deintercalated germanane structure.

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

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

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 mode for 1T and 6R GeH occurs at 301.6 and 301.8 cm–1, respectively, while the 2H phase shifts to 300.2 cm–1. Furthemore, the out-of-plane A1 mode for the 1T and 6R GeH phase 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 provides an excellent approximation of the relative band gaps for germanane materials, as Tauc-Davis Mott models are often ambiguous due to the reduced 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


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Figure 7. a) Synchrotron XRD and c) Raman spectroscopy data at different temperatures and b) temperature dependent a(triangles) and c- (circles) lattice parameters and 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.

the germanane framework itself along with the electronwithdrawing –OH termination.28, 41 The small change in band gaps of GeH depending on polytype were further confirmed using Density Functional Theory simulations. As described above, LeBail and Rietveld analysis found layer spacings of 5.78 Å, 5.84 Å, 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 With the experimental lattice constant but altering the Geframework to have different degrees of buckling and GeGe 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. In order 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 band structures. The simulated band gaps for 1T, 2H and 6R are 1.56 eV, 1.61 eV, 1.59 eV (Fig. 6). Taking into account that 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, shifting of the and 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 Xray diffraction with synchrotron radiation. Capillary mode powder diffraction patterns of 6R GeH, with an internal Ge standard werecollected 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 (Fig S7). 6R GeH exhibits an expansion of the inter-plane spacing as the temperature increases from 100 to 295 K, since the c lattice parameter increases from 32.939(4) Å 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) Å 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-ofplane 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

Conclusion

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Here we have demonstrated that each of the 3 precursor Zintl phase materials can be synthesized and subsequently deintercalated to functionalize the layers with hydrogen and hydroxide. Each structure was characterized using X-ray diffraction, showing that the stacking sequence is retained through the deintercalation process. Each structure shows the hallmark Ge-H stretches in FTIR and similar Raman vibrations, where 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 In retained on the Ge lattice, which is required to access the 2H phase via melt synthesis. Lastly, the thermal expansion of the 6R germanane phase was studies, along with the temperature dependent Raman parameters, showing the negative thermal expansion of the in-plane constants, while the out-of-plane shows a positive thermal expansion. The germanane system also shows the trend of decreasing wavenumber as a function of temperature, common in layered materials.

ASSOCIATED CONTENT Supporting Information. Rietveld and Lebail refinements and 6R, 2H and 1T GeH, as well as EuGe2, α-CaGe2, and β-CaGe2. Temperature dependent lattice constants of Ge standard. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * Email: [email protected]

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

ACKNOWLEDGMENT 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 No. DE-AC0206CH11357. Primary funding for this research was provided by NSF EFRI-1433467. Partial Funding for this research was provided by the Center for Emergent Materials: an NSF MRSEC under award number DMR-1420451. Work at the University of Delaware received financial support from the National Science Foundation, grants DMR-0743916 (CAREER) and DMR-1709813. J.E.G. acknowledges the Camille and Henry Dreyfus Foundation for partial support.

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