Synthesis, Structure, and Electronic Properties of Germanane and Layered Materials

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

Nicholas David Cultrara

Graduate Program in Chemistry

The Ohio State University

2018

Professor Joshua E. Goldberger, Advisor

Professor Shiyu Zhang

Professor Yiying Wu

Copyrighted by

Nicholas David Cultrara

2018

Abstract

In the past decade, since the isolation of a single layer of graphene, materials research has been dominated by exploration and characterization of layered materials with an interesting range of electronic, optical, spintronic and spin-orbit coupling properties. Herein, we present research into the exploration and advancement in research of two-dimensional materials, focusing on the layered germanium structure, germanane and transport properties of other exfoliatable layered Zintl phase materials. Chapter 2 focuses on the ability to add external dopants to the layered 6R structure of the germanane family.

Gallium and arsenic were successfully introduced to the precursor Zintl phase CaGe2 and

2.3% Ga and 1.1% As were retained through the topotactic deintercalation process. Single flake transport measurements show a reduction of 3-orders of magnitude in H2O containing atmosphere for the As doped samples and >4-orders of magnitude in the highest doped Ga samples in inert atmospheres. These structures were also found to be relatively stable over

1-day exposure to ambient conditions, while showing signs of oxidation between 1-4 days in ambient atmosphere. Chapter 3 focuses on the ability to selectively grow 3 different polymorphs of layered germanium Zintl phases which can be used as precursors in the topotactic deintercalation reaction to form the corresponding hydrogen terminated layered germanane structures. Both the 6R and 1T structures are obtained from annealing elements in quartz but the 2H structure must be obtained from a synthesis using an indium flux. A

ii small portion of indium is retained by the germanane following deintercalation, which causes the formation of the two layered structure. Both the 6R and 1T show remarkable similar properties, while the 2H shows a shrinking of the band gap and shifting of Raman and FTIR peaks, associated with the retention of the heavier indium on the germanium lattice. Finally the thermal expansion of the 6R phase was investigated and found to negative in the in-plane direction while the out-of-plane direction was positive. Chapter 4 focuses on the device fabrication and electronic measurements of layered materials. The role of contact resistance and contact size on highly resistive germanane was optimized to allow to electronic characterization. Bulk resistivity measurements of both the NaSn2As2 and EuSn2As2 were conducted using indium and silver epoxy contacts respectively. Both were found to behave as metals with phonon mediated resistivity, while the EuSn2As2 shows an increase of resistivity around the magnetic transition temperature, associated with the spin-scattering of conduction electrons by Eu2+ f-electrons.

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Acknowledgments

Firstly, my journey though graduate school would have been possible without the years of mentorship, especially by Prof. Joshua E. Goldberger. As a young chemist, I would never have been able to accomplish what I have accomplished in my Ohio State career without the hands on work and with the countless hours spend conceptualizing experiments and analyzing data. I would also like to thank Professor Yiying Wu and

Professor Shiyu Zhang for their scientific input to my dissertation.

Secondly, I would like to acknowledge the members of the Goldberger group and collaborators who mentored me in the finer points of being a research scientist. Dr.

Shishi Jiang, Dr. Maxx Arguilla, Mr. Chuanchuan Sun and Dr. Basant Chitara of the

Goldberger groups, Dr. Justin Young of the Johnson-Halperin group, and Mr. Bin He of the Heremans group, along with countless others in chemistry, physics, and mechanical engineering at The Ohio State. To the present and former members of the Goldberger group, thank you for the support, entertainment, and scientific conversations, especially

Mr. Tianyang Li, Mr. Zachary Baum, Mr. Rick Morasse, Dr. Ashley Wallace, Ms.

Elisabeth Bianco, Mr. Michael Scudder, Mr. Dom Ross, Mr. Fan Fan, and Mr. Ben

Redemann.

Lastly, I would like to thank the friends and family who were always there to support me through my time as a graduate student.

Vita

2008...... Kenmore West Senior High School

2012...... B.S. Chemistry, State University of New

York at Buffalo

2012 to Present ...... Graduate Research Assistant, Department of

Chemistry and Biochemistry, The Ohio

State University

Publications

1. N. D. Cultrara, M. Q. Arguilla, S. Jiang, C. Sun, M. R. Scudder, R. D. Ross, J. E.

Goldberger, (2017) “Group 13 and 15 Doping of Germanane.” Beilstein Journal

of Nanothechnology 8(1), 1642-1648

2. M. Q. Arguilla, N. D. Cultrara, M. R. Scudder, S. Jiang, R. D. Ross, J. E.

Goldberger, (2017) “Optical Properties and Raman-Active Phonon Modes in

Two-Dimensional Honeycomb Zintl Phases.” Journal of Materials Chemistry C,

2017 v

3. M. Q. Arguilla, N. D. Cultrara, Z. J. Baum, S. Jiang, R. D. Ross, J. E. Goldberger

(2017). "EuSn2As2: An Exfoliatable Magnetic Layered Zintl–Klemm Phase."

Inorganic Chemistry Frontiers, 4, 378-386

4. M. Q. Arguilla, J. Katoch, K. Krymowski, N. D. Cultrara, J. Xu, X. Xi, A. Hanks,

S. Jiang, R. D. Ross, R. J. Koch, S. Ulstrup, A. Bostwick, C. Jozwiak, E.

Rotenberg, D. McComb, J. Shan, W. Windl, R. K. Kawakami and J. E.

Goldberger (2016). "NaSn2As2: An Exfoliatable Layered van der Waals Zintl

Phase." ACS Nano 10(10): 9500-9508.

5. G. Coloyan, N. D. Cultrara, A. Katre, J. Carrete, M. Heine, E. Ou, J. Kin, S.

Jiang, L. Lindsay, N. Mingo, D. Broido, J. P. Heremans, J. E. Goldberger, L. Shi,

(2016). "Basal-plane Thermal Conductivity of Nanocrystalline and Amorphized

Thin Germanane." Applied physics letters 109(13): 131907.

6. S. Jiang, M. Q. Arguilla, N. D. Cultrara, J. Goldberger (2016). "Improved

Topotactic Reactions for Maximizing Organic Coverage of Methyl Germanane."

Chemistry of Materials 28(13): 4735-4740.

7. S. Jiang, K. Krymowski, T. Asel, M.Q. Arguilla, N. D. Cultrara, E. Yanchenko,

X. Yang, L. Brillson, W. Windl, J. E. Goldberger, (2016). "Tailoring the

Electronic Structure of Covalently Functionalized Germanane via the Interplay of

Ligand Strain and Electronegativity." Chemistry of Materials 28(21): 8071-8077.

8. J. R. Young, B. Chitara, N. D. Cultrara, S. Jiang, F. Fan, E. Johnston-Halperin, J.

E. Goldberger, (2015). "Water Activated Doping and Transport in Multilayered

Germanane Crystals." Journal of Physics: Condensed Matter 28(3): 034001.

9. S. Jiang,* M. Q. Arguilla,* N. D. Cultrara,* J. E. Goldberger, (2014).

"Covalently-controlled Properties by Design in Group IV Graphane Analogues."

Accounts of Chemical Research 48(1): 144-151. (*equal contribution)

Fields of Study

Major Field: Chemistry

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Table of Contents

Abstract ...... ii

Acknowledgments...... iv

Vita ...... v

Publications ...... v

Fields of Study ...... vii

Table of Contents ...... viii

List of Tables ...... xi

List of Figures ...... xii

Chapter 1: Introduction ...... 1

1.1 Introduction ...... 1

1.2 Topotactic Synthesis ...... 6

1.3 Single and Few-Layer Thick Materials ...... 13

1.4 Covalently Modifiable Building Blocks ...... 16

1.5 Tuning the Electronic Structure ...... 21

1.6 Thermal and Air Stability ...... 30

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1.7 Conclusion and Outlook ...... 34

1.8 References ...... 34

Chapter 2: Group 13 and 15 Doping of Germanane ...... 41

2.1 Introduction ...... 41

2.2 Results and Discussion ...... 43

2.3 Conclusion ...... 59

2.4 Experimental ...... 60

2.5 References ...... 62

Chapter 3: Synthesis of 1T, 2H and 6R Germanane Polytypes ...... 65

3.1 Introduction ...... 65

3.2 Experimental Methods ...... 68

3.2.1 Synthesis ...... 68

3.2.2 Characterization ...... 69

3.2.3 Electronic Structure Calculations ...... 70

3.3 Results and Discussion ...... 70

3.4 Thermal Parameters of the 6R Phase ...... 85

3.5 Conclusion ...... 88

3.6 References ...... 89

3.7 Supplemental Information ...... 100

Chapter 4: Electronic characterization of layered systems ...... 109

4.1 Introduction ...... 109

4.2 Highly resistive measurements of germanane ...... 111

4.3 Bulk crystal measurements of NaSn2As2 ...... 113

4.4 Bulk crystal measurements of EuSn2As2 ...... 115

4.5 Conclusions ...... 118

4.6 References ...... 119

Chapter 5: Conclusions and Outlook ...... 122

Bibliography ...... 125

Chapter 1 ...... 125

Chapter 2 ...... 131

Chapter 3 ...... 136

Chapter 4 ...... 142

List of Tables

Table 1. The band gaps of different sp3-hybridized Group IV elements in bulk and in 2D.

...... 27

Table 2. Structural parameters from Rietveld (for 6R and 2H) and Le Bail (1T) refinement of deintercalated GeH ...... 74

Table 3. Crystal data and refinement results for EuGe2...... 103

Table 4. Fractional atomic coordinates and isotropic displacement parameters based on the refined EuGe2 structure...... 104

Table 5. Crystal data and refinement results for α-CaGe2 ...... 105

Table 6. Fractional atomic coordinates and isotropic displacement parameters based on the refined α-CaGe2 structure...... 106

Table 7. Crystal data and refinement results for β-CaGe2 ...... 107

Table 8. Fractional atomic coordinates and isotropic displacement parameters based on the refined β-CaGe2 structure...... 108

List of Figures

Figure 1 Model of GeH. View from (a) the (100) and (b) the (001) directions. (Ge: blue and H: black)...... 4

Figure 2 Schematic illustration of topotactic deintercalation of (a) CaGe2 to (b) GeH (Ca: yellow, Ge: purple and H: black). Optical images of (c) CaGe2 and (d) GeH crystals with

12 select crystals on a 1 mm grid graph paper. Powder XRD of (e) CaGe2 and (f) GeH. ... 9

Figure 3. (a) A single (111) plane of crystalline germanium, representing a single layer of

GeH. The distance of the germanium atoms on a certain colored ring from the central germanium atom corresponds to the same color peak in (b). (b) PDFs of GeH and Ge.

The starred peaks correspond to the interactions between germanium atoms in different layers.29 (c,d) Ultra-STEM images of two (c) and three (d) layers of GeH. The top insets are Fourier transforms of the corresponding image confirming the hexagonal unit cell. The bottom insets are molecular models of the corresponding images. The two-layer

GeH in (c) has alternatively stacked AB GeH layers, while the three-layer structure in (d) has a third defective layer rotated by 30o...... 10

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Figure 4. (a) Topotactic deintercalation of CaGe2-2xSn2x to Ge1-xSnxH1-x(OH)x in HCl

(Ca: yellow, Ge: blue, H: black, O: red and Sn: green). (b) FTIR and (c) Raman spectra

30 of Ge1-xSnxH1-x(OH)x (x = 0 to 0.09)...... 12

Figure 5. AFM images of (a) an exfoliated single layer of GeH and (b) GeH thin film after deintercalation of a 5 nm thick CaGe2 film grown via MBE on Ge(111). Inset in (a) is an optical micrograph of the single-layer flake.12, 33 ...... 15

Figure 6. (a) Model of GeCH3. (Ge: blue, C: black and H: grey) (b) Optical images of

GeCH3 crystals with select crystals on 1 mm grid graph paper. (c) Single-crystal XRD pattern of GeCH3 collected down the [001] zone axis. (d) Powder XRD pattern of GeH and GeCH3. The starred peaks correspond to diffraction reflections of an internal Ge

13 standard. (e) FTIR spectra of GeH, GeCH3, Ge CH3 and GeCD3. The intensities of the four spectra are multiplied by 0.5 from 400-900 cm-1.13 ...... 18

Figure 7. (a) Powder XRD pattern (b) and FTIR spectra of GeCH3 (black), GeCH2CH3

(blue), and GeCH2CH=CH2 (red)...... 20

Figure 8. DFT simulations of the electronic band structure of (a) silicene and (b) silicane and the density orbitals at the high symmetry k-points: (A) Si-H σ*, (B) Si-Si σ*, (C) and

(D) Si-Si σ states. Inset in (a) is the hexagonal Brillouin zone.36 ...... 23

Figure 9. Electronic band structure of an isolated single layer of GeH calculated using

HSE-06 theory including spin–orbit coupling predicting a 1.56 eV direct band gap. The hole and electron effective masses for each extrema are indicated in red.12 ...... 26

Figure 10. (a) DRA spectra of Ge1-xSnxH1-x(OH)x (x = 0 to 0.09) plotted in terms of the Kubelka-Munk function illustrating the consistent red shift in the absorption onset.30

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(b) Sn-dependent optical band gap of Ge1-xSnxH1-x(OH)x alloys obtained via a linear

30 approximation of the absorption edge. (c) DRA spectra of GeH compared to GeCH3 and

(d) Absorption and PL spectra of GeCH3 with the actual PL observed in isopropyl alcohol as an inset.13 ...... 29

Figure 11. (a) FTIR spectra of GeH after exposure to air for up to 60 days highlighting the absence of any Ge-O vibrational mode. (b) XPS spectra of GeH exposed to air for up to five months, followed by 0.5 nm Ar etch. DRA spectra of (c) GeH and (d) GeCH3

12, 13 annealed in Ar/H2 gas at various temperatures...... 33

Figure 12. Schematic over-representation of doped a) CaGe2 and b) GeH after detintercalation. Red represents the dopant atom, blue is germanium, yellow is calcium and black is hydrogen. The number of dopants depicted here is purposefully inflated for visual effect. c) Powder XRD, d) Raman spectra, and e) FTIR spectra of GeH (black),

2.3% Ga:GeH (red), and 1.1% As:GeH (gray). The starred peaks in the XRD show residual germanium...... 45

Figure 13. XRF calibration curve of Ga:GeH (red) and As:GeH (gray). The stars represent measured concentrations from XRF. Inset shows XRF spectrum of highest dopes samples of Ga:GeH (red) and As:GeH. Asterisks denotes a silicon escape peak from the major Ge Kα peak...... 48

Figure 14. a) EDX spectrum of 2.3% Ga:GeH platelet show in b) SEM. EDX map of c) gallium (yellow) and d) germanium (red) show retention and uniform distribution though germanane crystal...... 50

xiv

Figure 15. a) Ge 2p3/2 Ge wafer (black) with surface oxide and 0 day exposed 2.3%

Ga:GeH (red) and 1.1% As:GeH (gray). b) Time dependent XPS of 2.3% Ga:GeH and c,d) 1.1% As:GeH at 0 day (black), 1 day (yellow), 4 days (green), and 8 day (red) ...... 52

Figure 16. a) Diagram and b) photo of typical device for 2-probe I-V measurements.

Representative curves of 2.3% Ga:GeH (red) in c) vacuum and d) air and 1.1% As:GeH

(gray) in e) vacuum and f) air...... 55

Figure 17. Sheet resistance of GeH (black), 1.1% As:GeH (gray), and 2.3% Ga:GeH (red) in vacuum (unfilled) and air (filled). Each doped data point represents 20-30 device measurements...... 58

Figure 18. a) X-ray diffraction patterns of the 6R β-CaGe2 (black), the 2H α-CaGe2 (red) and 1T EuGe2 phase (blue). X-ray diffraction patterns are shown for 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 and 1T-phase to accentuate their differences...... 71

Figure 19. Projections of each precursor Zintl phase and deintercalated phase. Interlayer atomic bonding of germanium (purple) is observed in each phase, while ionic bonding is observed though the planes, with cationic calcium (yellow) and europium (magenta) and the corresponding hydrogen (pink) terminated structures...... 73

Figure 20. FTIR spectrum of 6R (black), 2H (red) and 1T (blue) show expected results of germanane structure ...... 76

Figure 21. 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 ...... 79

Figure 22. a) Raman b) DRA and spectroscopy of 6R (black), 2H (red) and 1T (blue) shows expected results for the deintercalated germanane structure...... 81

Figure 23. DFT calculations of a) 1T b) 2H and c) 6R GeH. The band gaps at and A are shown...... 82

Figure 24. a) Synchrotron XRD and c) Raman spectroscopy 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 ...... 87

Figure 24. Le Bail refinement results of 1T-GeH using GSAS ...... 100

Figure 25. Rietveld refinement results of 2H-GeH using GSAS ...... 101

Figure 26. Rietveld refinement results of 6R-GeH using GSAS. Phase purity 91%. .... 102

Figure 27. Powder XRD Topas refinement results of EuGe2 ...... 103

Figure 28. Powder XRD Topas refinement results of α-CaGe2 ...... 105

Figure 29. Powder XRD Topas refinement results of β-CaGe2 ...... 107

Figure 30. Germanium lattice constants as a function of temperature ...... 108

xvi

Figure 25. a) Schematic dimensions measured in b)bulk crystal devices in determining sheet resistance...... 113

Figure 26. Temperature dependent transport measurements of bulk and exfoliated

NaSn2As2 crystals down to 2.1 K ...... 115

Figure 27. Temperature-dependent resistivity of EuSn2As2 showing its metallic behavior along the temperature range with a cusp near the magnetic ordering temperature. Top left: Characteristic micrograph of a EuSn2As2 crystal in a four-probe geometry. Bottom right: high temperature resistivity which increases with increasing temperatures ...... 117

xvii

Chapter 1: Introduction

1.1 Introduction

As a result of the widespread integration of semiconductor technology into all facets of life, the Group IV semiconductors, silicon and germanium, are the most important and ubiquitous materials of the current era. Not only are they the workhorse materials of transistor technology, silicon and germanium are the most prevalent materials employed in photovoltaics1 and photodetectors2, and have attracted considerable attention as thermoelectric energy generators3, 4. Still, the neverending push towards device miniaturization calls for the need to understand the nature of these materials when reduced below the nanoscale. The creation of single-atom thick layers provides an avenue for the discovery of new phenomena and properties that can potentially overcome some inherent limitations in the parent three-dimensional (3D) semiconductors. For example, the indirect nature of silicon and germanium’s band gap limits their efficiency in optoelectronics, and prevents their implementation into light emitting applications.

1 Parts of this chapter were reproduced from Jiang, S.,* Arguilla, M. Q.,* Cultrara, N. D.,* and Goldberger,

J. E. Covalently-Controlled Properties by Design in Group IV Graphane Analogues. Acc. Chem. Res., 48,

Graphene’s discovery5 has shown that it is possible to prepare single atom thick layers of a two-dimensional (2D) material, and as a consequence, numerous methods6 have been developed to facilitate the understanding of the unique properties that emerge in single layers. Graphene is a single layer of graphite, and is a comprised of a π-bonded honeycomb lattice of carbon atoms. Graphene, with its linear dispersion at the K point and massless

Dirac Fermions, has unique properties like high carrier mobilites (~200,000 cm2 V-1 s-1), leading to the observation of the quantum hall effect at room temperature, as well as high thermal conductivity, and an exceptional mechanical strength.7-9 Nevertheless, the fact that this high mobility state only appears as a result of the linear Fermi-Dirac dispersion of carbon’s half-filled 2pz orbitals and semimetallic zero band gap, limits the ability to readily integrate graphene into current semiconductor technology, which require materials with band gaps for optimal performance. For example, the lack of a band gap prevents graphene transistors to have large ratios in current between the on and off state. Functionalization of graphene to make hydrogen-terminated graphene, or graphane, opens a sizable band gap,

2 -1 -1 but dramatically decreases the carrier mobility to 10 cm V s , by bonding to the C 2pz orbitals thereby eliminating the Fermi Dirac state.10 While graphene is a fascinating and promising material, the limitations of its electronic structure has inspired researchers to explore other 2D materials beyond graphene.

An entire field of research has emerged investigating other similar 2D Van der Waals solids.6 These materials allow for manual exfoliation to single and few-layers, breaking the weak interlayer interactions while maintaining the strong in-plane bonding. While there is an ever expanding class of these materials, this account will primarily focus on the Group

IV (silicon, germanium, and tin) analogues of graphane. These structures are comprised of

2D puckered honeycomb networks of sp3-hybridized Group IV atoms, and are terminated with hydrogen or other ligands (Figure 1).

Figure 1 Model of GeH. View from (a) the (100) and (b) the (001) directions. (Ge: blue and H: black).

Although they all belong to Group IV, heavier elements like Si, Ge, and Sn do not readily form π-bonds. This arises from their larger atomic size, which increases their bond distances, thereby reducing overlap between nearest neighbor π-bonding p orbitals. In other words, each Si/Ge/Sn atom would preferentially bond to another atom or ligand rather than form a π-bond with its neighbor. Since every atom in the 2D network has a covalently bound ligand, the identity of this ligand can provide a versatile synthetic handle for tuning the electronic structure and properties of this class of materials. For example, with the appropriate surface terminating ligand these 2D materials can feature direct band gaps,11-

14 potentially enhancing silicon and germanium’s performance in photovoltaics, photodetectors, light-emitting diodes, and lasers. Furthermore, these puckered honeycomb networks are structurally analogous to Si and Ge(111) surfaces, potentially allowing the use of established surface functionalization chemistries15-19 to modify the ligand.

In this account, we highlight the various routes towards synthesizing these Group IV graphane analogues. We show that this system can be covalently modified with various organic functional groups. We also discuss the influence of the surface-terminating ligand and the main group element on the electronic structure and stability of these 2D materials.

Finally, we provide an outlook on this emerging class of covalently modifiable device building blocks.

1.2 Topotactic Synthesis

To date, there have been no synthetic routes for preparing Group IV graphane analogues from small molecule precursors due to the lack of a mechanism for controlling growth in

2D. However, the layered Zintl phases, CaSi2, CaGe2, and BaSn2, are comprised of

- - - puckered honeycomb [Si ]n, [Ge ]n, and [Sn ]n graphane-like layers held together by the

Group II [M2+] cations (Figure 2a). Consequently, the preparation of Group IV graphane analogues relies on developing soft chemical processes that can topotactically deintercalate the M2+ cations while maintaining the structure and covalently terminating the anionic

Group IV layers.

The topotactic deintercalation of CaSi2 using HCl can be traced back to Wöhler in the

1860’s, Kautsky in the 1920’s, and the structure and properties were partially resolved in

20-24 the 1980’s and 1990’s. Here, it was reported that siloxene (Si6H3(OH)3) was

o preferentially formed at temperatures greater than 0 C and layered polysilane (Si6H6-

o 11, 23, 24 x(OH)1-x (x<1)) formed at -30 C. These structures are silicon graphane analogues, or silicanes, with -H/-OH, and -H terminal substituents, respectively. Compared to the indirect band gap of crystalline silicon (1.1 eV), siloxene has a direct band gap at 2.4 eV with strong photoluminescence (PL).24 Other silicon graphane analogues terminated with organic functional groups were also reported in the past decade that feature PL ranging from 2.7 to 2.9 eV.25-27 However, all these reactions rely on the topotactic deintercalation of CaSi2 in aqueous HCl, which readily produces partially OH-terminated SiHx(OH)1-x due to the significantly stronger Si-O bond (800 kJ/mol) compared with the Si-H bond (300 kJ/mol). This ambiguity in surface functionalization convolutes efforts to correlate the

6 effects of surface functionalization on the optoelectronic properties of these single-atom thick semiconductors. In contrast, the difference of bond strength is much smaller between

Ge-O (660 kJ/mol) and Ge-H (320 kJ/mol) and furthermore, any native germanium oxide or hydroxide termination is readily dissolved in aqueous HCl, thereby producing pure germanane (GeH). Indeed it was reported by Brandt and Stutzmann that CaGe2 thin films grown on germanium wafers can be topotactically deintercalated to form GeH, with little surface oxidation.28

Recently our group has synthesized for the first time millimeter-scale crystals of GeH via

12 the topotactic deintercalation of large CaGe2 single crystals in aqueous HCl (Figure 2).

2+ - Here, Ca is removed via the formation of a soluble CaCl2 species and the anionic [Ge ]n layer is terminated by H atoms. The X-ray diffraction (XRD) confirms that the layered hexagonal germanium lattice is maintained, and an increase in the interlayer distance occurs (5.1 Å to 5.5 Å) upon replacing the Ca2+ with two Ge-H bonds (Figure 2e,f). The large full-width half maximum of all the peaks that contain any c-axis reflections indicates that there exists a significant amount of disorder in the c-axis, which is common in layered materials. Pair distribution function (PDF) analysis collected from synchrotron measurements directly confirms the honeycomb 2D network of germanium atoms.29

Compared with the PDF of crystalline Ge, GeH has systematic absences at 5.66 Å, 7.35 Å, and 8.95 Å (Figure 3) that arise in 3D crystalline Ge. In Ge, these peaks correspond to Ge–

Ge pairs between atoms in different (111) layers. All other peaks can be indexed to the Ge–

Ge pairs within a single Ge(111) plane. The interlayer disorder of GeH prevents the observation of scattering between any interlayer Ge–Ge pairs. Furthermore, the 2D honeycomb lattice can be visualized via aberration corrected scanning transmission

7 electron microscopy (STEM) images of exfoliated two- and three-layer sheets. Finally, the most conclusive technique for determining the nature of the ligand bonded to each germanium atom is Fourier transform infrared spectroscopy (FTIR). Every vibrational mode observed in the FTIR corresponds to a Ge-H bond, with no peaks corresponding to any Ge-O vibrational mode. The identity of each mode was readily verified with deuterium labeling.12

Figure 2 Schematic illustration of topotactic deintercalation of (a) CaGe2 to (b) GeH (Ca: yellow, Ge: purple and H: black). Optical images of (c) CaGe2 and (d) GeH crystals with

12 select crystals on a 1 mm grid graph paper. Powder XRD of (e) CaGe2 and (f) GeH.

Figure 3. (a) A single (111) plane of crystalline germanium, representing a single layer of

GeH. The distance of the germanium atoms on a certain colored ring from the central germanium atom corresponds to the same color peak in (b). (b) PDFs of GeH and Ge. The starred peaks correspond to the interactions between germanium atoms in different layers.29 (c,d) Ultra-STEM images of two (c) and three (d) layers of GeH. The top insets are Fourier transforms of the corresponding image confirming the hexagonal unit cell. The bottom insets are molecular models of the corresponding images. The two-layer GeH in

(c) has alternatively stacked AB GeH layers, while the three-layer structure in (d) has a third defective layer rotated by 30o.

The deintercalation of Zintl phases that contain multiple Group IV elements enables the synthesis of alloy graphane analogues. We have been able to substitute up to 9% of the Ge

30 atoms with Sn (Figure 4a) in the precursor Zintl phase (CaGe2-2xSn2x). Deintercalation in aqueous HCl produces a 2D honeycomb network where tin is OH-terminated while germanium remains H-terminated (Figure 4b). Raman spectroscopy confirms that alloying occurs, as the shifts in both the in-plane phonon (E2) and cross-plane (A1) vibrational modes are consistent with the expected differences based on the changes in the reduced mass (Figure 4c).

Figure 4. (a) Topotactic deintercalation of CaGe2-2xSn2x to Ge1-xSnxH1-x(OH)x in HCl (Ca: yellow, Ge: blue, H: black, O: red and Sn: green). (b) FTIR and (c) Raman spectra of Ge1-

30 xSnxH1-x(OH)x (x = 0 to 0.09).

1.3 Single and Few-Layer Thick Materials

The synthesis of large millimiter-scale flakes of GeH enables the isolation of single- and few-layer thick sheets via mechanical exfoliation using PDMS and scotch tape (Figure 5a), by adapting the procedure developed by Frindt.31, 32 These single-layer flakes were exfoliated onto 110 nm thick and 300 nm thick SiO2/Si substrates, which provide suitable optical contrast. However, mechanical exfoliation is a labor-intensive and non-scalable process for producing single and few-layer thick materials. The ability to synthesize precise layer thicknesses of these Group IV graphane analogues on Si and Ge substrates, would enable their seamless integration into existing semiconductor fabrication infrastructure.

This can be achieved by first epitaxially growing the precursor Zintl phases onto Si and Ge substrates. The a,b-parameters of the precursor CaGe2 Zintl phase closely match the spacing of the Ge(111) surface to less than <0.5%. This enables the direct epitaxial growth of CaGe2 thin films on Ge(111) wafers, which can be subsequently topotactically deintercalated to obtain GeH. Indeed, we have prepared 5 nm thick films of CaGe2 on

Ge(111) via molecular beam epitaxy (MBE), which would correspond to ~10 layers.33

These co-deposited CaGe2 thin films have grain sizes on the order of a few micrometers, which is the typical sizes of terrace formed due to the miscut of the Ge(111) growth substrate. Upon treatment in HCl, these thin films exhibit the same XRD and Raman profiles as those produce from single crystals of CaGe2. Consequently, the combination of epitaxial growth and topotactic deintercalation represents a promising and scalable route

13 for the preparation of precise layer films of Group IV graphane analogues and simplifies subsequent VLSI processing.

1.4 Covalently Modifiable Building Blocks

The presence of a covalently bound surface ligand on every atom in these Group IV graphane analogues opens up the possibility of tuning the properties by varying this surface ligand. There has been extensive work during the past few decades showing that every atom on Si and Ge(111) surfaces can be terminated with small organic substituents such as

15-18, 34 –CH3 and –CCH. In contrast to H-terminated Si/Ge(111) surfaces, which oxidize within 30 minutes of exposure to air, these organic-terminated surfaces have been shown to be resistant towards oxidation for at least 30 days.18, 34, 35 Consequently, it is easy to envision 2D derivatives of these same organic functionalized surfaces.

To these ends, we have developed a one-step metathesis approach that directly converts

CaGe2 crystals into organic-terminated germananes by topotactically reacting it with organoiodines. For instance, we have prepared ~1 mm flakes of GeCH3 (Figure 6a,b) by

13 - reacting CaGe2 with CH3I. Through this reaction, Ge anions bond to the CH3 group, and

2+ the iodide reacts with Ca to form a soluble CaI2 species, which is easily separated. Single crystal and powder XRD analysis show that the hexagonal unit cell of CaGe2 is retained and GeCH3 has a similar 2H unit cell (two GeCH3 layers per unit cell) as GeH (Figure

6c,d). The interlayer distance of GeCH3 is increased by 3.1 Å compared to GeH, which is close to the estimated increase (~2.5 Å) based on the bond length and van der Waals radii differences of these two ligands. The methyl-termination is further confirmed by FTIR measurements (Figure 6e). Compared with GeH, the intense Ge-H stretching mode at 2000

16 cm-1 is almost entirely gone, while a Ge-C stretching mode at 573 cm-1 is observed. Other vibrational modes like -CH3 stretching, bending and rocking modes are also detected. The identity of each mode can be further verified upon comparison with the FTIR spectra of

13 Ge CH3 and GeCD3.

Figure 6. (a) Model of GeCH3. (Ge: blue, C: black and H: grey) (b) Optical images of

GeCH3 crystals with select crystals on 1 mm grid graph paper. (c) Single-crystal XRD pattern of GeCH3 collected down the [001] zone axis. (d) Powder XRD pattern of GeH and

GeCH3. The starred peaks correspond to diffraction reflections of an internal Ge standard.

13 (e) FTIR spectra of GeH, GeCH3, Ge CH3 and GeCD3. The intensities of the four spectra are multiplied by 0.5 from 400-900 cm-1.13

This one-step metathesis method is a general route for prepare organic ligand terminated germananes. By substituting CH3I with other organoiodine reagents like CH3CH2I and

CH2=CHCH2I, we have prepared CH3CH2Ge and CH2=CHCH2Ge, respectively. The interlayer spacing is expected to increase by 3.5 and 6.2 Å, upon replacing -H in GeH with

-CH2CH3 and -CH2CH=CH2, respectively. This is in close agreement with the increases in interlayer spacing of 3.7 and 5.8 Å observed via XRD (Figure 7a). In FTIR spectra, Ge-C stretching can be detected in all three spectra along with the near elimination of the Ge-H stretching mode. All the other vibrational modes can be assigned to the corresponding organic functional groups. The observation of H-C=C- stretching and bending modes at

-1 3076 and 1626 cm further confirms the termination with -CH2CH=CH2. The versatility of this reaction scheme enables the grafting of functional ligands with tunable polarity, reactivity, and mechanical strength for a wide variety of applications.

Figure 7. (a) Powder XRD pattern (b) and FTIR spectra of GeCH3 (black), GeCH2CH3

(blue), and GeCH2CH=CH2 (red).

1.5 Tuning the Electronic Structure

The rich surface functionalization chemistry allows these Group IV graphane analogues to be highly tunable electronic and optoelectronic building blocks for next generation devices.

By substituting the main group element and varying the surface ligand, one can tune the electronic structure of these materials to produce unique properties that do not exist in the parent 3D semiconductor structure. To understand how the presence of a surface ligand influences the electronic structure of these 2D graphane analogues, here we describe density functional theory (DFT) simulations illustrating the difference between silicene and silicane as a model system.36

Silicene is a 2D material comprised of a honeycomb arrangement of Si atoms in which every Si shares three σ and one π bond with the three neighboring Si atoms.36, 37 Similar to graphene, this structure exhibits Fermi-Dirac behavior at the K point on account of the half- filled 3pz orbitals (Figure 8a). Adding hydrogen as a surface terminating ligand to silicene

36- produces silicane (SiH) through the formation of a covalent bond with the Si 3pz orbital.

38 This bonding and anti-bonding interactions splits the Dirac cone at the K point thus opening a sizeable band gap (Figure 8b). The electronic band structure of silicane calculated at the the HSE-0639, 40 level predicts it to have an indirect band gap of 2.94 eV from Γ to M and a direct band gap of 3.14 eV at Γ.36 The conduction band valley at Γ is comprised of Si-H σ* states, the conduction band minimum (CBM) at M corresponds to

Si-Si σ* states, whereas the valence band maximum (VBM) corresponds to Si-Si σ states

(Figure 8b).

Figure 8. DFT simulations of the electronic band structure of (a) silicene and (b) silicane and the density orbitals at the high symmetry k-points: (A) Si-H σ*, (B) Si-Si σ*, (C) and

(D) Si-Si σ states. Inset in (a) is the hexagonal Brillouin zone.36

The band structure of GeH and SnH are closely related to SiH.36, 37 However, in the case of GeH, the CBM occurs at Γ and not M, leading to a 1.56 eV direct band gap with an

* effective electron mass of me,Γ = 0.09. This is consistent with the observed absorption onset at 1.59 eV and with the observation of PL in GeH at 1.56 eV at low temperature. The direct band gap of GeH is in sharp contrast to the 0.67 eV indirect bap gap of crystalline germanium.41 In crystalline germanium, the CBM occur in the four equivalent valleys at

the L <111> point, which has a much higher effective mass (me,L* =1.64) than the

conduction band valleys at Γ (me,Γ* = 0.041). However, since GeH can be thought of as hydrogen-terminated isolated (111) sheets of germanium, we are effectively eliminating the L wavevector in the Ge Brillouin zone, resulting in a material that has a direct gap and a considerably higher electron mobility. Electron mobility is inversely proportional to effective mass. We calculated from first-principles the phonon-limited electronic mobility for an isolated single layer of GeH obtaining a room temperature mobility of ~18,000 cm2

V-1 s-1. This 5× increase in electron mobility from bulk Ge (3,900 cm2 V-1 s-1) is consistent with the reduced electron effective mass in GeH. These Group IV graphane analogues also feature significantly larger band gaps compared to the parent 3D material (Table 1).

The 2D sp3-hybridized Sn is predicted to be a topological insulator when terminated by electronegative groups such as -OH and various halides, yet remains as a trivial insulator when terminated by smaller less electronegative ligands such as -H.14 This topological phase emerges when the Sn 5s σ* bands drop below the 4px and 4py VBM, which are split

24 on account of the large spin-orbit coupling in Sn. Consequently, 2D tin is predicted to exhibit the quantum spin Hall effect.

Figure 9. Electronic band structure of an isolated single layer of GeH calculated using HSE-

06 theory including spin–orbit coupling predicting a 1.56 eV direct band gap. The hole and electron effective masses for each extrema are indicated in red.12

Table 1. The band gaps of different sp3-hybridized Group IV elements in bulk and in 2D.

3D sp3 (eV) 2D sp3 (eV) C 5.48 (indirect)42 3.5* (-H, direct)43 Si 1.12 (indirect)44 2.4 (-OH/-H, direct)11 2.94* (-H, indirect)36 Ge 0.67 (indirect)41 1.59 (-H, direct)12 13 1.71 (-CH3, direct) Sn 0±0.05 (β-, ~0.30* (-H,-halides, direct)45 direct)14 *Theoretical

One way to tune the electronic structure of these 2D materials is through alloying different

Group IV elements in the framework. We have systematically tuned the band gap of GeH from 1.59 eV down to 1.38 eV (Figure 10a,b) through alloying up to 9% of Sn into the

CaGe2-2xSn2x lattice, and topotactically deintercalating the lattice with aqueous HCl, based off of diffuse reflectance absorption (DRA) measurements.30 Increasing the Sn percentage in the lattice can potentially lead to band gaps down to 0.3 eV, which would open up their application as tunable photodetectors suitable for telecommunications.

Figure 10. (a) DRA spectra of Ge1-xSnxH1-x(OH)x (x = 0 to 0.09) plotted in terms of the

Kubelka-Munk function illustrating the consistent red shift in the absorption onset.30 (b)

Sn-dependent optical band gap of Ge1-xSnxH1-x(OH)x alloys obtained via a linear

30 approximation of the absorption edge. (c) DRA spectra of GeH compared to GeCH3 and

(d) Absorption and PL spectra of GeCH3 with the actual PL observed in isopropyl alcohol as an inset.13

Aside from alloying the main group element, another route that can be used to tune the properties of these materials is through the variation of the surface ligand. Because the

CBM at Γ corresponds to the Group IV-ligand σ* antibond, this energy level can be raised or lowered depending on the electronegativity of the surface terminating ligand. Since the ligand also influences the geometry of the entire lattice, this also leads to change in the bond lengths and bond angles in the 2D network, significantly affecting the rest of the band structure. Therefore, substituting one ligand for another will affect the entire electronic band structure, and can significantly change the observed band gap and direct/indirect nature. A systematic experimental study comparing the electronic and optical properties of uniformly functionalized 2D Group IV graphane analogues is needed to foster a deeper understanding of the influence of surface ligands on the electronic structure of these materials.

Experimentally, we have demonstrated that the band gap of GeH can be increased by 0.1

13 eV (Figure 10c) by replacing the -H with -CH3. Furthermore, very intense band edge PL is observed in GeCH3 (Figure 10d) with a quantum yield of 0.2%. In contrast to other layered materials like MoS2, this PL is independent of layer thickness, thus obviating the need for large area single layers for practical devices. Also, the PL indicates that these 2D

Group IV graphane analogues have great potential in light-emitting applications. In summary, the combination of alloying and covalent functionalization provides an unprecedented level of control over electronic structure in a 2D crystal, rendering it a versatile platform for a multitude of electronic and optoelectronic applications.

1.6 Thermal and Air Stability

The potential utility of these Group IV graphane analogues for any functional device strongly hinges on their air and temperature stability. It has been established that silicon- based 2D systems such as SiH readily oxidize upon exposure to air.26, 38 GeH, however, shows a remarkable resistance to oxidation.12 The lack of oxidation in the bulk material is observed in FTIR measurements (Figure 11a) after air exposure to up to 60 days where the

46 -1 absence of different GeOx vibrational modes from 800 to 1000 cm indicate that the bulk of GeH is unchanged. On the other hand, X-Ray Photoelectron Spectroscopy (XPS) (Figure

11b) is the most sensitive technique in determining the presence of oxidation on the surface.

+/3+ After five months of exposure to air, a Ge2 shoulder emerges at 1219.3 eV (29.7%

+/3+ Ge2 ) indicating surface oxidation is prevalent. After Ar ion etching the top 0.5 nm (~1

+/3+ +/3+ 12 layer), the Ge2 peak almost completely disappears with 10.1% Ge2 remaining.

Together, the XPS and FTIR results suggest that GeH is resilient towards oxidation and only the surface layer slowly oxidizes.

The thermal stability of these 2D materials strongly depend on the identity of the surface terminating ligand. We observed that in GeH, the emergence of thermal-induced amorphization begins at 75 oC.12 We have observed that measurements focused on changes in properties are the most sensitive methods for detecting amorphization, compared to bulk structural analyses like XRD and Raman spectroscopy. The absorption of the new and emerging amorphous germanium species lower than the 2D germanane band gap is more readily observable than the gradual disappearance of a crystalline phase in XRD or Raman.

Upon annealing at 75 oC and above, there is an increasing red shift in the absorption onset, consistent with the formation of amorphous GeH, and at higher temperatures, amorphous

Ge (Figure 11c).12 This amorphization process was further confirmed with temperature

29 dependent PDF measurements. In contrast, GeCH3 starts to amorphize at a significantly higher annealing temperature, at around 250 oC (Figure 11d).13 In summary, the oxidation resistance of germanane and the improvement in thermal stability by methyl termination makes these materials attractive candidates for next generation devices, and can be potentially robust enough to withstand the demands for fabrication and operation.

12, 13 in Ar/H2 gas at various temperatures.

1.7 Conclusion and Outlook

These Group IV graphane analogues are real materials that can be synthesized as robust single crystals in gram-scale quantities, exfoliated into single layers, and prepared on conventional VLSI substrates as few-layer thin films. Compared to all other 2D van der

Waals solids, these Group IV graphane analogues offer the capability for covalent surface termination, providing a versatile handle for tailoring the structure, stability, and electronic properties in these single atom-thick materials. Such exquisite control over the material properties makes these system attractive candidates for a multitude of applications such as sensing, optoelectronics, and thermoelectrics and also offers the potential for new physical phenomena such as the quantum spin Hall effect. Leveraging previous established semiconductor surface functionalization routes will enable the creation of entire libraries of organic ligand terminated 2D materials. Furthermore, the ability to vary band offsets and band gaps in a single 2D sheet solely by selecting different surface functional groups provides the opportunity to define and study 1D interfaces within a single 2D layer by spatially patterning the surface termination. Overall, this new class of 2D materials is posed to have a great impact not only in the traditional sectors of nanoscience, but also in opening up a new research paradigm in covalently-controlled properties by design.

1.8 References

1. Lalović, B.; Kiss, Z.; Weakliem, H., A hybrid amorphous silicon photovoltaic and

thermal solar collector. Solar Cells 1986, 19 (2), 131-138.

2. Auston, D. H.; Lavallard, P.; Sol, N.; Kaplan, D., An amorphous silicon

photodetector for picosecond pulses. Appl. Phys. Lett. 1980, 36 (1), 66-68.

3. Joshi, G.; Lee, H.; Lan, Y.; Wang, X.; Zhu, G.; Wang, D.; Gould, R. W.; Cuff, D.

C.; Tang, M. Y.; Dresselhaus, M. S.; Chen, G.; Ren, Z., Enhanced Thermoelectric

Figure-of-Merit in Nanostructured p-type Silicon Germanium Bulk Alloys. Nano

Lett. 2008, 8 (12), 4670-4674.

4. Hochbaum, A. I.; Chen, R.; Delgado, R. D.; Liang, W.; Garnett, E. C.; Najarian,

M.; Majumdar, A.; Yang, P., Enhanced thermoelectric performance of rough

silicon nanowires. Nature 2008, 451 (7175), 163-167.

5. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Zhang, Y.; Dubonos, S.

V.; Grigorieva, I. V.; Firsov, A. A., Electric field effect in atomically thin carbon

films. Science 2004, 306 (5696), 666-669.

6. Butler, S. Z.; Hollen, S. M.; Cao, L.; Cui, Y.; Gupta, J. A.; Gutierrez, H. R.; Heinz,

T. F.; Hong, S. S.; Huang, J.; Ismach, A. F.; Johnston-Halperin, E.; Kuno, M.;

Plashnitsa, V. V.; Robinson, R. D.; Ruoff, R. S.; Salahuddin, S.; Shan, J.; Shi, L.;

Spencer, M. G.; Terrones, M.; Windl, W.; Goldberger, J. E., Progress, challenges,

and opportunities in two-dimensional materials beyond graphene. ACS Nano 2013,

7 (4), 2898-2926.

7. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.;

Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A., Two-dimensional gas of massless

Dirac Fermions in graphene. Nature 2005, 438 (7065), 197-200.

8. Balandin, A. A.; Ghosh, S.; Bao, W.; Calizo, I.; Teweldebrhan, D.; Miao, F.; Lau,

C. N., Superior Thermal Conductivity of Single-Layer Graphene. Nano Lett. 2008,

8 (3), 902-907.

9. Lee, C.; Wei, X.; Kysar, J. W.; Hone, J., Measurement of the Elastic Properties and

Intrinsic Strength of Monolayer Graphene. Science 2008, 321 (5887), 385-388.

10. Elias, D. C.; Nair, R. R.; Mohiuddin, T. M. G.; Morozov, S. V.; Blake, P.; Halsall,

M. P.; Ferrari, A. C.; Boukhvalov, D. W.; Katsnelson, M. I.; Geim, A. K.;

Novoselov, K. S., Control of graphene's properties by reversible hydrogenation:

evidence for graphane. Science 2009, 323 (5914), 610-613.

11. Yamanaka, S.; Matsu-ura, H.; Ishikawa, M., New deintercalation reaction of

calcium from calcium disilicide. Synthesis of layered polysilane. Mater. Res. Bull.

1996, 31, 307-316.

12. Bianco, E.; Butler, S.; Jiang, S.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Stability and Exfoliation of Germanane: A Germanium Graphane Analogue. ACS

Nano 2013, 7 (5), 4414-4421.

13. Jiang, S.; Butler, S.; Bianco, E.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Improving the stability and optical properties of germanane via one-step covalent

methyl-termination. Nat. Commun. 2014, 5, 3389.

14. Xu, Y.; Yan, B.; Zhang, H.-J.; Wang, J.; Xu, G.; Tang, P.; Duan, W.; Zhang, S.-C.,

Large-gap quantum spin hall insulators in tin films. Physical Review Letters 2013,

111, 136804.

15. Knapp, D. Chemistry and Electronics of the Ge(111) Surface. Dissertation (Ph.D.),

California Institute of Technology, 2011.

16. Bansal, A.; Li, X.; Lauermann, I.; Lewis, N. S.; Yi, S. I.; Weinberg, W., Alkylation

of Si surfaces using a two-step halogenation/Grignard route. J. Am. Chem. Soc.

1996, 118 (30), 7225-7226.

17. Han, S. M.; Ashurst, W. R.; Carraro, C.; Maboudian, R., Formation of alkanethiol

monolayer on Ge (111). J. Am. Chem. Soc. 2001, 123 (10), 2422-2425.

18. Linford, M. R.; Fenter, P.; Eisenberger, P. M.; Chidsey, C. E., Alkyl monolayers

on silicon prepared from 1-alkenes and hydrogen-terminated silicon. J. Am. Chem.

Soc. 1995, 117 (11), 3145-3155.

19. Buriak, J. M., Organometallic chemistry on silicon and germanium surfaces. Chem.

Rev. 2002, 102 (5), 1271-1308.

20. Wöhler, F., Ueber Verbindungen des Siliciums mit Sauerstoff und Wasserstoff.

Liebigs. Ann. 1863, 127 (3), 257-274.

21. Kautsky, H., Über einige ungesättigte Siliciumverbindungen. Z. anorg. Chem.

1921, 117 (1), 209-242.

22. Kautsky, H.; Herzberg, G., Über das Siloxen und seine Derivate. Z. anorg. allg.

Chem. 1924, 139 (1), 135-160.

23. Weiss, A.; Beil, G.; Meyer, H., The Topochemical Reaction of CaSi2 to a Two-

Dimensional Subsiliceous Acid Si6H3(OH)3. Z. Naturforsch. 1980, 35b, 25-30.

24. Dettlaff-Weglikowska, U.; Hönle, W.; Molassioti-Dohms, A.; Finkbeiner, S.;

Weber, J., Structure and optical properties of the planar silicon compounds

polysilane and Wöhler siloxene. Phys. Rev. B 1997, 56 (20), 13132-13140.

25. Okamoto, H.; Kumai, Y.; Sugiyama, Y.; Mitsuoka, T.; Nakanishi, K.; Ohta, T.;

Nozaki, H.; Yamaguchi, S.; Shirai, S.; Nakano, H., Silicon nanosheets and their

self-assembled regular stacking structure. Journal of the American Chemical

Society 2010, 132 (8), 2710-2718.

26. Nakano, H.; Nakano, M.; Nakanishi, K.; Tanaka, D.; Sugiyama, Y.; Ikuno, T.;

Okamoto, H.; Ohta, T., Preparation of alkyl-modified silicon nanosheets by

hydrosilylation of layered polysilane (Si6H6). Journal of the American Chemical

Society 2012, 134 (12), 5452-5455.

27. Sugiyama, Y.; Okamoto, H.; Mitsuoka, T.; Morikawa, T.; Nakanishi, K.; Ohta, T.;

Nakano, H., Synthesis and Optical Properties of Monolayer Organosilicon

Nanosheets. J. Am. Chem. Soc. 2010, 132 (17), 5946-5947.

28. Vogg, G.; Brandt, M. S.; Stutzmann, M., Polygermyne—A Prototype System for

Layered Germanium Polymers. Adv. Mater. 2000, 12 (17), 1278-1281.

29. Jiang, S.; Bianco, E.; Goldberger, J. E., The structure and amorphization of

germanane. J. Mater. Chem. C 2014, 2 (17), 3185-3188.

30. Arguilla, M. Q., Jiang, S., Chitara, B., Goldberger, J. E., Synthesis and Stability of

Two-dimensional Ge/Sn Graphane Alloys. Submitted 2014.

31. Frindt, R.; Yoffe, A., Physical properties of layer structures: optical properties and

photoconductivity of thin crystals of molybdenum disulphide. Proc. R. Soc.

London, Ser. A. 1963, 273 (1352), 69-83.

32. Novoselov, K.; Jiang, D.; Schedin, F.; Booth, T.; Khotkevich, V.; Morozov, S.;

Geim, A., Two-dimensional atomic crystals. Proc. Natl. Acad. Sci. U. S. A. 2005,

102 (30), 10451-10453.

33. Pinchuk, I. V.; Odenthal, P. M.; Ahmed, A. S.; Amamou, W.; Goldberger, J. E.;

Kawakami, R. K., Epitaxial co-deposition growth of CaGe2 films by molecular

beam epitaxy for large area germanane. J. Mater. Res. 2014, 29 (03), 410-416.

34. He, J.; Lu, Z.-H.; Mitchell, S. A.; Wayner, D. D., Self-Assembly of Alkyl

Monolayers on Ge (111). J. Am. Chem. Soc. 1998, 120 (11), 2660-2661.

35. Nemanick, E. J.; Hurley, P. T.; Brunschwig, B. S.; Lewis, N. S., Chemical and

electrical passivation of silicon (111) surfaces through functionalization with

sterically hindered alkyl groups. J. Phys. Chem. B 2006, 110 (30), 14800-14808.

36. Restrepo, O. D.; Mishra, R.; Goldberger, J. E.; Windl, W., Tunable gaps and

enhanced mobilities in strain-engineered silicane. J. Appl. Phys. 2014, 115 (3),

033711.

37. Lew Yan Voon, L. C.; Sandberg, E.; Aga, R. S.; Farajian, A. A., Hydrogen

compounds of group-IV nanosheets. Applied Physics Letters 2010, 97 (16),

163114.

38. Nakano, H.; Mitsuoka, T.; Harada, M.; Horibuchi, K.; Nozaki, H.; Takahashi, N.;

Nonaka, T.; Seno, Y.; Nakamura, H., Soft synthesis of single-crystal silicon

monolayer sheets. Angew. Chem., Int. Ed. 2006, 45, 6303-6306.

39. Paier, J.; Marsman, M.; Hummer, K.; Kresse, G.; Gerber, I. C.; Angyan, J. G.,

Screened hybrid density functionals applied to solids. J. Chem. Phys. 2006, 124,

154709.

40. Heyd, J.; Scuseria, G. E.; Ernzerhof, M., Hybrid functionals based on a screened

Coulomb potential. [Erratum to document cited in CA139:042043]. J. Chem. Phys.

2006, 124, 219906.

41. Macfarlane, G. G.; Roberts, V., Infrared absorption of germanium near the lattice

edge. Phys. Rev. 1955, 97, 1714.

42. Clark, C. D.; Dean, P. J.; Harris, P. V., Intrinsic edge absorption in diamond. Proc.

R. Soc. London, Ser. A. 1964, 277, 312-329.

43. Lebegue, S.; Klintenberg, M.; Eriksson, O.; Katsnelson, M. I., Accurate electronic

band gap of pure and functionalized graphane from GW calculations. Phys. Rev. B:

Condens. Matter Mater. Phys. 2009, 79, 245117.

44. Bludau, W.; Onton, A.; Heinke, W., Temperature dependence of the band gap of

silicon. J. Appl. Phys. 1974, 45, 1846-1848.

45. Hoechst, H.; Hernandez-Calderon, I., Angular resolved photoemission of indium

antimonide(001) and heteroepitaxial films of a-Sn(001). Surf. Sci. 1983, 126, 25-

31.

46. Rivillon, S.; Chabal, Y. J.; Amy, F.; Kahn, A., Hydrogen passivation of germanium

(100) surface using wet chemical preparation. Appl. Phys. Lett. 2005, 87, 253101.

Chapter 2: Group 13 and 15 Doping of Germanane

2.1 Introduction

Since the discovery of graphene[1], the quest to discover and measure novel two dimensional and layered materials has led to the investigation of group 14 and group 15 allotropes of graphene and graphane[1-14], transition metal dichalcogenides[15-19], and layered van der Waals materials[20-22]. Germanane, a hydrogen-terminated germanium graphane analogue, has garnered considerable attention in the 2D materials on account of its direct band gap[5, 23, 24], large predicted electron mobility, and the ability to controllably tune the optoelectronic properties via covalent modification with surface ligands.[3, 24-28] While the room temperature electron mobility of germanane has been predicted to be >18,000 cm2 V-1s-1, transport measurements on non-extrinsically doped crystals were highly resistive, indicating the need of extrinsic dopants to access lower resistivity devices. In previous studies, the resistivity of germanane was reduced through the incorporation of phosphorus only when activated in the presence of atmospheric water[29].

Parts of this chapter were reproduced from Cultrara, N.*, Arguilla, M., Jiang, S., Sun, C. Scudder,

M. Ross, R., Goldberger, J., Group-13 and Group-15 doping of germanane, Beilsteil J.

Recently[30], undoped germanane field effect transistors were reported with device hole mobilites ranging from 70 to 150 cm2 V-1 s-1 from room temperature to low temperature, indicating that germanane has the potential to be a viable electronic building block for 2D transistors. Together, this emphasizes the need for further control of doping behaviour in these materials.

The preparation of GeH requires the synthesis of a CaGe2 precursor phase followed by its topotactic deintercalation in HCl. GeH is a metastable phase, and begins to amorphize when annealing above 75o C. Consequently, traditional doping processes such as the direct ion-implantation of GeH cannot be used as they require a high-temperature post annealing to heal the lattice. Due to the existence of a large number of closely related layered Zintl phases with group 13 and group 15 elements that are structurally similar to CaGe2,, dopant elements can be partially substituted into the germanium lattice in the CaGe2 precursor[29,

31]. Providing that these elements are retained in the germanium framework after the topotactic deintercalation process, the effect on electronic transport behaviour of GeH should be appreciable. Having previously grown phosphorus-doped GeH (P:GeH)[29] using this method, here, we explore whether other group 13 and 15 elements including (Al,

Ga, As and Sb) as dopants onto the germanane framework, and how these dopants affect the stability and electronic properties of GeH.

Herein, we show that Ga and As can be doped into the CaGe2 precursor phase and are retained on the germanane lattice after topotactic deintercalation. Using X-ray fluorescence (XRF) and X-ray Photoelectron Spectroscopy (XPS) we show that up to 1.1% and 2.3% of As and Ga, respectively, can be substituted onto the germanane lattice. In

42 contrast to pristine GeH, these materials begin to oxidize between 24 and 96 hours in ambient atmosphere. In both cases, the incorporation of more dopants produced lower sheet resistances in H2O, containing ambient conditions, while only the gallium doped samples continue to show dopant activation under vacuum and H2O-free conditions.

2.2 Results and Discussion

First, we explored whether crystals of CaGe2 that were doped with Al, Ga, As, and Sb at

0.1% could be synthesized (Figure 12a). Of these dopants only Al, Ga, and As were able to be successfully incorperated onto the CaGe2 framework. After topotactic deintercalation in HCl, GeH platelets doped with Ga and As were successfully obtained (Figure 12b), while the Al-doped CaGe2 crystals disintegrated into small micron-sized particles not sutable for bulk transport measurements. Subsequently, we synthesized Ga and As doped

CaGe2 at 0.1-9% atomic substitutions, however, CaGe2 crystals only formed with less than

3% of added dopant. Figure 12c shows the powder X-ray diffraction (XRD) pattern for undoped GeH reported by Bianco et al.[25], and the highest doped Ga:GeH and As:GeH samples. All the deintercalated phases can be indexed[32] to a 6-layer rhombohedral unit cell with a=3.97 Å and c=33.22 Å lattice parameters while showing no significant difference between the phases and other peaks indiciative of impurity phases are observed.

The asterisks labeled peaks show residual germanium in the sample. The Raman spectra

(Figure 1d) of these doped crystals all exhibit A1 (out-of-plane) and E2 (in-plane) modes at

-1 -1 228 cm and 302 cm respectively, with no change in peak location, shape or A1/E2 intensity relative to undoped GeH. Fourier Transform Infrared Spectra of these samples

(FTIR) also further show clear spectroscopic signatures for the formation of GeH. An extremely strong Ge–H stretching mode is observed at ∼2000 cm-1 as well as characteristic

-1 wagging modes at 570, 507, and 475 cm , and the edge/defect Ge–H2 defect modes at 770 and 820 cm-1. While no additional features indicative of As–H or Ga–H, were observed, the small concentration of Ga and As makes it impossible to completely determine the actual chemical environment of these dopants using a bulk technique such as IR spectroscopy.

Ga:GeH (red), and 1.1% As:GeH (gray). The starred peaks in the XRD show residual germanium.

The retention and concentration of Ga and As dopants in the lattice was determined for each system using XRF (Figure 13). A calibration curve usng the ratio of Ga/As Kα to Ge

Kα was prepared in standards comprised of elemental Ge and As or Ga2O3. These measurements showed that the highest concentration of Ga in Ga:GeH to be 2.3%, and the highest concentration of As in As:GeH was 1.1%. XRF analysis of GeH synthesized with greater than 1% As substitution, always yielded an As:Ge of ~1.1% in GeH indicating that this is the maximum amount of As that can be substituted on CaGe2. The lack of any other distinguishing phase in the XRD suggests that the Ga and As as part of the germanane lattice.

Kα peak.

Scanning electron microscopy (SEM) with energy-dispersive X-ray spectroscopy (EDX) provided further verification of the incorporation of dopant atoms into GeH. As a representative example, the EDX spectrum when the electron beam is localized on a single

2.3% Ga:GeH platelet, shows the presence of both Ge K and Ga K peaks (Figure 14a).

Figure 14b,c,d shows an SEM image, the map of Ga Ka signal, and Ge Ka signal, respectively, of a corner of a Ga:GeH platelet. These EDX maps show that there is a uniform distribution of gallium and germanium throughout the germanane crystal. This confirms the retention of the Ga dopant into the germanane lattice.

Figure 14. a) EDX spectrum of 2.3% Ga:GeH platelet show in b) SEM. EDX map of c) gallium (yellow) and d) germanium (red) show retention and uniform distribution though germanane crystal.

X-ray photoelecton spectroscopy (XPS) measurements confirmed dopant retention in the lattice, and elucidate the local chemical environment of the dopant (Figure 15). The Ge

2p3/2 peak for the 2.3% Ga:GeH and 1.1% As:GeH occur at 1217.7 and 1217.6 eV (Figure

15a) respectively, which is relatively close to undoped GeH at 1217.8 eV[25]. These Ge

1+ 2p3/2 energies are indicative of a Ge oxidation state. For comparison, a Ge(111) wafer

0 + 4+ having surface oxide contains Ge 2p3/2 peaks at 1216.3 for Ge , and oxidized Ge2 -Ge peaks that range from 1218.2 to 1220.6 eV. Figure 15b shows the Ga 3d5/2 and Ge 3d5/2

XPS spectra for the 2.3% Ga:GeH crystals after exposure to ambient conditions for 0-8 days. Immediately after synthesis (0 days air exposure) the Ga and Ge 3d5/2 peaks can be fit to single peaks at 19.9 eV and 30.3 eV, respectively. These binding energies occur in the range expected for Ga3+,[33] and Ge1+[25] oxidation states. Minimal changes are observed after one day of exposure to air. However, after four days of ambient air exposure, the XPS spectra shows the emergence of higher binding energy Ge 3d5/2 peaks that are indicative of surface oxidation. Fitting the higher energy spectra shows that 83%

1+ of Ge at the surface is not oxidized. Conversely, the binding energies of the Ga 3d5/2 peak and the Ga 2p3/2 peak that occurs at 1117.5 eV[34] do not change after exposure to the ambient atmosphere. As a dopant in GeH, Ga is bonded to three more electronegative Ge atoms, and locally exists in an electron-deficient state. Consequently, minimal changes in the Ga XPS spectra would be expected if Ga:GeH were to become oxidized to form Ga2O3.

Figure 15. a) Ge 2p3/2 Ge wafer (black) with surface oxide and 0 day exposed 2.3%

Ga:GeH (red) and 1.1% As:GeH (gray). b) Time dependent XPS of 2.3% Ga:GeH and c,d)

1.1% As:GeH at 0 day (black), 1 day (yellow), 4 days (green), and 8 day (red)

Figure 15c shows the XPS spectra for As:GeH after exposure to air for 1-8 days. Again, immediately after synthesis (0 days air exposure) the As and Ge 3d5/2 peaks can be fit to single peaks at 41.8 eV and 30.0 eV, respectively. Also, in As:GeH minimal changes one the observed Ge 2p3/2 peak after one day of exposure to air. However, surface oxidation is prevalent after four days, evident from the emergence of a higher binding energy peak, indicative of an oxidized Ge 3d5/2 envirnment. Fitting the intensity of the peaks shows 84% of Ge remains as Ge1+. The similarity of the change in Ge binding energy for both As:GeH and Ga:GeH implies that the rate of oxidation of Ge for both samples are similar. In contrast to Ga:GeH, the changes in the As 2p3/2 binding energy (Figure 15d) indicates that significant oxidation of As occurs. The As 2p3/2 peak centered at 1323 eV in as-grown

As:GeH starts disappearing in favor of a 1326.1 eV[35] oxidized state.

To probe the effect of dopants to the electronic transport transport of single crystal flakes of Ga:GeH and As:GeH were carried out with contacts fabricated using a shadow mask technique (Figure 16b). 2-probe I-V measurements were measured on single crystals with device geometries that typically featured 25 µm channel length, 2-4 mm in width and 5-20

µm in thickness (Figure 16b). After exploring numerous metals, nearly ohmic contacts

(under vacuum) to Ga:GeH were observed using 100 nm Au as a contact metal.

Furthermore, the highest ambient and vacuum conductivities in As:GeH were achieved when contacting with Ag (80 nm)/Au (20 nm). Since Au, a higher work function metal is needed to make ohmic contacts for Ga:GeH, as opposed to Ag for As:GeH, suggests that

Ga and As are likely to act as p-type and n-type dopants, respectively. . I-V measurements were carried out with a direct probe contact to each metal pad, and measured from a range of -5 to 5 V. 20-30 devices were fabricated for each doping concentration. Each measurement was normalized to a sheet resistance. Undoped GeH exhibited sheet resistances approaching the noise limit of instrumentation on the order of ~1015 Ω/☐ in both vacuum and ambient lab atmosphere conditions, similar to previous studies[29].

Figure 16. a) Diagram and b) photo of typical device for 2-probe I-V measurements.

Representative curves of 2.3% Ga:GeH (red) in c) vacuum and d) air and 1.1% As:GeH

(gray) in e) vacuum and f) air.

Figure 16c and Figure 16 d show a representative I-V plot for 2.3% Ga:GeH under vacuum and air, respectively. The I-V plot for the device measured under vacuum in Figure 16c shows ohmic contact behavior, with a typical sheet resistance of 9.5x1010 Ω/☐. The IV behavior when measured in H2O-containing atmosphere such as air, is highly hysteretic, non-ohmic, and with much higher current. Previously, we had shown that for P:GeH[29],

H2O containing atmospheres are necessary to activate the phosphorus, group 15, dopants.

For Ga:GeH, the presence of atmospheric water, while significantly increasing conductivity, also introduces signficantly non-linear behaviour in the IV plots, suggesting that H2O plays an additional role in these gallium samples in addition to dopant activation.

This also makes it difficult to extract an accurate value for sheet resistance, leading to the ommiting of its use as a metric for these samples. Regardless of the chemical state of the dopant atom in H2O atmosphere, under vacuum there is a systematic decrease of sheet resistance for Ga:GeH with increasing amounts of Ga doping. Specifically, with 0.08%,

0.14%, and 2.3% of gallium doping, the sheet resistance drops to 8.4x1012, 1.3x1012, and

9.5x1010 Ω/☐, respectively (Figure 17). This demonstrates a marked improvement of >4- orders of magnitude over undoped GeH with our most highly doped samples.

Conversely, As:GeH does not exhibit any dopant activation under vacuum. Figure 16e shows a representative I-V plot of a 1.1% As:GeH device under vacuum. With the application of ± 5V, <1 pA of current is observed, which is again at the actual noise-limit of our instrumentation. However, in air, there is at least a 3-order of magnitude increase in conductivity. Figure 16f shows a represent I-V plot of As:GeH in ambient conditions.

In contrast to Ga:GeH, the I-V plot is linear, with minimal hysteresis. The IV behavior both in air and vacuum is similar to what was previously reported to occur in P:GeH, another group 15-dopant. In air, the average sheet resistance for 30 As:GeH devices was

5.0x1011 Ω/☐ (Figure 17).

Figure 17. Sheet resistance of GeH (black), 1.1% As:GeH (gray), and 2.3% Ga:GeH (red) in vacuum (unfilled) and air (filled). Each doped data point represents 20-30 device measurements.

2.3 Conclusion

Here we have demonstrated that gallium and arsenic can be incorporated into the precursor

CaGe2 Zintl phase and are retained in the 2D germanium framework after the topotactic deintercalation process. These dopants do not significantly change the structure of germanane. These doped materials are stable in ambient atmosphere conditions for at least

24 hours but start to oxidize between 1-4 days. The introduction of Ga and As to the lattice decreases the resistance in ambient conditions with large amounts of hysteresis, suggesting that the presence of water can activate these dopants. As was previously observed with

P:GeH,[29] As:GeH is highly resistive under vacuum, indicating the presence of water is required to activate group 15 dopants. In contrast, Ga:GeH exhibited decreased sheet resistances in vacuum by over 4-orders of magnitude, which was proportional to the amount of gallium and exhibited minimal hysteretic behaviour. This indicates that the dopant activated state in Ga:GeH is stable under vacuum, enabling robust electronic properties through encapsulation. Overall, this work provides a pathway to dope germanane and enable future explorations of electronic devices.

2.4 Experimental

Single crystalline platelets of doped and undoped GeH were synthesized using methods adapted from those reported previously.[25] For undoped GeH, Stoichiometric amounts of calcium (Acros, 99%) and germanium(Acros, 99.999%) were sealed in quartz tubes under vacuum of <60 mTorr. The sample was annealed at 950o C for 18 hours, and slowly cooled to room temperature over the course of 2-10 days. CaGe2 crystals were collected and placed in concentrated HCl, at -40o C for >8 days resulting in flakes having lateral dimensions of 5 x 5 mm. To prepare extrinsically doped CaGe2, elemental aluminum

(Johnson Matthey Electronics 99.9%), gallium (Acros 99.9%), arsenic (Sigma 99.999%), or antimony (Strem 99%) were stoichiometrically replaced germanium in the initial calcium and germanium mixture. Again, these materials were sealed in quartz tubes under vacuum, and annealed following the same procedures as undoped germanane. The Sb containing experiments resulted in the formation of a mixture of different phases, none of which were structurally similar to any known layered CaGe2 polymorph. Subsequently, the

o single crystals of the x:CaGe2 (x= Al, Ga, As) were placed in -40 C HCl for at least 8 days, until deintercalation was complete and a lack of crystalline CaGe2 peaks appeared in the

XRD. The products were first rinsed with deionized water followed by rinsing with methanol, three times each. The crystals were collected via slow centrifugation and subsequently dried in vacuum.

The structure of doped GeH were confirmed using capillary X-Ray Diffraction using Cu

Kα1 λ=1.54 nm on a Bruker D8 Powder X-ray Diffractometer. XRD was performed using with finely ground powders packed in capillaries. Raman spectroscopy was used to

60 confirm vibrational modes using a Renishaw InVia Raman equipped with a CCD detector exciting with a 633 nm (He-Ne red laser). The Relative elemental composition was measured using X-Ray Fluorescence using an Olympus X-5000 Mobile XRF System.

SEM and EDX was performed using FEI Helios Nanolab 600 Dual Beam Focussed Ion

Beam/Scanning Electron. X-Ray Photoelectron Spectroscopy was performed using a

Kratos Axis Ultra X-ray photoelectron spetrometer with a monochromatic Aluminum X- ray gun. Samples were mounted in a glovebox and then stored in ambient atmosphere conditions for 1, 4 and 8 days to determine the stability in air. Fourier Transform Infrared

Spectra were collected with a Perkin-Elmer Frontier Dual-Range FIR/MidIR spectrometer that was loaded in an Ar-filled glovebox and using an attenuated total internal reflection

(ATR) sample geometry.

Electrical properties were measured in a top-contact device geometry, where metal contacts were first deposited via E-beam deposition using a shadow mask resulting in a 25 µm channel length. The contact metals used for undoped GeH and As:GeH were 80nm/20nm

(Ag/Au). Ga:GeH device contacts were prepared using 100nm Au. Contact materials were selected after testing with multiple metals and selecting the metal which gave the highest and most linear I-V. Additionally, four probe measurements indicate that the contact resistance for the 2.3% Ga:GeH is at least 2 orders of magnitude lower than the resistance of the material, indicating that contact resistance is neglible for these samples. All devices were stored in an Ar-filled glovebox until atmospheric measurements were carried out.

Electronic measurements were conducted using a Keithley 4200-SCS attached to a Lake

Shore Cryonics Inc. probe station. Two-probe current-voltage measurements were performed in both vacuum (~10-4 mbar) and under ambient conditions in the dark.

2.5 References

1. Novoselov, K. S.; Geim, A. K.; Morozov, S.; Jiang, D.; Katsnelson, M.; Grigorieva, I.; Dubonos, S.; Firsov, A., Nature (London, U.K.) 2005, 438 (7065), 197-200.

2. Garcia, J. C.; de Lima, D. B.; Assali, L. V.; Justo, J. F., J. Phys. Chem. C 2011, 115 (27), 13242-13246.

3. Jiang, S.; Arguilla, M. Q.; Cultrara, N. D.; Goldberger, J. E., Acc. Chem. Res. 2015, 48 (1), 144-151.

4. Pulci, O.; Gori, P.; Marsili, M.; Garbuio, V.; Del Sole, R.; Bechstedt, F., Europhys. Lett. 2012, 98 (3), 37004.

5. Vogg, G.; Brandt, M.; Stutzmann, M., Adv. Mater. 2000, 12 (17), 1278-1281.

6. Vogt, P.; De Padova, P.; Quaresima, C.; Avila, J.; Frantzeskakis, E.; Asensio, M. C.; Resta, A.; Ealet, B.; Le Lay, G., Phys. Rev. Lett. 2012, 108 (15), 155501.

7. Wei, W.; Dai, Y.; Huang, B.; Jacob, T., Phys. Chem. Chem. Phys. 2013, 15 (22), 8789-8794.

8. Bonaccorso, F.; Colombo, L.; Yu, G.; Stoller, M.; Tozzini, V.; Ferrari, A. C.; Ruoff, R. S.; Pellegrini, V., Science (Washington, DC, U. S.) 2015, 347 (6217), 1246501.

9. De Padova, P.; Vogt, P.; Resta, A.; Avila, J.; Razado-Colambo, I.; Quaresima, C.; Ottaviani, C.; Olivieri, B.; Bruhn, T.; Hirahara, T., Appl. Phys. Lett. 2013, 102 (16), 163106.

10. Zhang, R. W.; Zhang, C. W.; Ji, W. X.; Li, S. S.; Hu, S. J.; Yan, S. S.; Li, P.; Wang, P. J.; Li, F., New J. Phys. 2015, 17 (8), 083036.

11. Zhang, R. W.; Zhang, C. W.; Ji, W. X.; Li, S. S.; Yan, S. S.; Hu, S. J.; Li, P.; Wang, P. J.; Li, F., Sci. Rep. 2016, 6.

12. Zhao, H.; Ji, W. X.; Zhang, C. W.; Li, P.; Zhang, S. F.; Li, F.; Wang, P. J.; Li, S. S.; Yan, S. S., J. Mater. Chem. C 2017, 5 (10), 2656-2661.

13. Zhang, R. W.; Ji, W. X.; Zhang, C. W.; Li, S. S.; Li, P.; Wang, P. J., J. Mater. Chem. C 2016, 4 (10), 2088-2094.

14. Wang, Y. P.; Ji, W. X.; Zhang, C. W.; Li, P.; Li, F.; Ren, M. J.; Chen, X. L.; Yuan, M.; Wang, P. J., Sci. Rep. 2016, 6, 20342.

15. Radisavljevic, B.; Kis, A., Nat. Mater. 2013, 12 (9), 815-820.

16. Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, i. V.; Kis, A., Nat. Nanotechnol. 2011, 6 (3), 147-150.

17. Wang, Q. H.; Kalantar-Zadeh, K.; Kis, A.; Coleman, J. N.; Strano, M. S., Nat. Nanotechnol. 2012, 7 (11), 699-712.

18. Mak, K. F.; Lee, C.; Hone, J.; Shan, J.; Heinz, T. F., Phys. Rev. Lett. 2010, 105 (13), 136805.

19. Zhang, R. W.; Ji, W. X.; Zhang, C. W.; Li, P.; Wang, P. J., 2016, arXiv:1607.02973 [cond-mat.mtrl-sci]. arXiv.org e-Print archive. https://arxiv.org/abs/1607.02973 (accessed 07/05/2017) 2016.

20. Arguilla, M.; Cultrara, N.; Baum, Z.; Jiang, S.; Ross, R.; Goldberger, J., Inorg. Chem. Front. 2017, (4), 378–386.

21. Arguilla, M. Q.; Katoch, J.; Krymowski, K.; Cultrara, N. D.; Xu, J.; Xi, X.; Hanks, A.; Jiang, S.; Ross, R. D.; Koch, R. J., ACS Nano 2016, 10 (10), 9500-9508.

22. Geim, A. K.; Grigorieva, I. V., Nature (London, U.K.) 2013, 499 (7459), 419-425.

23. Dávila, M.; Xian, L.; Cahangirov, S.; Rubio, A.; Le Lay, G., New J. Phys. 2014, 16 (9), 095002.

24. Butler, S. Z.; Hollen, S. M.; Cao, L.; Cui, Y.; Gupta, J. A.; Gutiérrez, H. R.; Heinz, T. F.; Hong, S. S.; Huang, J.; Ismach, A. F., ACS Nano 2013, 7 (4), 2898-2926.

25. Bianco, E.; Butler, S.; Jiang, S.; Restrepo, O. D.; Windl, W.; Goldberger, J. E., ACS Nano 2013, 7 (5), 4414-4421.

26. Jiang, S.; Bianco, E.; Goldberger, J. E., J. Mater. Chem. C 2014, 2 (17), 3185-3188.

27. Jiang, S.; Butler, S.; Bianco, E.; Restrepo, O. D.; Windl, W.; Goldberger, J. E., Nat. Commun. 2014, 5, 4389.

28. Jiang, S.; Krymowski, K.; Asel, T.; Arguilla, M. Q.; Cultrara, N. D.; Yanchenko, E.; Yang, X.; Brillson, L. J.; Windl, W.; Goldberger, J. E., Chem. Mater. 2016, 28 (21), 8071-8077.

29. Young, J. R.; Chitara, B.; Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Fan, F.; Johnston-Halperin, E.; Goldberger, J. E., J. Phys.: Condens. Matter 2015, 28 (3), 034001.

30. Madhushankar, B.; Kaverzin, A.; Giousis, T.; Potsi, G.; Gournis, D.; Rudolf, P.; Blake, G.; van der Wal, C.; van Wees, B., 2D Mater. 2017, 4 (2), 021009.

31. Arguilla, M. Q.; Jiang, S.; Chitara, B.; Goldberger, J. E., Chem. Mater. 2014, 26 (24), 6941-6946.

32. Luo, X.; Zurek, E., J. Phys. Chem. C 2015, 120 (1), 793-800.

33. Leonhardt, G.; Berndtsson, A.; Hedman, J.; Klasson, M.; Nilsson, R.; Nordling, C., Phys. Status Solidi B 1973, 60 (1), 241-248.

34. Wagner, C., Faraday Discuss. Chem. Soc. 1975, 60, 291-300.

35. Taylor, J. A., J. Vac. Sci. Technol. (N. Y., NY, U. S.) 1982, 20 (3), 751-755.

Chapter 3: Synthesis of 1T, 2H and 6R Germanane Polytypes

3.1 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 numbers1, 2, 3, 4, 5, 6, 7. 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, 4, 8, 9, 10. This has led to a more recent efforts focused on stacking and understanding how to couple neighboring layers together to create exotic physical phenomena.11, 12, 13, 14 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

65 per rhombohedral unit cell).15 The 1T phases are all metallic with Mo4+ and W4+ in octahedral coordination to S. 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,15, 16 and computational predictions indicate

17 the 2H- and 3R-MoS2 polytypes have 1.29 eV and 1.33 eV band gap, respectively. The

2H and 3R phases also feature slightly different bond lengths, different Raman modes, exciton binding energies, and temperature dependent lattice constant changes9, 15, 17, 18.

Another 2D material system that has attracted considerable interest is hydrogen- terminated germanane (GeH)19, 20, 21, 22, 23. Germanane is a 2-dimensional 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 electrode24, 25. In general GeH has a direct band-gap of

~1.6 eV, which can be controllably 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 the germanane framework26. 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 has 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 elements27, whose structure and bonding can be rationalized using the Zintl-

Klemm concept28. 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 (6-layers per rhombohedral unit cell) stacking arrangements of germanium layers, respectively, which are separated by the divalent Eu2+ and Ca2+ cations

29, 30, 31. Previous studies have solely focused on studying GeH transformed from the 6R

19, 20, 21, 22 β-CaGe2 phase.

Herein, we have 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 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.

3.2 Experimental Methods

3.2.1 Synthesis

All reactions were carried out in evacuated fused silica tubes which were loaded in an argon filled glovebox using methods adopted previously.19, 29, 30 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 <60 mTorr. These tubes were then heated at 950o C for 18 hr, followed by slowly cooling to room temperature over the course of 2-10 days. 2-H CaGe2 was grown via indium flux, by adding calcium, germanium, and indium in a 1:1:10 ratio inside of small crucible, with a quartz wool plug on top. Tubes were heated at 960o C for 20 hr, followed by a slow cool to 600o C over >50 hr. The tubes were then inverted and centrifuged separating excess In from the crystalline CaGe2. 1-T

EuGe2 was obtained by loading stoichiometric amounts of 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 from 5-40 days until complete reaction of the precursor phase. 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.

3.2.2 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

D8 Diffractometer employing Cu Kα1 radiation with  = 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 exciting with a 785 nm diode laser 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 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.

3.2.3 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)32, 33. For structures where relaxations were necessary, we used PBE34 projector augmented-wave

(PAW) potentials35 and Grimme’s DFT-D2 method36 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 functional37, 38 was employed for accurate description of the band structures.

3.3 Results and Discussion

First, polycrystalline 6R and 2H CaGe2 and 1T EuGe2 were synthesized for their subsequent deintercalation into germanane. Figure 18a 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 (Figure 25, Figure 26,

Figure 27) indicated a phase purity of 91% for β-CaGe2, 89% for α-CaGe2, and 49% for

EuGe2 (Table 3, Table 4, Table 5, Table 6, Table 7, Table 8).

Figure 18. a) X-ray diffraction patterns of the 6R β-CaGe2 (black), the 2H α-CaGe2 (red) and 1T EuGe2 phase (blue). X-ray diffraction patterns are shown for 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 and 1T-phase to accentuate their differences.

The Zintl phases were then reacted at –40 oC in concentrated HCl for 1-4 weeks to transform them into the GeH phases. Figure 18 b-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-3m), 2H (P63mc) and 1T (P-3m1) 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 19).

Consequently, each phase can be identified using distinct reflections in the diffraction pattern. For instance, the 1T phase has both the 011 reflection for which the analogous 6R reflection (016) is forbidden, as well as the 110 reflection which is forbidden in the 2H structure. Similarly, the 6R-structure features the (012) and (01-4) reflections, which do not exist in the other unit cells. The GeH reflections also show significant broadening compared to the precursor Zintl phases, which is in-

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

6R-GeH 2H-GeH 1T-GeH Space Group R-3m P63mc P-3m1 a (A) 3.97142(5) 3.9535(2) 3.9499(5) c (A) 33.033(5) 11.64(9) 5.776(11) 0 , 0 , Ge(1) 0.17716(9) 0 , 0 , 0.00(1) 0 , 0 , 1/3 , 2/3 , Ge(2) 0.34183(8) 0.561(1) U11 = 2 U22 (Å ) 3.372(2) 0.413(4) 2 U33 (Å ) 58.3(6) 79.0(6) wRp/Rp 0.1217/0.0978 0.0737/0.0584 0.0215/0.0178

74 dicative of smaller crystalline domain sizes. Due to the broadness and presence of overlapping reflections, synchrotron powder diffraction data is essential for unambiguously distinguishing between the 6R and 2H GeH phases.29

The structures of these phases were confirmed using Rietveld analysis (Table 2, Figure 28,

Figure 29, Figure 30, Table 3, Table 4, Table 5, Table 6, Table 7). Excellent refinements can only be achieved when the GeH space group is the same as the parent Zintl

Figure 20. FTIR spectrum of 6R (black), 2H (red) and 1T (blue) show expected results of germanane structure

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 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.

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 19). 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

o are two different puckered honeycomb layers. These layers are rotated by 30 from each another. The 6R polymorph consists of 6 different layers having a stacking sequence that

77 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 energy cost for transforming between these different polytypes.

Infrared spectroscopy was used to further confirm hydrogen termination for the three GeH polytypes. In 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 21).

Figure 21. 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

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. We have previously demonstrated that hydroxide terminating the Ge-OH framework would appear between 850 cm–1, and Sn-OH to be centered at 560 cm–1.39 We attribute this peak to the presence In doping on the framework, which is retained upon deintecalation. XRF measurements confirm that after deintercalation, the 2H GeH phase has 0.027 moles of indium per mole of germanium in the framework (Figure 21). It is important to point out that no In peaks were observed in the XRD (Figure 18), 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. Indeed the broadness and red-shifting of the 2H Ge-H Infrared stretching frequency relative to the 6R and 1T phases is indicative of a heavier atom on the germanane framework. Such changes have been previously observed when Sn is substituted onto the germanium lattice.

Figure 22. a) Raman b) DRA and spectroscopy of 6R (black), 2H (red) and 1T

(blue) shows expected results for the deintercalated germanane structure.

Figure 23. DFT calculations of a) 1T b) 2H and c) 6R GeH. The band gaps at and A are shown.

The Raman spectroscopy 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 22), 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

–1 300.2 cm . 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 22). 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 the germanane framework itself along with the electron-withdrawing –OH termination.26, 39

The small change in band gaps of GeH depending on polytype were further confirmed using Density Functional Theory simulations. As described above, 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

83 were too large by more than 0.5 eV. 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, which then are the structures for which the final band structures were calculated.

The obtained band gaps for 1T, 2H and 6R are 1.56 eV, 1.61 eV, 1.59 eV (Figure 23).

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  point of the Brillouin zone while 1T has a direct band gap at the A point. For 6R, the energy gaps at  and A are only slightly different.

The presence of In on the 2H GeH framework terminated with -OH explains why it has the lowest 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 was also observed when small percentages of Sn (~3%) were

39 substituted onto the CaGe2 the framework. 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.

3.4 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

85 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 24. 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 measurements40 (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 24). 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. 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 temperatures41.

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

Additionally, the Raman spectra were collected at each temperature for which diffraction

–1 was obtained (Figure 24). The out-of-plane A1 vibration located at 227-229 cm and the

–1 in-plane E2 vibration located at 301-304 cm both decrease in wavenumber as the temperature of the system is increased (Figure 24). This relationship is directly correlated to the increase in Ge-Ge and Ge-H bond lengths as the temperature increases. Similar

42, 43, 44 trends are widely observed in other 2D materials such as MoS2 and phosphorene .

3.5 Conclusion

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. 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.

3.6 References

1. Li, L.; Yu, Y.; Ye, G. J.; Ge, Q.; Ou, X.; Wu, H.; Feng, D.; Chen, X. H.; Zhang, Y.,

Black phosphorus field-effect transistors. Nat. Nanotechnol. 2014, 9 (5), 372-377.

2. Novoselov, K. S.; Geim, A. K.; Morozov, S.; Jiang, D.; Katsnelson, M.; Grigorieva,

I.; Dubonos, S.; Firsov; AA, Two-dimensional gas of massless Dirac fermions in graphene. arXiv preprint cond-mat/0509330 2005.

3. Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, i. V.; Kis, A., Single-layer

MoS2 transistors. Nat. Nanotechnol. 2011, 6 (3), 147-150.

4. Splendiani, A.; Sun, L.; Zhang, Y.; Li, T.; Kim, J.; Chim, C.-Y.; Galli, G.; Wang,

F., Emerging photoluminescence in monolayer MoS2. Nano Lett. 2010, 10 (4), 1271-1275.

5. Butler, S. Z.; Hollen, S. M.; Cao, L.; Cui, Y.; Gupta, J. A.; Gutiérrez, H. R.; Heinz,

T. F.; Hong, S. S.; Huang, J.; Ismach, A. F., Progress, challenges, and opportunities in two- dimensional materials beyond graphene. ACS Nano 2013, 7 (4), 2898-2926.

6. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.;

Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A., Two-dimensional gas of massless Dirac fermions in graphene. Nature (London) 2005, 438 (7065), 197-200.

7. Molle, A.; Goldberger, J.; Houssa, M.; Xu, Y.; Zhang, S.-C.; Akinwande, D.,

Buckled two-dimensional Xene sheets. Nat. Mater. 2017, 16 (2), 163-169.

8. Mak, K. F.; Lee, C.; Hone, J.; Shan, J.; Heinz, T. F., Atomically thin MoS2: a new direct-gap semiconductor. Phys. Rev. Lett. 2010, 105 (13), 136805.

9. Schönfeld, B.; Huang, J.; Moss, S., Anisotropic mean-square displacements (MSD) in single-crystals of 2H-and 3R-MoS2. Acta Crystallogr., Sect. B: Struct. Crystallogr.

Cryst. Chem. 1983, 39 (4), 404-407.

10. He, Q.; Zeng, Z.; Yin, Z.; Li, H.; Wu, S.; Huang, X.; Zhang, H., Fabrication of

Flexible MoS2 Thin‐Film Transistor Arrays for Practical Gas‐Sensing Applications. Small

2012, 8 (19), 2994-2999.

11. Chang, C.-Z.; Zhang, J.; Feng, X.; Shen, J.; Zhang, Z.; Guo, M.; Li, K.; Ou, Y.;

Wei, P.; Wang, L.-L., Experimental observation of the quantum anomalous Hall effect in a magnetic topological insulator. Science 2013, 340 (6129), 167-170.

12. Yu, R.; Zhang, W.; Zhang, H.-J.; Zhang, S.-C.; Dai, X.; Fang, Z., Quantized anomalous Hall effect in magnetic topological insulators. Science 2010, 329 (5987), 61-

64.

13. Geim, A. K.; Grigorieva, I. V., Van der Waals heterostructures. Nature (London)

2013, 499 (7459), 419-425.

14. Lee, G.-H.; Yu, Y.-J.; Cui, X.; Petrone, N.; Lee, C.-H.; Choi, M. S.; Lee, D.-Y.;

Lee, C.; Yoo, W. J.; Watanabe, K., Flexible and transparent MoS2 field-effect transistors on hexagonal boron nitride-graphene heterostructures. ACS Nano 2013, 7 (9), 7931-7936.

15. Song, I.; Park, C.; Choi, H. C., Synthesis and properties of molybdenum disulphide: from bulk to atomic layers. RSC Adv. 2015, 5 (10), 7495-7514.

16. Wilson, J.; Yoffe, A., The transition metal dichalcogenides discussion and interpretation of the observed optical, electrical and structural properties. Adv. Phys. 1969,

18 (73), 193-335.

17. Coutinho, S. S.; Tavares, M. S.; Barboza, C. A.; Frazão, N. F.; Moreira, E.;

Azevedo, D. L., 3R and 2H polytypes of MoS2: DFT and DFPT calculations of structural, optoelectronic, vibrational and thermodynamic properties. J. Phys. Chem. Solids 2017, 111

(Supplement C), 25-33.

18. Lee, J.-U.; Kim, K.; Han, S.; Ryu, G. H.; Lee, Z.; Cheong, H., Raman Signatures of Polytypism in Molybdenum Disulfide. ACS Nano 2016, 10 (2), 1948-1953.

19. Bianco, E.; Butler, S.; Jiang, S.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Stability and exfoliation of germanane: a germanium graphane analogue. ACS Nano 2013,

7 (5), 4414-4421.

20. Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Sun, C.; Scudder, M. R.; Ross, R. D.;

Goldberger, J. E., Group-13 and group-15 doping of germanane. Beilstein J. Nanotechnol.

2017, 8 (1), 1642-1648.

21. Jiang, S.; Bianco, E.; Goldberger, J. E., The structure and amorphization of germanane. J. Mater. Chem. C. 2014, 2 (17), 3185-3188.

22. Jiang, S.; Butler, S.; Bianco, E.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Improving the stability and optical properties of germanane via one-step covalent methyl- termination. Nat. Commun. 2014, 5, 3389.

23. Young, J. R.; Chitara, B.; Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Fan, F.;

Johnston-Halperin, E.; Goldberger, J. E., Water activated doping and transport in multilayered germanane crystals. J. Phys. Condens. Matter 2015, 28 (3), 034001.

24. Madhushankar, B.; Kaverzin, A.; Giousis, T.; Potsi, G.; Gournis, D.; Rudolf, P.;

Blake, G.; van der Wal, C.; van Wees, B., Electronic properties of germanane field-effect transistors. 2D Mater. 2017, 4 (2), 021009.

25. Serino, A. C.; Ko, J. S.; Yeung, M. T.; Schwartz, J. J.; Kang, C. B.; Tolbert, S. H.;

Kaner, R. B.; Dunn, B. S.; Weiss, P. S., Lithium-Ion Insertion Properties of Solution-

Exfoliated Germanane. ACS Nano 2017, 11 (8), 7995-8001.

26. Jiang, S.; Krymowski, K.; Asel, T.; Arguilla, M. Q.; Cultrara, N. D.; Yanchenko,

E.; Yang, X.; Brillson, L. J.; Windl, W.; Goldberger, J. E., Tailoring the Electronic

Structure of Covalently Functionalized Germanane via the Interplay of Ligand Strain and

Electronegativity. Chem. Mater. 2016, 28 (21), 8071-8077.

27. Arguilla, M.; Cultrara, N.; Baum, Z.; Jiang, S.; Ross, R.; Goldberger, J., EuSn2 As2: an exfoliatable magnetic layered Zintl–Klemm phase. Inorganic Chemistry Frontiers

2017, 4 (2), 378-386.

28. Nesper, R., The Zintl‐Klemm Concept–A Historical Survey. Z. Anorg. Allg. Chem

2014, 640 (14), 2639-2648.

29. Bobev, S.; Bauer, E. D.; Thompson, J. D.; Sarrao, J. L.; Miller, G. J.; Eck, B.;

Dronskowski, R., Metallic behavior of the Zintl phase EuGe2: combined structural studies,

92 property measurements, and electronic structure calculations. J Solid State Chem 2004, 177

(10), 3545-3552.

30. Tobash, P. H.; Bobev, S., Synthesis, structure and electronic structure of a new polymorph of CaGe2. J Solid State Chem 2007, 180 (5), 1575-1581.

31. Luo, X.; Zurek, E., Crystal Structures and Electronic Properties of Single-Layer,

Few-Layer, and Multilayer GeH. J. Phys. Chem. C 2015, 120 (1), 793-800.

32. Kresse, G.; Hafner, J., Ab initio molecular dynamics for liquid metals. Phys. Rev.

B 1993, 47 (1), 558.

33. Kresse, G.; Hafner, J., Ab initio molecular-dynamics simulation of the liquid- metal–amorphous-semiconductor transition in germanium. Phys. Rev. B 1994, 49 (20),

14251.

34. Perdew, J. P.; Burke, K.; Ernzerhof, M., Generalized gradient approximation made simple. Phys. Rev. Lett. 1996, 77 (18), 3865.

35. Blöchl, P. E., Projector augmented-wave method. Phys. Rev. B 1994, 50 (24),

17953.

36. Grimme, S., Semiempirical GGA‐type density functional constructed with a long‐ range dispersion correction. J. Comput. Chem. 2006, 27 (15), 1787-1799.

37. Heyd, J.; Scuseria, G. E.; Ernzerhof, M., Hybrid functionals based on a screened

Coulomb potential. J. Chem. Phys. 2003, 118 (18), 8207-8215.

38. Paier, J.; Marsman, M.; Hummer, K.; Kresse, G.; Gerber, I. C.; Ángyán, J. G.,

Screened hybrid density functionals applied to solids. J. Chem. Phys. 2006, 124 (15),

154709.

39. Arguilla, M. Q.; Jiang, S.; Chitara, B.; Goldberger, J. E., Synthesis and stability of two-dimensional Ge/Sn graphane alloys. Chem. Mater. 2014, 26 (24), 6941-6946.

40. Novikova, S., Thermal expansion of germanium at low temperatures. Phys. Solid

State 1960, 2 (1), 37-38.

41. Munn, R., Role of the elastic constants in negative thermal expansion of axial solids. J Phys C 1972, 5 (5), 535.

42. Taube, A.; Judek, J.; Łapińska, A.; Zdrojek, M., Temperature-Dependent Thermal

Properties of Supported MoS2 Monolayers. ACS Appl. Mater. Interfaces 2015, 7 (9), 5061-

5065.

43. Zhang, S.; Yang, J.; Xu, R.; Wang, F.; Li, W.; Ghufran, M.; Zhang, Y.-W.; Yu, Z.;

Zhang, G.; Qin, Q.; Lu, Y., Extraordinary Photoluminescence and Strong

Temperature/Angle-Dependent Raman Responses in Few-Layer Phosphorene. ACS Nano

2014, 8 (9), 9590-9596.

44. Lanzillo, N. A.; Glen Birdwell, A.; Amani, M.; Crowne, F. J.; Shah, P. B.; Najmaei,

S.; Liu, Z.; Ajayan, P. M.; Lou, J.; Dubey, M., Temperature-dependent phonon shifts in monolayer MoS2. Appl. Phys. Lett. 2013, 103 (9), 093102.

1. Li, L.; Yu, Y.; Ye, G. J.; Ge, Q.; Ou, X.; Wu, H.; Feng, D.; Chen, X. H.; Zhang, Y.,

Black phosphorus field-effect transistors. Nat. Nanotechnol. 2014, 9 (5), 372-377.

2. Novoselov, K. S.; Geim, A. K.; Morozov, S.; Jiang, D.; Katsnelson, M.; Grigorieva,

I.; Dubonos, S.; Firsov; AA, Two-dimensional gas of massless Dirac Fermions in

graphene. arXiv preprint cond-mat/0509330 2005.

3. Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, i. V.; Kis, A., Single-layer

mos2 transistors. Nat. Nanotechnol. 2011, 6 (3), 147-150.

4. Splendiani, A.; Sun, L.; Zhang, Y.; Li, T.; Kim, J.; Chim, C.-Y.; Galli, G.; Wang,

F., Emerging photoluminescence in monolayer mos2. Nano Lett 2010, 10 (4), 1271-

1275.

5. Butler, S. Z.; Hollen, S. M.; Cao, L.; Cui, Y.; Gupta, J. A.; Gutiérrez, H. R.; Heinz,

T. F.; Hong, S. S.; Huang, J.; Ismach, A. F., Progress, challenges, and opportunities

in two-dimensional materials beyond graphene. ACS nano 2013, 7 (4), 2898-2926.

6. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.;

Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A., Two-dimensional gas of massless

Dirac Fermions in graphene. Nature (London) 2005, 438 (7065), 197-200.

7. Molle, A.; Goldberger, J.; Houssa, M.; Xu, Y.; Zhang, S.-C.; Akinwande, D.,

Buckled two-dimensional xene sheets. Nat. Mater. 2017, 16 (2), 163-169.

8. Mak, K. F.; Lee, C.; Hone, J.; Shan, J.; Heinz, T. F., Atomically thin mos 2: A new

direct-gap semiconductor. Phys Rev Lett 2010, 105 (13), 136805.

9. Schönfeld, B.; Huang, J.; Moss, S., Anisotropic mean-square displacements (msd)

in single-crystals of 2h-and 3r-mos2. Acta Crystallogr Sec B 1983, 39 (4), 404-407.

10. He, Q.; Zeng, Z.; Yin, Z.; Li, H.; Wu, S.; Huang, X.; Zhang, H., Fabrication of

flexible mos2 thin‐film transistor arrays for practical gas‐sensing applications.

Small 2012, 8 (19), 2994-2999.

11. Chang, C.-Z.; Zhang, J.; Feng, X.; Shen, J.; Zhang, Z.; Guo, M.; Li, K.; Ou, Y.;

Wei, P.; Wang, L.-L., Experimental observation of the quantum anomalous hall

effect in a magnetic topological insulator. Science 2013, 340 (6129), 167-170.

12. Yu, R.; Zhang, W.; Zhang, H.-J.; Zhang, S.-C.; Dai, X.; Fang, Z., Quantized

anomalous hall effect in magnetic topological insulators. Science 2010, 329 (5987),

61-64.

13. Geim, A. K.; Grigorieva, I. V., Van der waals heterostructures. Nature 2013, 499

(7459), 419-425.

14. Lee, G.-H.; Yu, Y.-J.; Cui, X.; Petrone, N.; Lee, C.-H.; Choi, M. S.; Lee, D.-Y.;

Lee, C.; Yoo, W. J.; Watanabe, K., Flexible and transparent mos2 field-effect

transistors on hexagonal boron nitride-graphene heterostructures. ACS nano 2013,

7 (9), 7931-7936.

15. Song, I.; Park, C.; Choi, H. C., Synthesis and properties of molybdenum disulphide:

From bulk to atomic layers. RSC Adv. 2015, 5 (10), 7495-7514.

16. Wilson, J.; Yoffe, A., The transition metal dichalcogenides discussion and

interpretation of the observed optical, electrical and structural properties. Advances

in Physics 1969, 18 (73), 193-335.

17. Coutinho, S. S.; Tavares, M. S.; Barboza, C. A.; Frazão, N. F.; Moreira, E.;

Azevedo, D. L., 3r and 2h polytypes of mos2: Dft and dfpt calculations of structural,

optoelectronic, vibrational and thermodynamic properties. J. Phys. Chem. Solids

2017, 111 (Supplement C), 25-33.

18. Lee, J.-U.; Kim, K.; Han, S.; Ryu, G. H.; Lee, Z.; Cheong, H., Raman signatures of

polytypism in molybdenum disulfide. ACS Nano 2016, 10 (2), 1948-1953.

19. Bianco, E.; Butler, S.; Jiang, S.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Stability and exfoliation of germanane: A germanium graphane analogue. Acs Nano

2013, 7 (5), 4414-4421.

20. Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Sun, C.; Scudder, M. R.; Ross, R. D.;

Goldberger, J. E., Group-13 and group-15 doping of germanane. Beilstein J.

Nanotechnol. 2017, 8 (1), 1642-1648.

21. Jiang, S.; Bianco, E.; Goldberger, J. E., The structure and amorphization of

germanane. J. Mater. Chem. C. 2014, 2 (17), 3185-3188.

22. Jiang, S.; Butler, S.; Bianco, E.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Improving the stability and optical properties of germanane via one-step covalent

methyl-termination. Nat. Commun. 2014, 5, 3389.

23. Young, J. R.; Chitara, B.; Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Fan, F.;

Johnston-Halperin, E.; Goldberger, J. E., Water activated doping and transport in

multilayered germanane crystals. J. Phys. Condens. Matter 2015, 28 (3), 034001.

24. Madhushankar, B.; Kaverzin, A.; Giousis, T.; Potsi, G.; Gournis, D.; Rudolf, P.;

Blake, G.; van der Wal, C.; van Wees, B., Electronic properties of germanane field-

effect transistors. 2D Mater. 2017, 4 (2), 021009.

25. Serino, A. C.; Ko, J. S.; Yeung, M. T.; Schwartz, J. J.; Kang, C. B.; Tolbert, S. H.;

Kaner, R. B.; Dunn, B. S.; Weiss, P. S., Lithium-ion insertion properties of solution-

exfoliated germanane. ACS Nano 2017, 11 (8), 7995-8001.

26. Jiang, S.; Krymowski, K.; Asel, T.; Arguilla, M. Q.; Cultrara, N. D.; Yanchenko,

E.; Yang, X.; Brillson, L. J.; Windl, W.; Goldberger, J. E., Tailoring the electronic

structure of covalently functionalized germanane via the interplay of ligand strain

and electronegativity. Chem. Mater. 2016, 28 (21), 8071-8077.

27. Bobev, S.; Bauer, E. D.; Thompson, J. D.; Sarrao, J. L.; Miller, G. J.; Eck, B.;

Dronskowski, R., Metallic behavior of the zintl phase euge 2: Combined structural

studies, property measurements, and electronic structure calculations. J Solid State

Chem 2004, 177 (10), 3545-3552.

28. Tobash, P. H.; Bobev, S., Synthesis, structure and electronic structure of a new

polymorph of cage 2. J Solid State Chem 2007, 180 (5), 1575-1581.

29. Luo, X.; Zurek, E., Crystal structures and electronic properties of single-layer, few-

layer, and multilayer geh. J. Phys. Chem. C 2015, 120 (1), 793-800.

30. Arguilla, M. Q.; Jiang, S.; Chitara, B.; Goldberger, J. E., Synthesis and stability of

two-dimensional ge/sn graphane alloys. Chem. Mater. 2014, 26 (24), 6941-6946.

31. Munn, R., Role of the elastic constants in negative thermal expansion of axial

solids. J Phys C 1972, 5 (5), 535.

32. Taube, A.; Judek, J.; Łapińska, A.; Zdrojek, M., Temperature-dependent thermal

properties of supported mos2 monolayers. ACS Appl. Mater. Interfaces 2015, 7 (9),

5061-5065.

33. Zhang, S.; Yang, J.; Xu, R.; Wang, F.; Li, W.; Ghufran, M.; Zhang, Y.-W.; Yu, Z.;

Zhang, G.; Qin, Q.; Lu, Y., Extraordinary photoluminescence and strong

temperature/angle-dependent raman responses in few-layer phosphorene. ACS

Nano 2014, 8 (9), 9590-9596.

34. Lanzillo, N. A.; Glen Birdwell, A.; Amani, M.; Crowne, F. J.; Shah, P. B.; Najmaei,

S.; Liu, Z.; Ajayan, P. M.; Lou, J.; Dubey, M., Temperature-dependent phonon

shifts in monolayer mos2. Appl. Phys. Lett. 2013, 103 (9), 093102.

3.7 Supplemental Information

Figure 25. Le Bail refinement results of 1T-GeH using GSAS

100

Figure 26. Rietveld refinement results of 2H-GeH using GSAS

101

Figure 27. Rietveld refinement results of 6R-GeH using GSAS. Phase purity 91%.

102

EuGe2

Calc Intensity (a.u.) Intensity

10 20 30 40 50 60 70 2-Theta (o)

Figure 28. Powder XRD Topas refinement results of EuGe2

Table 3. Crystal data and refinement results for EuGe2

EuGe2 Cu Kα1 a/b = 4.10095(32) λ = 1.5406 Å c = 4.99811(44) 2θ = 10-70 α/β = 120 T = 298 K γ = 90 V = 72.796(13) Rp/Rwp = .0650/.0864

103

Table 4. Fractional atomic coordinates and isotropic displacement parameters based on the refined EuGe2 structure.

Atom Wychkoff X Y Z 2 Position Beq (Å ) Eu 6c 0 0 0.000 2.41(11) Ge (1) 6c 1/3 2/3 .4018(9) 0.418(87)

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Figure 29. Powder XRD Topas refinement results of α-CaGe2

Table 5. Crystal data and refinement results for α-CaGe2

α-CaGe2 Cu Kα1 a/b = 3.98542(5) λ = 1.5406 Å c = 4.3767(2) 2θ = 10-70 α/β = 120 T = 298 K γ = 90 V = 141.606(4) Rp/Rwp =.1001/ .1315

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Table 6. Fractional atomic coordinates and isotropic displacement parameters based on the refined α-CaGe2 structure.

Atom Wychkoff X Y Z 2 Position Beq (Å ) Ca 6c 0 0 0.3325(82) 9.7(25) Ge (1) 6c 1/3 2/3 .06084(59) 6.7(17) Ge (2) 6c 0 0 0 6.8(16)

106

Obs Calc

Diff Intensity (a.u.)Intensity

10 20 30 40 50 60 70 2-Theta (o)

Figure 30. Powder XRD Topas refinement results of β-CaGe2

Table 7. Crystal data and refinement results for β-CaGe2

β-CaGe2 Cu Kα1 a/b = 3.9986(9) λ = 1.5406 Å c = 30.62(15) 2θ = 10-70 α/β = 120 T = 298 K γ = 90 V = 421.8(30) Rp/Rwp = .0565/.0798

107

Table 8. Fractional atomic coordinates and isotropic displacement parameters based on the refined β-CaGe2 structure.

Atom Wychkoff X Y Z 2 Position Beq (Å ) Ca 6c 0 0 0.0833(26) 9.7(25) Ge (1) 6c 0 0 .1914(30) 6.7(17) Ge (2) 6c 0 0 0.3545(34) 6.8(16)

5.6580 5.6570 5.6560 5.6550

5.6540 lattice constant lattice

- 5.6530 a 5.6520 5.6510 75 125 175 225 275 325 Temperature (k)

Figure 31. Germanium lattice constants as a function of temperature

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Chapter 4: Electronic characterization of layered systems

4.1 Introduction

Since the discovery of graphene in 2004, a boom in research has been underway in a quest to identify and isolate single and few layered 2-dimensional materials from layered crystals. Graphene, a single honeycombed layer of sp2 hybridized carbon, was first isolated though mechanical exfoliation of graphite crystals, physically pulling apart layers with

Scotch tape1. By thinning to a single layer, it allows for the study of a new regime of physics in which allows for the exploration of two-dimensional density of states as a result of the Dirac-band crossing at the Fermi-level. This allows suspended single layer devices to be fabricated which show 2D electron gas-like properties with an ultra-high carrier mobility of >200,000 cm2 V-1 s -1.2 A room-temperature quantized Hall Effect was measured in this system in 20053, continuing progress in these fundamentally thin systems as sources of interesting physical phenomena. Furthermore, due the sensitivity of graphene to the surrounding environment, the quantum anomalous Hall Effect has been predicted by the interaction graphene heterostructures with ferromagnetic materials4.

While graphene shows remarkable properties, there are some inherent limitations as a materials system. One of the most restrictive properties of graphene is the lack of inherent band-gap, which limits ability to use it in traditional transistor devices. Band-gap engineering in graphene is accomplished through altering the chemical environment

109 around the graphene sheets, with chemical functionalization5, doping6, 7, and with coupling substrates8 and other layered structures9, 10. While this accomplishes the task of opening a band-gap, it typically drastically reduces the mobility of the system.

The interest in the graphene system has drastically increased the visibility of other layered materials systems which were already known which have the ability to be thinned down.

One family of systems is the transition metal dichalcogenides (TMDs), which have been used in the past a lubricants and which exist as natural compounds. MoS2 is a known crystal systems, which contains layers of Mo-S oriented in trigonal prismatic layers, which are separated by a van der Waals gap. The presence of this gap, much like in graphene, allows MoS2 to be mechanically exfoliated down to single layers. This thinning causes a band-gap shift from an indirect gap of 1.29 to a desirable direct gap of 1.90 eV, a trend which is followed by other TMDs.

A third system heavily studied is the single layer structure of black phosphorus, a single armchair like layer called phosphorene. Again, these structures can be exfoliated down from a bulk crystal, allowing for the exploration of physics not observed in the bulk material. Phosphorene also shows a widening of the band-gap as thickness of the crystal is decreased, increasing the gap from 0.3 to 1.5 eV11, 12. Thinning also increases the mobility of measured devices, making it an intriguing systems for optoelectronics.

The previous systems highlight the potential in electronic application for layered semiconducting and metallic materials. It is important that when developing new crystal systems that we are able to characterize their electronic properties to fully understand the potential application and inherent properties. This chapter focuses on the device fabrication and bulk electronic measurements of three of these systems, GeH, NaSn2As2,

110 and EuSn2As2. Finally, it discusses the progress in reducing thinning materials to few layers for device fabrication.

4.2 Highly resistive measurements of germanane

Many materials which have been grown, though they may have high theoretical mobilities, for one reason or another turn out to be highly restive materials. When dealing with these materials, it is important to maximize current which can be measured without damaging the materials with high applied biases. There are multiple scenarios in which you can fabricate devices in order to increase the signal which is observed and to achieve a more real observation of a result. Both scenarios are present in the germanane system, due to the high resistivity and the presence of dopants which change the identity of the material.

Firstly, the dimensions of contacts are very important in determining the amount of current which can be measured. Traditional methods looks to minimize the contact size in order to decrease the role of contact resistance plays in a measurement. When dealing with highly resistive materials it helps to increase the size of the contact to increase the lateral dimension of the contact to increase the volume in which the flux of electrons occur during a transport measurement. The second way of maximizing observed current is to match the work function of the contact metal to that of the Fermi-level of material. This is especially important in semi-conductors.

Fabricating these sheet resistance devices is accomplished with a shadow mask technique.

First the crystallite of interest must be separated from larger crystallites. Razor blades are used to flake smaller crystallites from a larger crystal under a microscope. Different

111 crystalline domains can be observed and by placing the corner of a razor and applying force will separate the layers. The crystallite of interest is first placed on a small glass slide which has been diced to ~1 cm x 1 cm. Small pieces of wire, which determine the channel length of the contacts, are draped across the flake and attached to the glass slide with kapton tape. It is important that the wire is attached to be stretched tight and holding down the flake, since metal evaporation with take place in high vacuum and the wire is the only thing attaching the flake to the glass. Multiple glass slide pieces are attached to a larger glass slide for metal deposition. The slides are them loaded into a metal deposition instrument, loading directly into the chamber on a rotation arm to ensure that angle of deposition is completely normal the surface of the flakes. This minimized the effects of shadowing from the wire and a uniform contact.

The deposition of metal onto the surface also needs to be optimized to ensure the maximum response to an applied voltage. This is accomplished by matching the work function of the metal contact to the Fermi energy of the semiconductor. In the case of germanane, multiple depositions of test metals was used to determine which metal to surface interface gave the most ohmic of contact. In the doped germanane systems, depositions of Ag:Au, Au, Ti:Au,

Cr:Au, and Pd:Au13, 14 were carried out to determine which metal gives the most linear contacts. Any deviation from linearity will indicate a Schottky barrier and shows the devices are not functioning properly.

Due to the large contact area which must be used, only two probe measurements can be carried out on these highly resistive materials. To determine the resistance of a measurement, when the ohmic contacts are found, a linear trend line can be fit to the data.

The inverse of trendline is a direct calculation of the resistance, which is converted into

112

sheet resistance by measuring the respective dimensions (Figure 32) using optical

microscopy (Figure 32) and using equation 1:

퐿∗훺 휌 = 푊∗푇

The resistance is multiplied by the length of the channel of transport and divided by the

Figure 32. a) Schematic dimensions measured in b)bulk crystal devices in determining sheet resistance

area in which the electron flux can travel through. Due to the highly resistive nature of

these materials, multiple devices are measured and averaged.

4.3 Bulk crystal measurements of NaSn2As2

15 NaSn2As2, an electron deficient layered Zintl phase has been developed . These crystals

are achieve through annealing of elemental precursors in silica crucibles in evacuated

quartz tubes. This NaSn2As2 structure contains a van der Waals gap between neighboring

tin layers which allows it to be mechanically and chemically exfoliated down to few layer

structures. The fact that these materials are exfoliatable and metallic makes them an

intriguing system to explore, first using bulk electronic measurements.

113

Contact fabrication was done via indium stamping onto the crystals of NaSn2As2. This is accomplished by cutting small pieces of indium, and manipulating with sharpened tooth picks. These indium pieces have a tendency to stick to metallic surfaces and due to the soft nature of indium it allows for the pieces to be physically pressed onto the surface of a crystal. A small portion of indium in pressed with the toothpicks until an amount sticks to the tip of the wood. This allows for easier manipulation with larger pieces of indium which will be used to make contacts of the surface. Another small piece of indium is them moved to the surface of the crystal and second sharpened toothpick, without indium stuck to the tip, and the edges are teased down to surface of the crystal while keeping the contact as small and symmetric as possible. Once the indium is pressed onto the surface, a copper wire (25 µm) is placed on to the indium spot and the indium is folded around the copper wire. After contacts are made to the surface, wires are attached to contact pads on a measurement puck allowing for a 2-probe measurement to be conducted.

Transport measurements were collected at temperatures from 300 K down to 2 K. Figure

33 shows measurements of a bulk crystal of NaSn2As2. At high temperatures, a two-probe measurement shows very low resistivity of 2.9 µΩ·m, which decreases with a decreasing temperature. The decrease in resistivity as a function of temperature is consistent with phonon-mediated carrier scattering observed in materials which exhibit metallic behavior.

114

Figure 33. Temperature dependent transport measurements of bulk and exfoliated NaSn2As2 crystals down to 2.1 K

4.4 Bulk crystal measurements of EuSn2As2

16 EuSn2As2 is an exfoliatable magnetic Zintl phase . Bulk crystals of EuSn2As2 was

prepared via sealing elemental Eu, Sn, and As in evacuated quartz and annealed. Thin

crystals are separated from larger bulk crystals via fracturing between two pieces of weigh

paper. Once between weigh papers, the back of tweezers can be used to separate the

crystals, typically crystals which fracture along planes different crystal domains. Crystals

115 can be thinned further by placing the pointed edge of a razor between different growth domains and lifting the domains apart from one another.

Contacts are fabricated using silver epoxy. The epoxy is applied using a sharpened toothpick. Once a tooth pick is sharpened to a very fine point, the toothpick is placed into the epoxy and lifted out in one concerted motion. The goal is to have an epoxy tail which become a point at the end which is used to apply a small amount of epoxy to the surface of the crystal. Further precision to the application of the epoxy can be done by using an epoxy which has been partially cured, either by pulling a brief vacuum over it or by gently heating it. This increases the viscosity of the epoxy allowing for a smaller tail to be formed. After the epoxy is placed in four equally sized and spaced points, 25 µm Cu wire is placed in each droplet and subsequently covered with a small portion of silver epoxy and cured at

116

Figure 34. Temperature-dependent resistivity of EuSn2As2 showing its metallic behavior along the temperature range with a cusp near the magnetic ordering temperature. Top left: Characteristic micrograph of a EuSn2As2 crystal in a four-probe geometry. Bottom right: high temperature resistivity which increases with increasing temperatures

117

150o C for 5 minutes. Wires are attached to PPMS puck using indium stamping. Figure

34 shows temperature dependent in-plane 4-probe resistivity measurement of the bulk

EuSn2As2 crystal confirms metallic behavior with a phonon-meditated increase of resistivity at higher temperatures. At temperatures approaching the paramagnetic transition temperature of EuSn2As2 (30-24 K), an increase of ~1.2% in resistivity is observed do to the spin-scattering of conduction electrons due to the Eu2+ f-orbital electrons.

4.5 Conclusions

Here we have successfully measured the electronic properties a variety of electronic materials. Measurements of highly resistive materials is achieved by depositing contacts via metal deposition using a shadow masking technique. This allows for a larger response to the applied bias and can be measured above the noise floor of instrumentation. To allow for direct comparison of materials, sheet resistance is calculated. Bulk crystal measurements of NaSn2As2 and EuSn2As2 was achieved though the used of point contacts fabricated with indium stamping and silver epoxy, respectively. These measurements confirm predicted metallic behavior of both systems. Furthermore, these measurements provide valuable insight into the role which magnetic ordering in the EuSn2As2 system plays in the transport phenomena along with a secondary confirmation of measured magnetic data.

118

4.6 References

1. Novoselov, K. S.; Geim, A. K.; Morozov, S.; Jiang, D.; Katsnelson, M.; Grigorieva, I.;

Dubonos, S.; Firsov; AA, Two-dimensional gas of massless Dirac Fermions in

graphene. nature 2005, 438 (7065), 197-200.

2. Bolotin, K. I.; Sikes, K.; Jiang, Z.; Klima, M.; Fudenberg, G.; Hone, J.; Kim, P.;

Stormer, H., Ultrahigh electron mobility in suspended graphene. Solid State

Communications 2008, 146 (9), 351-355.

3. Zhang, Y.; Tan, Y.-W.; Stormer, H. L.; Kim, P., Experimental observation of the

quantum hall effect and berry's phase in graphene. nature 2005, 438 (7065), 201-

204.

4. Tse, W.-K.; Qiao, Z.; Yao, Y.; MacDonald, A.; Niu, Q., Quantum anomalous hall

effect in single-layer and bilayer graphene. Physical Review B 2011, 83 (15),

155447.

5. Haberer, D.; Vyalikh, D.; Taioli, S.; Dora, B.; Farjam, M.; Fink, J.; Marchenko, D.;

Pichler, T.; Ziegler, K.; Simonucci, S., Tunable band gap in hydrogenated quasi-

free-standing graphene. Nano letters 2010, 10 (9), 3360-3366.

6. Chang, C.-K.; Kataria, S.; Kuo, C.-C.; Ganguly, A.; Wang, B.-Y.; Hwang, J.-Y.;

Huang, K.-J.; Yang, W.-H.; Wang, S.-B.; Chuang, C.-H., Band gap engineering of

chemical vapor deposited graphene by in situ bn doping. Acs Nano 2013, 7 (2),

1333-1341.

119

7. Park, J.; Jo, S. B.; Yu, Y. J.; Kim, Y.; Yang, J. W.; Lee, W. H.; Kim, H. H.; Hong, B.

H.; Kim, P.; Cho, K., Single‐gate bandgap opening of bilayer graphene by dual

molecular doping. Advanced materials 2012, 24 (3), 407-411.

8. Zhou, S. Y.; Gweon, G.-H.; Fedorov, A.; First, P.; De Heer, W.; Lee, D.-H.; Guinea,

F.; Neto, A. C.; Lanzara, A., Substrate-induced bandgap opening in epitaxial

graphene. Nature materials 2007, 6 (10), 770-775.

9. Ramasubramaniam, A.; Naveh, D.; Towe, E., Tunable band gaps in bilayer graphene−

bn heterostructures. Nano letters 2011, 11 (3), 1070-1075.

10. Schwierz, F., Graphene transistors. Nature nanotechnology 2010, 5 (7), 487-496.

11. Li, L.; Yu, Y.; Ye, G. J.; Ge, Q.; Ou, X.; Wu, H.; Feng, D.; Chen, X. H.; Zhang,

Y., Black phosphorus field-effect transistors. Nature nanotechnology 2014, 9 (5),

372-377.

12. Tran, V.; Soklaski, R.; Liang, Y.; Yang, L., Layer-controlled band gap and

anisotropic excitons in few-layer black phosphorus. Physical Review B 2014, 89

(23), 235319.

13. Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Sun, C.; Scudder, M. R.; Ross, R. D.;

Goldberger, J. E., Group-13 and group-15 doping of germanane. Beilstein Journal

of Nanotechnology 2017, 8, 1642-1648.

14. Young, J. R.; Chitara, B.; Cultrara, N. D.; Arguilla, M. Q.; Jiang, S.; Fan, F.;

Johnston-Halperin, E.; Goldberger, J. E., Water activated doping and transport in

multilayered germanane crystals. Journal of Physics: Condensed Matter 2015, 28

(3), 034001.

120

15. Arguilla, M. Q.; Katoch, J.; Krymowski, K.; Cultrara, N. D.; Xu, J.; Xi, X.;

Hanks, A.; Jiang, S.; Ross, R. D.; Koch, R. J., Nasn2as2: An exfoliatable layered

van der waals zintl phase. ACS nano 2016, 10 (10), 9500-9508.

16. Arguilla, M.; Cultrara, N.; Baum, Z.; Jiang, S.; Ross, R.; Goldberger, J., Eusn 2 as

2: An exfoliatable magnetic layered zintl–klemm phase. Inorganic Chemistry

Frontiers 2017, 4 (2), 378-386.

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Chapter 5: Conclusions and Outlook

The discovery of graphene has fundamentally altered the landscape and trajectory of materials for scientists who are interested in discovery at the nanoscale regime. It started a boom on research on layered materials which show different properties at few layers in comparison to what is observed in the bulk crystals. Herein, we show the synthesis and characterization of layered materials, focusing on: 1) improving the electronic properties of 6R germanane though the introduction of dopants 2) synthesis and characterization of structural polymorphs of the layered germanane structure and 3) device fabrication and measurements of layered electronic materials.

In chapter 2, we demonstrated that gallium and arsenic can be incorporated into the precursor CaGe2 Zintl phase and are retained in the 2D germanium framework after the topotactic deintercalation process. These dopants do not significantly change the structure of germanane. These doped materials are stable in ambient atmosphere conditions for at least 24 hours but start to oxidize between 1-4 days. The introduction of Ga and As to the lattice decreases the resistance in ambient conditions with large amounts of hysteresis, suggesting that the presence of water can activate these dopants. As was previously observed with P:GeH,[29] As:GeH is highly resistive under vacuum, indicating the presence of water is required to activate group 15 dopants. In contrast, Ga:GeH exhibited decreased sheet resistances in vacuum by over 4-orders of magnitude, which was proportional to the amount of gallium and exhibited minimal hysteretic behaviour. This indicates that the dopant activated state in Ga:GeH is stable under vacuum, enabling robust

122 electronic properties through encapsulation. Overall, this work provides a pathway to dope germanane and enable future explorations of electronic devices.

In chapter 3, we 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. 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.

In chapter 4, we demonstrated the ability to create devices for electronic characterization of bulk materials. We show how to maximize the response to applied bias in highly resistive materials by selecting the contact material to minimize contact resistance along with creating large contacts to maximize the area which allows electron flux through. The characterization temperature dependent transport properties of similar 1:2:2 phases,

NaSn2As2 and EuSn2As2, both demonstrate phonon mediated electronic transport which is expected for metallic crystals. The presence of the magnetic europium in the 1:2:2 systems can be directly observed at the low temperature measurements, where and increase in resistivity in the 24-30 K range demonstrates spin-scattering caused by the paramagnetic

123 transition at low temperatures. Finally we show the progress to making few layer devices of exfoliatable magnetic metals.

124

Bibliography

Chapter 1

1. Lalović, B.; Kiss, Z.; Weakliem, H., A hybrid amorphous silicon photovoltaic and

thermal solar collector. Solar Cells 1986, 19 (2), 131-138.

2. Auston, D. H.; Lavallard, P.; Sol, N.; Kaplan, D., An amorphous silicon

photodetector for picosecond pulses. Appl. Phys. Lett. 1980, 36 (1), 66-68.

3. Joshi, G.; Lee, H.; Lan, Y.; Wang, X.; Zhu, G.; Wang, D.; Gould, R. W.; Cuff, D.

C.; Tang, M. Y.; Dresselhaus, M. S.; Chen, G.; Ren, Z., Enhanced Thermoelectric

Figure-of-Merit in Nanostructured p-type Silicon Germanium Bulk Alloys. Nano

Lett. 2008, 8 (12), 4670-4674.

4. Hochbaum, A. I.; Chen, R.; Delgado, R. D.; Liang, W.; Garnett, E. C.; Najarian,

M.; Majumdar, A.; Yang, P., Enhanced thermoelectric performance of rough

silicon nanowires. Nature 2008, 451 (7175), 163-167.

5. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Zhang, Y.; Dubonos, S.

V.; Grigorieva, I. V.; Firsov, A. A., Electric field effect in atomically thin carbon

films. Science 2004, 306 (5696), 666-669.

6. Butler, S. Z.; Hollen, S. M.; Cao, L.; Cui, Y.; Gupta, J. A.; Gutierrez, H. R.; Heinz,

T. F.; Hong, S. S.; Huang, J.; Ismach, A. F.; Johnston-Halperin, E.; Kuno, M.;

Plashnitsa, V. V.; Robinson, R. D.; Ruoff, R. S.; Salahuddin, S.; Shan, J.; Shi, L.;

Spencer, M. G.; Terrones, M.; Windl, W.; Goldberger, J. E., Progress, challenges,

125

and opportunities in two-dimensional materials beyond graphene. ACS Nano 2013,

7 (4), 2898-2926.

7. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.;

Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A., Two-dimensional gas of massless

Dirac Fermions in graphene. Nature 2005, 438 (7065), 197-200.

8. Balandin, A. A.; Ghosh, S.; Bao, W.; Calizo, I.; Teweldebrhan, D.; Miao, F.; Lau,

C. N., Superior Thermal Conductivity of Single-Layer Graphene. Nano Lett. 2008,

8 (3), 902-907.

9. Lee, C.; Wei, X.; Kysar, J. W.; Hone, J., Measurement of the Elastic Properties and

Intrinsic Strength of Monolayer Graphene. Science 2008, 321 (5887), 385-388.

10. Elias, D. C.; Nair, R. R.; Mohiuddin, T. M. G.; Morozov, S. V.; Blake, P.; Halsall,

M. P.; Ferrari, A. C.; Boukhvalov, D. W.; Katsnelson, M. I.; Geim, A. K.;

Novoselov, K. S., Control of graphene's properties by reversible hydrogenation:

evidence for graphane. Science 2009, 323 (5914), 610-613.

11. Yamanaka, S.; Matsu-ura, H.; Ishikawa, M., New deintercalation reaction of

calcium from calcium disilicide. Synthesis of layered polysilane. Mater. Res. Bull.

1996, 31, 307-316.

12. Bianco, E.; Butler, S.; Jiang, S.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Stability and Exfoliation of Germanane: A Germanium Graphane Analogue. ACS

Nano 2013, 7 (5), 4414-4421.

13. Jiang, S.; Butler, S.; Bianco, E.; Restrepo, O. D.; Windl, W.; Goldberger, J. E.,

Improving the stability and optical properties of germanane via one-step covalent

methyl-termination. Nat. Commun. 2014, 5, 3389.

126

14. Xu, Y.; Yan, B.; Zhang, H.-J.; Wang, J.; Xu, G.; Tang, P.; Duan, W.; Zhang, S.-C.,

Large-gap quantum spin hall insulators in tin films. Physical Review Letters 2013,

111, 136804.

15. Knapp, D. Chemistry and Electronics of the Ge(111) Surface. Dissertation (Ph.D.),

California Institute of Technology, 2011.

16. Bansal, A.; Li, X.; Lauermann, I.; Lewis, N. S.; Yi, S. I.; Weinberg, W., Alkylation

of Si surfaces using a two-step halogenation/Grignard route. J. Am. Chem. Soc.

1996, 118 (30), 7225-7226.

17. Han, S. M.; Ashurst, W. R.; Carraro, C.; Maboudian, R., Formation of alkanethiol

monolayer on Ge (111). J. Am. Chem. Soc. 2001, 123 (10), 2422-2425.

18. Linford, M. R.; Fenter, P.; Eisenberger, P. M.; Chidsey, C. E., Alkyl monolayers

on silicon prepared from 1-alkenes and hydrogen-terminated silicon. J. Am. Chem.

Soc. 1995, 117 (11), 3145-3155.

19. Buriak, J. M., Organometallic chemistry on silicon and germanium surfaces. Chem.

Rev. 2002, 102 (5), 1271-1308.

20. Wöhler, F., Ueber Verbindungen des Siliciums mit Sauerstoff und Wasserstoff.

Liebigs. Ann. 1863, 127 (3), 257-274.

21. Kautsky, H., Über einige ungesättigte Siliciumverbindungen. Z. anorg. Chem.

1921, 117 (1), 209-242.

22. Kautsky, H.; Herzberg, G., Über das Siloxen und seine Derivate. Z. anorg. allg.

Chem. 1924, 139 (1), 135-160.

23. Weiss, A.; Beil, G.; Meyer, H., The Topochemical Reaction of CaSi2 to a Two-

Dimensional Subsiliceous Acid Si6H3(OH)3. Z. Naturforsch. 1980, 35b, 25-30.

127

24. Dettlaff-Weglikowska, U.; Hönle, W.; Molassioti-Dohms, A.; Finkbeiner, S.;

Weber, J., Structure and optical properties of the planar silicon compounds

polysilane and Wöhler siloxene. Phys. Rev. B 1997, 56 (20), 13132-13140.

25. Okamoto, H.; Kumai, Y.; Sugiyama, Y.; Mitsuoka, T.; Nakanishi, K.; Ohta, T.;

Nozaki, H.; Yamaguchi, S.; Shirai, S.; Nakano, H., Silicon nanosheets and their

self-assembled regular stacking structure. Journal of the American Chemical

Society 2010, 132 (8), 2710-2718.

26. Nakano, H.; Nakano, M.; Nakanishi, K.; Tanaka, D.; Sugiyama, Y.; Ikuno, T.;

Okamoto, H.; Ohta, T., Preparation of alkyl-modified silicon nanosheets by

hydrosilylation of layered polysilane (Si6H6). Journal of the American Chemical

Society 2012, 134 (12), 5452-5455.

27. Sugiyama, Y.; Okamoto, H.; Mitsuoka, T.; Morikawa, T.; Nakanishi, K.; Ohta, T.;

Nakano, H., Synthesis and Optical Properties of Monolayer Organosilicon

Nanosheets. J. Am. Chem. Soc. 2010, 132 (17), 5946-5947.

28. Vogg, G.; Brandt, M. S.; Stutzmann, M., Polygermyne—A Prototype System for

Layered Germanium Polymers. Adv. Mater. 2000, 12 (17), 1278-1281.

29. Jiang, S.; Bianco, E.; Goldberger, J. E., The structure and amorphization of

germanane. J. Mater. Chem. C 2014, 2 (17), 3185-3188.

30. Arguilla, M. Q., Jiang, S., Chitara, B., Goldberger, J. E., Synthesis and Stability of

Two-dimensional Ge/Sn Graphane Alloys. Submitted 2014.

31. Frindt, R.; Yoffe, A., Physical properties of layer structures: optical properties and

photoconductivity of thin crystals of molybdenum disulphide. Proc. R. Soc.

London, Ser. A. 1963, 273 (1352), 69-83.

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