<<

AN ABSTRACT OF THE DISSERTATION OF

Sean W. Muir for the degree of Doctor of Philosophy in presented on April 17, 2012.

Title: Structure-property Relationships of Layered

Abstract approved:______M. A. Subramanian

Investigating the structure-property relationships of solid state materials can help improve many of the materials we use each day in life. It can also to the discovery of materials with interesting and unforeseen properties. In this work the structure property relationships of newly discovered layered phases are presented and discussed. There has generally been worldwide interest in layered oxypnictide materials following the discovery of up to 55 K for such as LnFeAsO1-xFx (where Ln = Lanthanoid). This work presents efforts to understand the structure and physical property changes which occur to LnFeAsO materials when Fe is replaced with Rh or Ir and when As is replaced with Sb. As part of this work the solid solution between LaFeAsO and LaRhAsO was examined and superconductivity is observed for low Rh content with a maximum critical temperature of 16 K. LnRhAsO and LnIrAsO compositions are found to be metallic; however Ce based compositions display a resistivity temperature dependence which is typical of Kondo lattice materials. At low temperatures a sudden drop in resistivity occurs for both CeRhAsO and CeIrAsO compositions and this drop coincides with an antiferromagnetic transition. The Kondo scattering temperatures and magnetic transition temperatures observed for these materials can be rationalized by considering the expected difference in N(EF)J parameters between them, where N(EF) is the density of states at the Fermi level and J represents the exchange interaction between the Ce 4f1 and the conduction electrons. In addition to studying these 4d and 5d substituted systems the LaFeSbO compositional system was investigated. While LaFeSbO has not been successfully synthesized the transition free layered oxypnictide composition La2SbO2 was discovered and its structural and physical properties have been examined along with the properties of La2BiO2.

Density functional theory was used to calculate the heats of formation for competing phases within the LaFeSbO system, in order to better understand the stability of

LaFeSbO and why it has not yet been observed. The materials La2SbO2 and La2BiO2 were investigated for the presence of vacancies using powder neutron diffraction. Structure refinement reveals that there is significant disorder within the a-b plane for Sb compositions.

©Copyright by Sean W. Muir

April 17, 2012

All Rights Reserved

Structure-property Relationships of Layered Oxypnictides

by

Sean W. Muir

A DISSERTATION

submitted to

Oregon State University

In partial fulfillment of the requirements for the degree of

Doctor of Philosophy

Presented April 17, 2012

Commencement June 2012 Doctor of Philosophy dissertation of Sean W. Muir presented on April 17, 2012.

APPROVED:

Major Professor, representing Chemistry

Chair of the Department of Chemistry

Dean of the Graduate School

I understand that my dissertation will become part of the permanent collection of Oregon State University libraries. My signature below authorizes release of my dissertation to any reader upon request.

Sean W. Muir, Author

ACKNOWLEDGMENTS

There is no way I can possibly list everyone that has helped me along the way with research advice, words of encouragement, and general kindness. Thus I apologize to all those I will miss in my attempt.

First off I need to thank my thesis advisor Mas Subramanian for giving me all the freedom, encouragement, and resources needed grow into the best possible scientist and chemist I can be. By providing such a great laboratory to work in I was allowed to truly work at the limits of my own ability.

Thanks to my three awesome office mates who helped to transform it from just an office into a tropical jungle and who tolerated the curious system of organization used on my desk; Geneva Laurita-Plankis, Rosa Grajczyk, and Whitney

Schmidt. Thanks to James Eilertsen, Romain Berthelot, and Alvin Gatimu for many a great discussion. Double thanks to these guys and thanks to Peng Jiang, Theeranun

Siritanon, and Krishnendu Biswas for help in the lab and for never telling me to turn down my music while working in the glove box right outside their offices. Thanks to

Jun Li for helpful conversations about x-ray diffraction. Thanks to Andy Smith for showing me around the lab when I first arrived at OSU.

Special thanks to Arthur Sleight, Michael Lerner, Brady Gibbons, and Guenter

Schneider for serving as my thesis committee and critically reviewing this work, with a double thanks to Guenter Schneider and to Jason Vielma as well for insightful discussions about band structure computations and for happily collaborating on the investigation of LaFeSbO.

I have also been blessed to encounter a number of amazing science educators along my path going all the way back way to my undergraduate days. Richard

Nafshun and Christine Pastorek have been a huge influence on me since arriving at

OSU. Their advice both as professional mentors and just genuinely nice people has been quite helpful. Extra thanks to Rick for guidance and encouragement during my time teaching one of the general chemistry courses and to Chris for letting me be a teaching assistant as part of the integrated chemistry series, both experiences helped immensely to solidify my love for teaching. Going back to my days before OSU at

Evergreen State College, where my passion for science was first realized, I need to thank Rebecca Sunderman, Dharshi Bopegedera, Judy Cushing, and Robert Knapp as each of them took part in steering me to the course I am now on.

A huge shout out needs to go out to Chris Knutson and Bill Cowell for many exciting and enlightening discussions about all topics under the sun. Also to all members past, present, and future of the SSA, a forum that has provided me with as many personal friends as it has connections to stellar scientists and academics.

Last and most importantly, I must thank my family. Thanks to my pops, for always being there. Each day as I grow older I understand more fully what it took for you to be the amazing person you are and to have made the sacrifices you did.

Thanks to my mom, for nurturing my body and mind as a child and for showing me the depth of human nature. Thanks for my brothers and sister for providing all the stories and memories that only growing up together can give. Finally I get to thank the most wonderful people I have in my life, my beautiful wife and children. It is for you that I try to succeed and it is you all that want to continue journeying through life with. Thanks for filling my dining room table office with enough music, art, tasty food, and love to see me through this scientific voyage. CONTRIBUTION OF AUTHORS

Chapter 2: Dr. Mas Subramanian and Dr. Arthur Sleight contributed in the preparation and review of the manuscript on which this chapter is primarily based,

Muir, S., Sleight, A.W., Subramanian, M.A., 2010. Mat. Res. Bull. 45, 392–395.

Chapter 3: Dr. Mas Subramanian and Dr. Arthur Sleight contributed in the preparation and review of the manuscript on which this chapter is primarily based,

Muir, S., Sleight, A.W., Subramanian, M.A., 2011. J. Solid State Chem. 184, 1972–

1976.

Chapter 4: Dr. Mas Subramanian and Dr. Arthur Sleight contributed in the preparation of the manuscript on which this chapter is primarily based. Dr. Guenter

Schneider and Jason Vielma performed the density functional theory calculations which provided the heats of formation and the phonon dispersion plot for LaFeSbO and related phases. They each also contributed to the writing and review of the resulting manuscript, Muir, S., Vielma, J., Schneider, G., Sleight, A.W., Subramanian,

M.A., 2012. J. Solid State Chem. 185, 156–159. Dr. Maxim Avdeev collected the neutron powder diffraction data for La2SbO1.5 and La2BiO1.5 samples on the Australian

Nuclear Science and Technology Organization - Bragg Institute’s Echidna diffractometer. Appendix A: Mas Subramanian and Omar Rachdi contributed to the preparation and review of the manuscript on which this chapter is based, Muir, S., Rachdi, O.D.,

Subramanian, M.A., 2012. Mat. Res. Bull. 47, 798–800. Omar Rachdi also contributed to the synthesis of these compositions as an undergraduate researcher under the guidance of S. Muir

Appendix B: Mas Subramanian, Chris Knutson, and Richard Nafshun contributed to the review of the manuscript on which this chapter is based. Additionally Chris

Knutson helped with the classroom outreach component of this project.

TABLE OF CONTENTS

Page

1 General Overview and Introduction to Layered Oxypnictides...... 1

1.1 Introduction ...... 1

1.2 Introduction to LnMAsO materials (where Ln = Lanthanoid, M = late transition

metal) ...... 4

1.3 References ...... 27

2 Superconductivity within the LaFe1-xRhxAsO system ...... 32

2.1 Abstract ...... 32

2.2 Introduction ...... 33

2.3 Results and Discussion ...... 36

2.3.1 Lattice parameters and unit cell volumes of LaFe1-xRhxAsO compositions 36

2.3.2 Electronic properties of LaFe1-xRhxAsO compositions ...... 39

2.3.3 Structural refinement of LaFe0.9Rh0.1AsO ...... 49

2.4 Conclusion ...... 54

2.5 Materials and Methods ...... 55

2.6 References ...... 57 TABLE OF CONTENTS (Continued)

Page

3 Synthesis and Electronic Properties of LnRhAsO and LnIrAsO (Ln = La, Ce, Nd)

Compositions ...... 60

3.1 Abstract ...... 60

3.2 Introduction ...... 61

3.3 Results and Discussion ...... 63

3.3.1 Crystal structure of LnRhAsO and LnIrAsO materials...... 63

3.3.2 Physical properties of LnRhAsO and LnIrAsO materials ...... 73

3.3.3 Electronic structure of LaRhAsO and LaIrAsO ...... 88

3.4 Conclusion ...... 94

3.5 Materials and Methods ...... 95

3.5.1 Experimental ...... 95

3.5.2 Computational ...... 96

3.6 References ...... 98

4 The search for LaFeSbO and the synthesis and properties of La2PnO2 compositions

(where Pn = Sb and Bi) ...... 101

4.1 Abstract ...... 101 TABLE OF CONTENTS (Continued)

Page 4.2 Introduction ...... 102

4.3 Results and discussion ...... 106

4.3.1 The search for LaFeSbO ...... 106

4.3.1.1 Investigation of LaFeSb1-xAsxO and LaFe1-yZnySbO solid solutions .... 106

4.3.1.2 Identification of La2SbO2 and a case of mistaken identity ...... 108

4.3.1.3 Calculation of heats of formation within the La-Fe-Sb-O system ...... 110

4.3.2 Further investigation of La2PnO2 compositions (where Pn = Sb and Bi)... 114

4.3.2.1 Synthesis and initial x-ray diffraction studies ...... 114

4.3.2.2 Neutron diffraction studies of the nominal compositions La2SbO1.5 and La2BiO1.5 ...... 117

4.3.2.3 Electronic properties of the nominal compositions La2SbO1.5 and La2BiO1.5 ...... 122

4.4 Conclusion ...... 127

4.5 Materials and Methods ...... 129

4.5.1 Experimental ...... 129

4.5.2 Calculations ...... 131

4.6 References ...... 133

General Conclusions and Future Work ...... 136

Bibliography ...... 140 TABLE OF CONTENTS (Continued)

Page

Appendices ...... 147

Appendix A Rapid Microwave Synthesis of the Iron Arsenides NdFeAsO and

NdFe0.9Co0.1AsO ...... 148

A.1 Abstract ...... 148

A.2 Introduction ...... 149

A.3 Materials and Methods ...... 151

A.4 Results and Discussion ...... 152

A.5 Conclusion ...... 157

A.6 References...... 158

Appendix B Celebrating superconductivity: A case study for developing high-

impact science education programs ...... 159

B.1 Abstract ...... 159

B.2 Introduction ...... 160

B.3 Results and discussion ...... 163

B.4 Conclusions ...... 169

B.5 References ...... 170

LIST OF FIGURES

Figure Page

1.1 The evolution from simple LnPn structure to the LnMPnO structure (upper). Atomic positions, formation of fluorite and anti-fluorite layers, and the M-M bonding which occurs within the MPn layers for LnMPnO materials (lower)...... 6

1.2 Schematic depiction of electronic energy bands for a metal, semimetal, and insulator. The Fermi energy is marked as EF...... 23

1.3 Electronic band structure and Fermi surface for LaFeAsO, calculated using crystallographic and calculation parameters as reported in Ref. [73]...... 24

2.1 The structure of LaFeAsO and related 1111-type compositions 3+ 2- 1+ 1- where in (a) the structure is considered as [Ln O ] [FeAs] stacking along the c-axis and (b) the structure considered as 3+ 2- 3- 2- 2+ [La O As ] stacking with metallic Fe sheets...... 36

2.2 Powder x-ray diffraction (PXRD) patterns for LaFe1-xRhxAsO compositions...... 37

2.3 Lattice parameters and unit cell volumes for the LaFe1-xRhxAsO solid solution; values shown for LaFeAsO are from Ref. [1]...... 38

2.4 Normalized resistivity for the LaFe1-xRhxAsO solid solution. Compositions stated are based on the nominal of the reaction...... 41

2.5 The normalized resistivity of LaFe1-xRhxAsO samples (where x=0-0.2) plotted as function of T2 and inset showing the normalized resistivity plotted as a function of T...... 42

2.6 An electronic phase diagram for the LaFe1-xRhxAsO solid solution as determined from changes in the resistivity data. Compositions are based on nominal reaction stoichiometry...... 44

2.7 DC magnetization volume susceptibility (field cooled and zero field cooled conditions) for LaFe0.9Rh0.1AsO...... 45 LIST OF FIGURES (Continued)

Figure Page

2.8 Resistivity normalized to 30 K for LaFe0.9Rh0.1AsO plotted as a function of magnetic field and with an inset detailing the change in temperature for 10%, 50%, and 90% resistance after application of the magnetic field...... 46

2.9 Hall coefficient for LaFe0.9Rh0.1AsO as well as LaFeAsO and LaFeAsO0.89F0.11 as determined from Refs. [21] and [37]...... 48

2.10 Rietveld refinement results LaFe0.9Rh0.1AsO detailing the average As-Fe-As angle found and its change relative to LaFeAsO, as reported in Ref. [1]. The tick marks from top to bottom represent the (hkl) reflections for LaFe0.9Rh0.1AsO, La2O3, and FeAs phases respectively...... 50

2.11 Critical temperature as a function of tetrahedral angle for Ln-1111 type materials, adapted from Ref. [39], as well as those determined for LaFe0.9Rh0.1AsO, NdFe0.9Rh0.1AsO (Ref. [28]), and SmFe0.9Rh0.1AsO (Ref. [29])...... 52

3.1 The structure of LnMPnO compositions portrayed as a stack of alternating MPn layers (antifluorite like) and LnO layers (fluorite like) along the c-axis...... 61

3.2 Powder x-ray diffraction patterns for LnRhAsO compositions where Ln = La, Ce, Nd. Open diamonds (◊) indicate the presence of RhAs while closed diamonds () represent CeO2-x phases...... 64

3.3 Powder x-ray diffraction patterns for LnIrAsO compositions where Ln = La, Ce, Nd. Open diamonds (◊) indicate the presence of IrAs2 while closed diamonds () represent Ir...... 64

3.4 Lattice parameters for LnRhAsO and LnIrAsO compositions...... 66

3.5 The structure of LnRhAsO and LnIrAsO materials highlighting the LnAs bond distances (a) and the MAs layer as viewed along the c-axis (b), detailing the M-M distance and its relationship to the unit cell a- parameter...... 67 LIST OF FIGURES (Continued)

Figure Page

3.6 Unit cell volumes and metal-metal (M-M) distances for LnRhAsO and LnIrAsO compositions...... 68

3.7 The relationships between M-As bond distance, a-parameter, As atomic position (zAs), and tetrahedral angles for LnMAsO materials...... 72

3.8 Tetrahedral bond angle, As-M-As, and As atomic position, zAs, evolution as a function of lanthanoid element for LnRhAsO and LnIrAsO compositions...... 73

3.9 Normalized resistivity, ρ/ρ300K, for LnRhAsO compositions (a) and LnIrAsO compositions (b)...... 74

3.10 Normalized resistivity, ρ/ρ300K, for the Ce based compositions CeRhAsO and CeIrAsO, as well as CeNiAsO as reported in Ref. [8]...... 75

3.11 A qualitative depiction of the Kondo lattice, TK, and RKKY temperature scales, TRKKY, as function of the parameter N(EF)J. Within the magnetic Kondo state (Mag. KS) the magnetic ordering temperature is labeled TM...... 77

3.12 The 4f resistivity for CeRhAsO and CeIrAsO compositions as well as the observed magnetic ordering temperatures, TN, and 푚푎푥 Kondo lattice temperature,푇퐾 ...... 80

3.13 Low temperature AC molar magnetic susceptibility and resistivity plots for CeRhAsO...... 81

3.14 Low temperature AC molar magnetic susceptibility and resistivity plots for CeIrAsO...... 82

3.15 DC magnetic susceptibility of LnRhAsO compositions. Inset is a table listing the results of a Curie-Weiss fitting,

휒 푚표푙 = 퐶/(푇-휃), for Ce and Nd based compositions...... 83

LIST OF FIGURES (Continued)

Figure Page

3.16 DC magnetic susceptibility of LnIrAsO compositions. Inset is a table listing the results of a Curie-Weiss fitting, 휒푚표푙 = 퐶/(푇 − 휃), for Ce and Nd based compositions...... 84

3.17 The positions of CeRhAsO, CeIrAsO, and CeNiAsO within the 푚푎푥 N(EF)J plot based on their observed 푇퐾 and TN values. Values for CeNiAsO were taken from Ref. [20]...... 86

3.18 The Hall coefficient temperature variation of LaRhAsO and CeRhAsO...... 88

3.19 band structure of LaRhAsO (a), total and partial atomic density of states (b), partial density of states for d orbitals only (c), and magnified view near the Fermi level (d). Band structure from energy minimized atomic positions is shown in black and red dashed lines indicate that determined using atomic positions from Rietveld refinement of PXRD data...... 90

3.20 Electron band structure of LaIrAsO (a), total and partial atomic density of states (b), partial density of states for d orbitals only (c), and magnified view near the Fermi level (d). Band structure from energy minimized atomic positions is shown in black and red dashed lines indicate that determined using atomic positions from Rietveld refinement of PXRD data...... 91

4.1 Crystal structure of La2PnO2 compositions (where Pn = Sb and Bi). The unit cell (a) contains O in tetrahedral coordination to La cations creating fluorite like slabs while the Sb anions adopta square prismatic coordination to La between the fluorite slabs...... 105

4.2 The lattice parameters for LaFeAs1-xSbxO and LaFe1-xZnxSbO compositions...... 108

4.3 The experimentally observed PXRD pattern for La:Fe:Sb:O reactions along with the simulated patterns for LaFeSbO (P4/nmm) and La2SbO2 (I4/mmm)...... 110

LIST OF FIGURES (Continued)

Figure Page

4.4 Phonon spectrum for LaFeSbO as found using the DFT energy minimized structural parameters, a = 4.168Å c = 9.214Å zLa = 0.1283 and zSb = 0.6595...... 113

4.5 The PXRD patterns for La2SbO2-x and La2BiO2-x compositions where nominally x = 0.5. Also shown are the general atomic positions for La2PnO2 materials and refined lattice parameters for each composition...... 115

4.6 The powder neutron diffraction pattern for the nominal composition La2SbO1.5. Refinement of atomic site occupancy results in a composition of La2SbO1.7...... 121

4.7 The powder neutron diffraction pattern for the nominal composition La2BiO1.5. Refinement of atomic site occupancy results in a composition of La2BiO1.9...... 121

4.8 Natural log of resistivity for La2SbO1.5,La2BiO1.5, La2Bi0.75Sb0.25O1.5, and La1.9Sr0.1SbO1.5 compositions. The slope, T0, is related to the band gap, T0 = Eg/2kB, and the calculated Eg values at both high and low temperatures are shown...... 122

4.9 Natural log of resistivity for La2SbO1.5,La2BiO1.5, and -1/4 La2Bi0.75Sb0.25O1.5 plotted as a function of T . The value T0 is obtained by evaluating the slope of the linear best fit line...... 124

4.10 The band structure for La2SbO2. A primitive cell was used for calculation cell rather than I-centered due to computational ease. The Brillouin zone is shown in the upper corner...... 125

LIST OF TABLES

Tables Page

1.1 Unit cell lattice parameters, c/a ratios, unit cell volumes, and M-M distances for all known LnMAsO compositions. Adapted from the reviews given in Ref. [5] and [11], and updated to include results from more recent investigations...... 10

3.1 Lattice parameters for LnCoAsO, LnRhAsO, and LnIrAsO compositions as determined by powder x-ray diffraction...... 66

3.2 Primary results of powder x-ray diffraction Rietveld refinement for LnRhAsO and LnIrAsO compositions...... 69

4.1 Zero temperature heat of formation energy per formula unit of LaFePnO relative to the standard state of the constituent elements and relative to more common starting reagents...... 111

4.2 Zero temperature heat of formation energy per formula unit of LaFePnO relative to the most stable competing phases...... 112

4.3 Lattice parameters determined for La2SbO2 and La2BiO2 compositions ...... 117

4.4 Powder neutron diffraction Rietveld refinement results for La2SbO1.5 and La2BiO1.5 nominal compositions...... 120

LIST OF APPENDIX FIGURES

Figure Page

A.1 Schematic of the setup used to synthesize NdFeAsO and NdFe0.9Co0.1AsO issuing reaction ampoules surrounded by an additional susceptor within a 2.45 GHz multi-mode microwave cavity...... 151

A.2 X-ray powder diffraction patterns for NdFeAsO and NdFe0.9Co0.1AsO synthesized using microwave radiation and an additional susceptor, CuO, surrounding the reaction ampoule. Shown right is the structure of Nd(Fe/Co)AsO materials depicted as layers of Nd2O2 and Fe2As2 stacked along the c-axis...... 154

A.3 Magnetic susceptibility of NdFe0.9Co0.1AsO collected under zero field cooled conditions with a magnetic field of 20 Oe; shown plotted as 4πχvol. The critical temperature, Tc, as determined by the dχvol/dT plot is 16 K; shown overlaid for comparison ...... 155

A.4 Normalized resistivity (ρ /ρ25K) of NdFe0.9Co0.1AsO collected under magnetic fields from 0 - 6 Tesla (T). The inset details how the critical temperatures for 10%, 50%, and 90% of the normal state resistivity change with increasing magnetic field; the rate of the change for each is listed below the label...... 156

B.1 A schematic detailing the four primary components of our superconductor outreach activities as part of the 2011 centennial anniversary of superconductivity...... 162

B.2 Details of the free superconductor demonstration kits which were made available to educators as part of the 2011 centennial anniversary of superconductivity and the 25th anniversary of high temperature superconductor materials...... 167

To Afton, your love, support, and encouragement is what keeps me going, and to Ella Mae and Moses, through you I have come to know true joy in life.

Structure-property Relationships of Layered Oxypnictides

Chapter 1

General Overview and Introduction to Layered Oxypnictides

1.1 Introduction The need for solid state materials is as ancient as the first human desire for better tools and yet as modern as human aspirations to colonize the moon. Solid state materials are everywhere and to know how they form, react, and respond to external stimuli is essential to understanding geology, life processes, and modern technology. The culmination of hundreds of years of scientific thought and experiment has led to the realization that the energy, structure, and properties of matter are deeply intertwined. To improve the solid state materials used in our lives we must search for connections between composition, atomic arrangement, and physical properties.

Many of the advanced solid state materials used in our lives today are found in electronics and computers. Society has been ushered into the ‘computer age’ of ever shrinking transistors and diodes in part because materials scientists and engineers have learned so much about the interconnectedness of composition, 2 structure, and properties. The dependence of modern society on electronic devices provides a current and continuing need to better understand the electronic properties of solid state materials and to develop new materials with electronic behaviors as of yet unseen. A major part of materials science is therefore focused on characterizing how electrons within a material act and how this changes as a function of some other stimulus such as thermal energy, pressure, or magnetic field.

Another major focus of materials science in general, is the investigation of new elemental combinations and spatial arrangements. For as much as we seem to know about designing materials with specific properties scientists are still largely unable to predict when certain electronic properties will exist in a material, one great example being superconductivity. Efforts to explore the properties of new compounds and atomic structures are therefore fundamentally important because each new stable arrangement of elements uncovered offers the possibility of discovering new properties which cannot be predicted a priori, superconductivity being just one.

Both synthesis and characterization aspects of materials science can be largely summed up as a game of invoking phase changes; solid to liquid changes, polymorph changes, structural symmetry changes, or electronic structure changes

(e.g. electron ordering, electron localization-delocalization, electron pairing). Most 3 often these phase changes are explored through control of temperature, and thus thermal energy, kBT. Naturally, the temperature at which a phase change occurs also depends on the composition of the materials system, i.e. its chemical energy. This means scientists have two key ways of exploring for new materials and interesting properties. One is to change the temperature of a material system while measuring its properties, perhaps electronic resistance or x-ray diffraction peak position, and another is to change the chemical composition then measure its properties. In reality both of these knobs, temperature and composition, as well as others, such as pressure and external magnetic field, are tuned in the search for interesting phenomena.

Layered oxypnictide compositions with ZrCuSiAs structure present many great examples of how changes in chemical composition between compounds with the same basic structural arrangement can dramatically alter physical properties. In particular the discovery of superconductivity up to 55 K for LaFeAsO (F-doped) provides yet another example of a layered atomic structure leading to record setting critical temperatures, similar to the cuprates. In addition the LaFeAsO discovery is a great example of how it is not always entirely new structures which turn out to have the most interesting properties. In other words, materials discovery is not simply about uncovering entirely new compositions. Sometimes the rediscovery and subtle 4 compositional tuning (F doping for LaFeAsO) of known materials can also yield exciting results.

The body of this thesis will explore the some of the property changes which occur for LaFeAsO as Fe is replaced by Rh. It will also introduce and discuss the physical properties of new LnMAsO compositions where Ln = La, Ce, and Nd and

M = Rh, Ir. In addition, the search for LaFeSbO, an as of yet observed layered oxypnictide, and the subsequent discovery of the layered oxypnictides, La2PnO2 where Pn = Sb and Bi, will be discussed. Before discussing these research results a general introduction to LnMAsO materials is given. Introduction to La2PnO2 materials is covered in Chapter 4 where these results are presented.

1.2 Introduction to LnMAsO materials (where Ln = Lanthanoid, M = late ) Most recently it has been the discovery of superconductivity up to 55 K for doped LnFeAsO compositions which has sparked a wildfire of research into LnMPnO

(Pn = P, As, and Sb) materials and has rekindled the scientific community’s fascination with similarly layered compositions such as BaFe2As2 and LiFeAs [1–4]. All these

LnMPnO materials actually belong to much larger class of compositions referred to as

ZrCuSiAs type, for which approximately 150 different compositions are known [5].

The ZrCuSiAs structure as first reported by Johnson and Jeitschko in 1973 and 5 described as a filled variant of the PbFCl structure [6]. The PbFCl structure may be regarded as a cubic close packed array of Pb with Cl atoms filling all octahedral positions and F atoms in one-half the tetrahedral, Td, voids in alternate layers; which leaves one-half the tetrahedral positions vacant. As described by these researchers, it was known that the compound NbSiAs adopted the PbFCl structure and could be

5+ 2- 3- described as Nb [Si] As where the Si atoms occupy ½ of the Td sites and the brackets indicate significant bonding between Si atoms within a layer. Johnson and

Jeitschko went on to suspect that another mono-valent cation, e.g. Cu1+, could occupy the vacant tetrahedral positions if the Nb5+ in NbSiAs was compensated with a 4+ cation such as Zr4+ or Hf4+. Thus the compounds ZrCuSiAs and HfCuSiAs were discovered and the “filled” PbFCl structure was reported.

The ZrCuSiAs structure can be described as having P4/nmm (#129) symmetry and in fact there are two settings for this space which are related by a (¼, -¼,

0) translation of the origin [7]. Within the literature it is commonly setting number 2 which is used and atoms are placed at Zr (¼, ¼, zZr), Cu (¾, ¼, ½), Si (¾, ¼, 0), and As

(¼, ¼, zAs). Additionally the structure is often described in terms of the anti-fluorite like and fluorite like layers which form. Fluorite layers have cations in tetrahedral coordination to anions while anti-fluorite layers have anions in tetrahedral coordination to cations. Figure 1.1 details the atomic arrangement of LnMPnO compositions which have the ZrCuSiAs type structure. 6

The upper portion shows how the LnMPnO structure can be viewed as an evolution from a simple close packed LnPn lattice (NaCl type structure) to one with tetrahedral sites filled and finally to one where the Ln and Pn sublattices undergo a tetragonal distortion and distinct layers are formed around the two tetrahedral positions. The lower portion of Figure 1.1 details the fluorite and anti-fluorite layers which result as well as the M-M bond distances within the MPn layer.

Figure 1.1 The evolution from simple LnPn structure to the LnMPnO structure (upper). Atomic positions, formation of fluorite and anti-fluorite layers, and the M-M bonding which occurs within the MPn layers for LnMPnO materials (lower). 7

Using a valence compensated approach to discover new compositions while maintaining a particular structure type is a very common technique used in solid state chemistry and in the case of ZrCuSiAs it can be taken in many different directions to synthesize numerous , fluorides, and in this family of compounds with this general structure. Compounds with simple ionic formulas such as Th4+Ag1+P3-O2-, Ce3+Co2+Si4-H1−, La3+Ag1+S2-O2-, Ba2+Cu1+S2-F-, Ba2+Fe2+As3-F-, and

La3+Fe2+As3-O2- have all been reported [5]. Considering their structure they are often rewritten to highlight the fluorite and anti-fluorite layering and the relative charge on each layer; [Th4+O2- ]2+[Ag+P3-]2-, [Ce3+H1−]2+[Co2+Si4-]2-, [La3+O2-]1+[ Ag+S2-]1-, [Ba2+F-]1+[

Cu1+S2-]1-, [Ba2+F-]1+[ Fe2+As3-]1-, and [La3+ O2-]1+[ Fe2+As3-]1-. For LnMPnO compositions the simple ionic binding scenario does not hold up entirely, as it does for BaCuSF for instance, and this will be discussed below through the use of bond valence analysis.

While LnMPnO compounds are stabilized by between layers, the charge transfer between the layers is significantly lower than the purely ionic picture. For instance computational studies have shown that in LaFePO, LaNiPO, and

LaFeAsO the layers have charges closer to [LaO]~0.35+[MPn]~0.35− [8–10]. This is partly because the M-Pn bonding within the anti-fluorite layers has a more covalent character than the largely ionic La-O bonding within the fluorite layers or the La-Pn bonding between the layers. Additionally the overall chemical bonding is further complicated by the presence of metal-metal interactions. Since the d shells of M2+ 8 ions M = Mn - Ni are partially filled and the M-M distances are relatively short these interactions can be bonding in nature and this M-M bonding in turn affects the M-Pn distances. As will be discussed throughout this work these interactions are all reflected in the observed bond lengths, bond valence sums, and band structures for these layered transition metal oxypnictides and justifies writing their general formulas as [La3+O2-]x+[MPn]x- where x ≈ 0.3-0.4 if the structural bonding is to be highlighted.

Narrowing our discussion further to just , LnMAsO, there are still many compositions known and they display a number of very interesting properties.

Table 1.1 below lists all the known LnMAsO compositions similar to what was done in previous reviews [5,11]; however, it has been updated to include data (or duplicate data) for a number of compositions which had not been reported at the time of the aforementioned reviews.

First off, by examining these tabulated data it can be seen that no LnMPnO compositions have been reported for transitions before Mn. Starting at Mn and moving across to Zn it can be seen that on transitioning from Mn to Fe there is a substantial decrease in the a-parameter as well as the c-parameter. However upon moving to Co the a-parameter increases once again while the c-parameter continues to contract substantially. By the time Ni compositions are reached the a-parameter 9 is once again nearly equal to that of Mn compositions however the c-parameter is reduced by almost 10%. The only Cu composition known contains Cu1+ and thus is valence compensated with a Th4+ rather than Ln3+. Moving from Ni to Zn however the c-parameter’s downward trend is broken and a and c lattice parameters very similar to Mn compositions are once again observed (i.e. the c-parameter suddenly jumps back up).

2+ Simply considering the ionic radii of M cations in tetrahedral, Td, coordination [12] the jump in lattice parameters is be expected for Zn compositions as its radii is known to increase relative to Fe - Ni. On the other hand the a- parameter’s expansion on moving from Fe - Ni might seem counterintuitive as Fe has the largest ionic radii; however, if the M-As bond length is considered then the reason for this trend becomes more obvious. Looking at just the Nd compositions across the series NdMAsO, where M = Mn - Zn, it is observed that the M-As bond lengths do in fact follow the ionic radii trend nicely. That is to say, NdMnAsO and

NdZnAsO have the largest ionic radii and the longest M-As bond lengths (2.546 Å and

2.559 Å respectively) [13,42]. Moving to NdFeAsO there is a relatively large decrease in the M-As bond length (2.399 Å) [28] and this decrease in M-As distance continues with NdCoAsO (2.352 Å) [34]. NdNiAsO has not been reported however the Ni-As bond length in LaNiAsO is ~2.346 Å [17] and thus in NdNiAsO it should be 2.31-

2.33 Å, based on the trend between other LaMAsO and NdMAsO compositions. 10

Table 1.1 Unit cell lattice parameters, c/a ratios, unit cell volumes, and M-M distances for all known LnMAsO compositions. Adapted from the reviews given in Ref. [5] and [11], and updated to include results from more recent investigations.

a-parameter c-parameter c/a Volume M-M Composition std. std. Reference (Å) (Å) ratio (Å3) (Å) YMnAsO 3.957 1 8.75 6 2.211 137.0 2.798 [13] 4.124 1 9.03 5 2.190 153.6 2.916 [13] LaMnAsO 4.11398 1 9.03044 2 2.195 152.8 2.909 [20] CeMnAsO 4.086 1 8.956 2 2.192 149.5 2.889 [13] PrMnAsO 4.067 1 8.919 3 2.193 147.5 2.876 [13] 4.049 2 8.893 1 2.196 145.8 2.863 [13] NdMnAsO 4.04396 7 8.8787 3 2.196 145.2 2.860 [20] 4.04979 5 8.89935 2 2.197 146.0 2.864 [21] 4.02 1 8.829 3 2.196 142.7 2.843 [13] SmMnAsO 4.0177 8.819 2.195 142.4 2.841 [22]

GdMnAsO 3.989 1 8.805 3 2.207 140.1 2.821 [13] TbMnAsO 3.978 1 8.743 4 2.198 138.4 2.813 [13] DyMnAsO 3.959 1 8.727 4 2.204 136.8 2.799 [13] UMnAsO 3.869 1 8.525 2 2.203 127.6 2.736 [13]

4.03552 8 8.7393 2 2.166 142.3 2.854 [1] 4.032 8.726 2.164 141.9 2.851 [23]

4.03007 9 8.7368 2 2.168 141.9 2.850 [24] LaFeAsO 4.03397 4 8.7449 4 2.168 142.3 2.852 [25] 4.038 1 8.753 6 2.168 142.7 2.855 [26] 4.03268 1 8.74111 4 2.168 142.2 2.852 [27] 4.0367 1 8.7218 4 2.161 142.1 2.854 [28] 4.000 1 8.655 1 2.164 138.5 2.828 [26] CeFeAsO 4.0058 1 8.6289 6 2.154 138.5 2.833 [28] 3.996 8.648 2.164 138.1 2.826 [29]

3.985 1 8.595 3 2.157 136.5 2.818 [26] PrFeAsO 3.9889 1 8.5966 8 2.155 136.8 2.821 [28] 3.97716 5 8.6057 2 2.164 136.1 2.812 [30] 3.965 1 8.575 2 2.163 134.8 2.804 [26] NdFeAsO 3.9713 1 8.5655 5 2.157 135.1 2.808 [28] 3.96594 1 8.59786 5 2.168 135.2 2.804 [31] 11

Table 1.1. Unit cell parameters for LnMAsO compositions continued. 3.940 1 8.496 3 2.156 131.9 2.786 [26] 3.9469 2 8.4965 6 2.153 132.4 2.791 [28] SmFeAsO 3.933 5 8.495 4 2.160 131.4 2.781 [2] 3.94 8.496 2.156 131.9 2.786 [32]

3.915 1 8.435 4 2.155 129.3 2.768 [26] GdFeAsO 3.9199 3 8.4451 8 2.154 129.8 2.772 [28] 3.9043 3 8.408 1 2.154 128.2 2.761 [28] TbFeAsO 3.898 1 8.404 2 2.156 127.7 2.756 [33]

LaRuAsO 4.119 1 8.488 1 2.061 144.0 2.913 [26] CeRuAsO 4.096 1 8.380 3 2.046 140.6 2.896 [26] PrRuAsO 4.085 1 8.337 1 2.041 139.1 2.889 [26] NdRuAsO 4.079 1 8.292 2 2.033 138.0 2.884 [26] SmRuAsO 4.050 2 8.191 7 2.022 134.4 2.864 [26] GdRuAsO 4.039 1 8.118 6 2.010 132.4 2.856 [26] TbRuAsO 4.027 1 8.078 1 2.006 131.0 2.848 [26] DyRuAsO 4.022 2 8.05 3 2.001 130.2 2.844 [26]

4.054 1 8.472 3 2.090 139.2 2.867 [26] LaCoAsO 4.0526 1 8.462 4 2.088 139.0 2.866 [15] CeCoAsO 4.015 1 8.364 2 2.083 134.8 2.839 [26] PrCoAsO 4.005 1 8.344 2 2.083 133.8 2.832 [26] 3.982 1 8.317 4 2.089 131.9 2.816 [26] NdCoAsO 3.98423 8 8.3333 3 2.092 132.3 2.817 [34] SmCoAsO 3.957 3 8.242 2 2.083 129.1 2.798 [35]

LaRhAsO 4.1226 2 8.4692 6 2.054 143.9 2.915 [36] CeRhAsO 4.089 1 8.365 2 2.046 139.9 2.891 [37] NdRhAsO 4.063 1 8.303 1 2.044 137.1 2.873 [37]

LaIrAsO 4.087 1 8.557 1 2.094 142.9 2.890 [37] CeIrAsO 4.063 1 8.444 2 2.078 139.4 2.873 [37] NdIrAsO 4.046 1 8.35 1 2.064 136.7 2.861 [37]

LaNiAsO 4.12309 1 8.18848 6 1.986 139.2 2.915 [17] 12

Table 1.1. Unit cell parameters for LnMAsO compositions continued. LaNiAsO 4.119 8.18 1.986 138.8 2.913 [38]

4.0767 8.1015 1.987 134.6 2.883 [39] CeNiAsO 4.0752 2 8.1134 4 1.991 134.7 2.882 [40]

ThCuAsO 3.9614 5 8.44 1 2.131 132.4 2.801 [41]

3.943 1 8.843 3 2.243 137.5 2.788 [42] YZnAsO 3.9423 3 8.84 1 2.242 137.4 2.788 [14] 4.095 1 9.068 3 2.214 152.1 2.896 [42] LaZnAsO 4.10492 13 9.08178 47 2.212 153.0 2.903 [43] 4.0956 3 9.07 1 2.215 152.1 2.896 [14] 4.069 1 8.995 3 2.211 148.9 2.877 [42] CeZnAsO 4.0679 5 8.998 1 2.212 148.9 2.876 [43] 4.0665 6 9.001 2 2.213 148.8 2.875 [14] 4.047 1 8.963 1 2.215 146.8 2.862 [42] PrZnAsO 4.0477 2 8.966 2 2.215 146.9 2.862 [43] 4.049 4 8.972 2 2.216 147.1 2.863 [14] 4.030 1 8.949 4 2.221 145.3 2.850 [42] NdZnAsO 4.0295 6 8.952 1 2.222 145.4 2.849 [43] 4.0309 4 8.957 1 2.222 145.5 2.850 [14] 4.003 1 8.903 2 2.224 142.7 2.831 [42] SmZnAsO 4.0002 4 8.902 2 2.225 142.4 2.829 [14] 3.976 1 8.894 3 2.237 140.6 2.811 [42] GdZnAsO 3.9762 4 8.892 2 2.236 140.6 2.812 [14] 3.957 1 8.841 2 2.234 138.4 2.798 [42] TbZnAsO 3.9625 5 8.852 2 2.234 139.0 2.802 [14] 3.947 1 8.838 1 2.239 137.7 2.791 [42] DyZnAsO 3.9496 4 8.837 1 2.237 137.9 2.793 [14] HoZnAsO 3.9379 4 8.829 1 2.242 136.9 2.785 [14] ErZnAsO 3.9272 3 8.815 1 2.245 136.0 2.777 [14]

LaCdAsO 4.219 2 9.23 3 2.188 164.3 2.983 [44] CeCdAsO 4.191 1 9.171 4 2.188 161.1 2.963 [44] PrCdAsO 4.172 1 9.136 4 2.190 159.0 2.950 [44] NdCdAsO 4.151 1 9.123 5 2.198 157.2 2.935 [44] 13

The overall trend in lattice parameters can be understood to come from competition between M-M bonding and M-As bonding and bond valence sum (BVS) analysis is insightful [18,19]. Bond length values for M-As distances can be used to calculate the expected cation valency through the use of tabulated bond valences.

The tabulated bond valence values are all determined from ionic compounds with clearly established valence states. In theory, the sum of each cations individual bond valence values should match well with the expected from a purely standpoint. In the case of LnMPnO compositions these oxidation states would be Ln3+M2+ Pn3-O2- as mentioned above.

푅0−푅 The individual BV values are found by using the formula BV = 푒 퐵 , where R0 and B are tabulated values for different cation-anion combinations and R is the experimental bond length (these values are typically tabulated for using bond lengths with units of Å) [18,19]. Since the M cations have four equivalent bonds to As in these materials the BVS value for Mn, BVSMn should be four times the bond valence

푅0−푅 퐵 found for an individual Mn-As bond. For NdMnAsO, BVSMn = 4 × 푒 =

2.36−2.55 4 × 푒 0.37 = 2.42. This value is fairly close to the expected value of 2 for a M2+ cation, which indicates that the M-As bonding is fairly ionic and that the M-M interaction must be minimal. However upon moving to Fe - Ni compositions (still with Nd) the observed BVS values increase substantially; BVSFe = 3.50, BVSCo = 3.29, 14

BVSNi = 3.22 respectively. This trend correlates well with the a-parameter observations, i.e. NdFeAsO has the shortest a-parameter and the strongest deviation from a M2+ state. The Co-As bond is expected to be more covalent, owing to the smaller difference in between M and As, therefore it might be expected that the calculated BVS values for Co should increase. The decrease in BVS observed therefore highlights the fact that M-M bonding is also affecting the M-As bond distance and BVS outcome. Another simultaneous consequence is that the lanthanoid atoms shift and the unit cell contracts along the c-axis so that the Ln-As and Ln-O bond preferences can be maintained. Upon getting to NdZnAsO the BVS is once again fairly close to that expected for ionic bonding, BVSZn = 1.69, and the M-M interaction must minimal.

Looking at the tabulated data for LnMAsO compositions it can also be seen that there are comparatively fewer 4d compositions known and that LnIrAsO compounds (reported in this body of work) are the only known compositions to contain a 5d metal. Both the a and c lattice parameters for 4d compositions are generally larger than those of their 3d counterparts, the exception being Co and Rh based compositions where the a-parameter is expanded but the c-parameter remains relatively unchanged. The expansion of lattice parameters is expected, given that a 4d metal ion will be considerably larger than a 3d ion thereby increasing both the M-M and M-As distances, but upon moving from 4d to 5d metal (Rh to Ir) the 15 a-parameter actually contracts. This structural change along with the physical properties of Rh and Ir compositions will be discussed further in Chapter 3.

Low dimensionality can often promote electronic correlation and this is partially why LnMAsO materials are known to possess rich electronic phase diagrams.

In addition it is the chemically flexibility of this structure which helps to make it a wonderful playground for exploring the effects of chemical pressure and aliovalent substitution. Many of these LnMAsO compositions had been only structurally characterized within the literature [13,26,42] since they were first discovered over 10 years ago until recently following the LnFeAsO superconductivity discoveries.

Looking at just LnMAsO compositions, where M = Mn - Zn, a number of interesting properties can be observed.

The parent materials LnMnAsO (where Ln = La - Sm) are all reported to be semiconducting; however, when doped with electrons or holes an insulator to metal transition can be induced [22,45]. Neutron diffraction experiments have shown that

LaMnAsO and NdMnAsO exhibit antiferromagnetic (AFM) ordering of Mn moments within the a-b plane weak but ferromagnetic (FM) ordering along the c-axis direction at room temperature [46]. Two later studies demonstrated that at low temperatures

(~24 K) Nd moments will align in an AFM order within the basal plane and at this temperature the Mn moments also realign to lie within the basal plane parallel to Nd 16 moments [21,20]. While neutron diffraction studies have not been reported yet for

CeMnAsO it has been demonstrated through magnetic susceptibility and heat capacity studies that there is also a change in the electronic state at about 34 K which is most likely due to spin reorientation [47]. Additionally it was reported that for

NdMnAsO compositions there is an approximately 10% drop (i.e. magneto-resistance ratio, MR ~ -10) in resistivity at room temperature, but only when the Nd position is slightly deficient (i.e. ~0.97 occupation fraction) [20].

It was the discovery of superconductivity at 26 K and then up to 55 K for

LnFeAsO1-xFx (x = 0 - 0.2) that kick started a worldwide investigation into the properties of LnMPnO materials and other structurally related compositions containing the FeAs anti-fluorite layer [1-4,48,23,49,32,29,50–52,36,53]. The parent compounds LnFeAsO are known to exhibit nonlinear resistivity temperature dependence.

The change in resistance with temperature for a material, i.e. the slope of a resistance vs. temperature plot, is commonly referred to as the temperature coefficient of resistance (TCR). For metals it is positive, as their resistance increases with increasing temperature. For semiconductors, it is negative as their resistance decreases with increasing temperature. 17

For LnFeAsO materials the TCR is negative from approximately 160 - 300 K, and then becomes positive from approximately 50 - 160 K. Finally it is once again negative at low temperatures; see for instance Figure 2.4 which shows the resistance of LaFeAsO as part of the LaFe1-xRhxAsO solid solution. This sort of nonlinear behavior has been attributed to the formation of an AFM spin density wave (SDW) state coinciding with a subtle transition from tetragonal (P4/nmm) to orthorhombic

(Cmma) symmetry [25,48]. Doping of the parent compounds has been shown to suppress the spin density wave and structural transition and to shift the ground state from AFM to superconducting [24,29,54,55]

Currently the mechanism by which electrons pair in LnFeAsO materials is not precisely known and as such is vaguely called ‘unconventional’. This essentially means that BCS theory and phonon coupling cannot explain their high critical temperatures. In fact there are a number of similarities between the , La2CuO4, and the iron pnictides, LaFeAsO1-xFx and CeFeAsO1-xFx

[56-59]. The parent compounds each undergo AFM ordering as the temperature is lowered and in both materials, with increased doping, the AFM temperature, TNeel, drops and then vanishes just before superconductivity emerges. One key difference is that La2CuO4 is semiconducting, while LaFeAsO and CeFeAsO are metallic (or semi metallic). While the exact details of the superconducting state, and its possible root cause, are beyond the scope of this introduction to LnMAsO materials; however a 18 few more words about their superconducting state should be mentioned along with a few background words on superconductivity.

There are two fundamental length scales used to describe the superconducting state. These parameters are coherence length, ζ, and magnetic penetration depth, λ, also known as the London depth [60–62]. The magnetic penetration depth describes how far the magnetic field extends into a superconductor material. An exact definition of coherence length is a little bit more complex. However since 1957, when Bardeen, Cooper, and Schrieffer put forth BCS theory [63], the coherence length is often simply described as the interaction distance between electrons in a Cooper pair. Within BCS theory electron pairs

(Cooper pairs) form, and thereby achieve a resistance free state, via their interaction with lattice vibrations, i.e. phonons. The magnetic penetration depth and coherence length are important because their ratio can be related to the type of superconductivity observed.

In addition another important parameter in the discussion of superconductors is their critical magnetic field, Hc. As the name might seem to imply, this is the magnetic field at which superconductivity is destroyed. Most elemental superconductors exhibit type 1 behavior, the first type discovered. These have one critical field, Hc1, and will abruptly transition from the superconducting to normal 19 state when it is reached. They are characterized as having a magnetic penetration depth to coherence length ratio of less than 1/2 and a positive surface free energy at a superconductor to normal material interface.

The most important class of superconductors is called type 2 and they are mainly alloys or compounds. This type can form a mixed superconductor-normal state, and therefore have two critical fields, Hc1 and Hc2. The lower one, Hc1, marks the onset of a mixed state and the upper one, Hc2, marks the complete transition to a normal state. For type 2 superconductors the magnetic penetration depth to coherence length ratio is greater than 1/2 and there is a negative surface free energy at a superconductor to normal interface. This negative free energy causes the magnetic flux that does enter a type 2 superconductor to disperse into many flux lines so as to maximize the superconductor and normal state contact area. So long as these threads, also called fluxions, don’t move too much a superconducting state can be maintained around them. Each fluxion can also be described as having a superconductor outer shell and normal core. The magnetic penetration depth is the overall radius of a fluxion and the coherence length is the radius of its normal core.

As the strength of the magnetic field increases the radius of these fluxions increases until eventually, at the upper critical field, the normal cores overlap, destroying superconductivity altogether. The ability to have pinned fluxions is what gives type 2 materials technological importance. The pinned flux threads essentially quarantine 20 the penetrating field and allow a superconducting state to be maintained even in very high fields for some type 2 materials.

As detailed in Figure 2.8 for LaFe0.9Rh0.1AsO, the upper of critical field and coherence length can be estimated from resistivity data using the Werthamer-

Helfand-Hohenberg (WHH) model [64]. The WHH model says that the upper critical magnetic field at zero Kelvin can be estimated as Hc2(0) = -0.693(Tc)(dHc2/dT)T=Tc .

The value used for (dHc2/dT)T=Tc is obtained by measuring the resistivity as a function of temperature and magnetic field and calculating the change in Tc as function of magnetic field. Ginzburg-Landau theory is known to relate the superconducting

Cooper pair coherence length, ξ, to the upper critical field [60]. The relationship is,

2 Hc2 = Φ0 / π ξ , where Φ0 represents a quantum of magnetic flux, or one fluxion, and equals approximately 2.07 x 10-7 Oe cm2. Using the WHH model, critical fields for polycrystalline LnFeAsO samples are found to be anywhere from 20 - 100 T and thus coherence lengths range from about 25 - 60 Å [52,65,66,67]. Magnetic penetration depths reported for LaFeAsO0.9F0.1 are approximately 1500 - 2000 Å [68,69]. It should be pointed out however that this is a crude estimate which does not take into consideration anisotropy of the material, something which must be considered if intrinsic properties are to be precisely determined. 21

Looking next at LnCoAsO compositions an itinerant FM state is observed

[15,51,70] rather than a semimetal AFM state as with LaFeAsO. The reason for this itinerant FM state can be related to the density of states at EF for LaCoAsO, as briefly mentioned below, and discussed in the context of LaRhAsO and LaIrAsO band structures (Chapter 3). Detailed electronic and magnetic studies of NdCoAsO and

SmCoAsO reveal a very complicated phase diagram where the magnetic ordering can be either FM or AFM depending on temperature as well as field strength [34,71].

For LnNiAsO compositions once again superconductivity is observed [17]. This time however the Tc is greatly reduced and the undoped parent compound displays superconductivity. It is generally understood that the mechanism of electron pairing in LaNiPO is phonon mediated (i.e. Bardeen-Cooper-Schrieffer (BCS) theory can explain the results) [72]; however in the case of LaNiAsO1-xFx it has been suggested that BCS theory cannot explain the large jump in heat capacity observed at the Tc [38] and thus it is in the strong coupling regime. CeNiAsO has been reported to display an

AFM Kondo lattice ground state where the 3d conduction band carriers undergo

Kondo scattering with the Ce 4f electrons from around 20 - 300 K [39]. At still lower temperature the Ce moments order and the 3d carriers are no longer scattered as strongly. The properties of CeNiAsO and how they relate to CeRhAsO and CeIrAsO compositions will be further discussed in Chapter 3. 22

Electronic and magnetic properties for any material are dependent on the bonding character (e.g. atomic distances or isotropic vs. anisotropic spatial arrangement) as well as the character of the electronic states close to the Fermi level. Therefore a brief word should be said about the band structure of LnMAsO materials before discussing research results. In Fe, Co, and Ni based compositions the Fermi level is dominated by orbital contributions from the transition metal cations and thus observed physical properties can change dramatically upon metal ion substitution as pointed out above. Band structure calculations of LaFeAsO show it to be a magnetic semimetal [73].

Figure 1.2 schematically details the difference between the band filling of a metal, a semimetal, and a semiconductor. Essentially a semimetal may be considered to be a semiconductor where the filled valence band is shifted up in energy and the conduction band is shifted down in energy, thus causing them to cross the Fermi level. Electrons are able to then move from the filled band valence band to the empty one and a metallic state emerges, but with reduced mobility. This means that semimetals can be multiband metals, with more than one active carrier type possible due to the presence of both hole and electron bands which cross the

Fermi level. 23

Before looking at the band structure of LaFeAsO it should also be reminded that carrier velocity, v, can be related to the slope, dE/dk, of the E vs k-space plot, via v = hk/2πm = (2π/h)(dE/dk), where h is Planck’s constant and k represent the crystal momentum, or coordinates in momentum space. The carrier effective mass, m*, is related to the inverse of the curvature, (second derivative, d2E/dk2) of the E vs k plot, m* = (2π/h)2(1/(d2E/dk2)). Thus for a wide band material we expected a higher carrier velocity (and mobility) than for a narrow band material. Additionally we expect m* to increase for electrons moving within narrow bands, owing to the inverse relationship between curvature and effective mass. The m* relationship also explains why the concave downward band are called hole bands, i.e. these carriers would appear to have a negative effective mass, suggestive of a moving “hole”, while concave upward bands would have a positive effective mass like a typical electron.

Figure 1.2 Schematic depiction of electronic energy bands for a metal, semimetal, and insulator. The Fermi energy is marked as EF. 24

Figure 1.3 Electronic band structure and Fermi surface for LaFeAsO, calculated using crystallographic and calculation parameters as reported in Ref. [73].

The band structure and Fermi surface for LaFeAsO is shown in Figure 1.3.

Both hole like bands, centered around the gamma point, Γ, of the Brillouin zone as well as electron like bands, centered around M, are observed. It can be seen from

Figure 1.3 that the electronic structure of LaFeAsO reflects the anisotropic crystallographic bonding, and this is true for all LnMAsO compositions. In LaFeAsO the bonding within layers is stronger than bonding between layers, i.e. along the z- direction of the unit cell there is less overlap between orbitals. This is seen in the 25 band structure as the bands have little to no dispersion along the Γ to Z or X to R directions. The Fermi surface is therefore comprised of cylinders. The curvature of the Fermi surface being within the a-b plane says that the velocity of carriers moving within the a-b plane is greater.

One of the most interesting aspects of these materials is the Fermi surface nesting which can occur due to their two dimensional character. The electron-like sheet around M can be shifted by what is called a nesting vector to coincide with the hole-like cylinder around the Γ point. This sort of nesting generally indicates an electronic or magnetic instability in the compound and is what to the formation of the spin density wave (SDW) state in LaFeAsO [55]. doping LaFeAsO is thought to slightly change the topology of the Fermi surface and therefore diminish the nesting, resulting in superconductivity instead of the SDW state.

The electronic band structures for Co and Ni based compositions look similar to that of LaFeAsO; however, owing to their extra electrons the Fermi level is shifted upward and cuts through bands of slightly different d orbital character. For LaCoAsO this upward shift results in a very sharp peak in the density of states at the Fermi level, commonly called a Van Hove singularity [74]. In essence a Van Hove singularity indicates a possible electronic instability similar to Fermi surface nesting, but instead 26 of a SDW state forming spin orientation splitting of electronic bands is generally favored and an itinerant FM state can result, as it does for LaCoAsO [15].

The details of the electronic structure for LaRhAsO and LaIrAsO are given in

Chapter 3, but in general the 4d compositions show similar band structures to their

3d counterparts, however with increased band dispersion owing to their larger atomic orbitals. The increased dispersion means that 4d analogs should have an increased carrier velocity and therefore display and increase in electronic conductivity. The physical property changes upon substitution for Rh and Ir are explored further in Chapters 2 and 3.

27

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32

Chapter 2

Superconductivity within the LaFe1-xRhxAsO system

2.1 Abstract The LaFe1-xRhxAsO solid-solution has been characterized using powder x-ray diffraction, resistivity and magnetization measurements. Superconductivity onset is

onset observed for x = 0.05, 0.10, 0.15 with a maximum resistivity Tc of ~16 K for the x = 0.1 composition. From the observed changes in resistivity an electronic phase diagram for the LaFe1-xRhxAsO solid-solution is proposed. While the critical temperature for iron superconductors has been correlated to the As-Fe-As tetrahedral angle, doping Rh within LnFeAsO systems (where Ln = La, Nd, Sm) appears to limit the Tc to less than 20 K even when the tetrahedral angle is nearly ideal. 33

2.2 Introduction Iron pnictide superconductors have presented the scientific community with a new platform for studying high temperature superconductivity and as such a materials renaissance in superconductivity has begun. Following the discovery of superconductivity for LaFeAsO1-xFx with Tc up to 26 K, it was found that lanthanoid substitutions can increase the Tc to as high as ~55 K [1–6]. These LnFeAsO

(Ln = La - Nd, Sm, Gd) compositions are now often referred to as 1111 or Ln-1111 type materials. This class of materials is relatively new to the worldwide scientific community but in fact had been known and structurally investigated for some time

[7,8].

Since the rebirth of LnFeAsO materials as high temperature superconductors an entire family of iron arsenide superconductors has been uncovered which now includes 111-type AFeAs (A = Li, Na) [9,10], 122-type MFe2As2 (M = Ca, Sr, Ba) [11–

14], and 21322-type representatives such as Sr2V1O3FeAs [15]. Additionally, the iron chalcogenide compositions, FeCh (where Ch = Se, Te) have been shown to superconduct at reasonably high temperatures (8 - 15 K) [16,17] and more recently these have been intercalated with resulting in a much higher critical temperature of 30 K [18].

Upon cooling LaFeAsO exhibits a tetragonal to orthorhombic structural transition at ~150 K closely followed by the establishment of a spin density wave 34

(SDW) state [19–21]. A superconducting ground state can be induced by suppressing the structural and SDW transitions via doping or high pressure. Chemical substitution studies of Ln-1111 compositions which suppress the structural/SDW transitions or modify the Fermi level DOS are a key part of testing the underlying mechanisms of superconductivity and thereby searching for further increased critical temperatures.

Here we report on the hitherto unknown LaFe1-xRhxAsO system (where x =

0.05, 0.1, 0.15, 0.2, 0.5, 0.75) [22]. For LaFeAsO electron doping with 3d transitions metals such as Co and Ni has resulted in superconductivity, albeit with lower critical temperatures [23,24]. Hole doping La-1111 with Mn suppressed the SDW transition and resulted in a metal-insulator transition at about 3% Mn, but neither p-type conductivity nor superconductivity were seen [25]. Contrary to the effects of Mn doping into the [FeAs] layer, hole doping with Sr into the [LaO] layer produces superconductivity with a Tc of 25 K, near the same as for LaFeAsO1-xFx [26]. For

PrFeAsO a suppression of the SDW state with substitution of isoelectronic Ru has been reported, but superconductivity did not emerge [27]. Nearly simultaneously to our work with doping Rh into LaFeAsO (Ref. [22]) came other reports which used Rh to electron dope NdFeAsO and SmFeAsO [28,29]. Additionally, Ir has been used to induce superconductivity [29,30]. 35

The defining structural feature of Ln-1111 compounds, and others within the growing iron arsenide family, is layers of edge sharing As tetrahedra with Fe centers within the a-b plane. From a bonding standpoint the general formula, LnFeAsO, can be rewritten as [Ln3+O2-]1+ [FeAs]1- which helps to highlight the c-axis stacking and charge transfer between the predominantly ionic ‘charge reservoir’ layers *LnO] and the more covalent [FeAs] conduction layers. On the other hand, density functional studies of LaFeAsO show only modest Fe d-orbital and As p-orbital hybridization and it has been suggested that the structure be described as square sheets of metallic

Fe2+ in an ionic matrix of [La3+O2-As3-]2- stacked along the c-axis [31]. These two different structural depictions are shown in Figure 2.1. In terms of symmetry the structure is derived with the tetragonal space group P4/nmm (#129, setting 2) and atoms at four positions (La at ¼, ¼, zLa, Fe/Rh at ¾ , ¼, ½ As at ¼, ¼, zAs, and O at ¾,

¼, 0). 36

Figure 2.1 The structure of LaFeAsO and related 1111-type compositions where in (a) 3+ 2- 1+ 1- the structure is considered as [Ln O ] [FeAs] stacking along the c-axis and (b) the 3+ 2- 3- 2- 2+ structure considered as [La O As ] stacking with metallic Fe sheets.

2.3 Results and Discussion

2.3.1 Lattice parameters and unit cell volumes of LaFe1-xRhxAsO compositions

The observed powder x-ray diffraction patterns for LaFe1-xRhxAsO compositions indicate that a full solid solution forms between LaFeAsO and LaRhAsO, as shown in Figure 2.2. Secondary peaks from La2O3, La(OH)3, and FeAs are present 37 for some samples and for Rh rich compositions (x = 0.75, 1) trace amounts of RhAs are also detected. The a and c lattice parameters as well as unit cell volumes are shown plotted in Figure 2.3; values shown for LaFeAsO are from Ref. [1].

Figure 2.2 Powder x-ray diffraction (PXRD) patterns for LaFe1-xRhxAsO compositions.

38

Figure 2.3 Lattice parameters and unit cell volumes for the LaFe1-xRhxAsO solid solution; values shown for LaFeAsO are from Ref. [1].

As the Rh content rises the a-parameter steadily increases, while the c-parameter decreases. The cell parameters for LaRhAsO are included in Figure 2.3 for completeness however, this composition will be discussed in more detail in

Chapter 3 within the context of other LnRhAsO and LnIrAsO compositions (where

Ln = La, Ce, and Nd). Unit cell volumes show a slight decrease initially due to contraction of the c-parameter and the minimal expansion of the a-parameter. 39

However, as Rh content increases the overall cell volume expands again due to the steady increase of a. The strong c-axis compression upon replacing Fe with Rh can be understood by considering that Rh substitution will increase the electron density within the Fe layer and thereby strengthen the interlayer Coulomb attraction (and decrease the interlayer distance), despite the increase in M-As distance upon substituting a 4d metal into the structure.

2.3.2 Electronic properties of LaFe1-xRhxAsO compositions

Normalized resistivity, ρ/ρ300 K, values for the LaFe1-xRhxAsO series, shown in

Figure 2.4, indicate that the samples become more metallic with increasing Rh content. This behavior is to be expected simply based on the orbital diffusivity increase from 3d to 4d metals causing increased band dispersion. In other words, the larger orbital size of Rh should increase the Rh-Rh overlap within the a-b plane as well as the Rh d and As p hybridization and both of these effects will shift the system towards increased metallic character. However, for the low Rh content samples

(x<0.5) the temperature coefficient of resistivity is observed to change from positive

(metallic) to negative (semiconducting) upon cooling. For LaFeAsO the change in resistivity upon cooling to 160 K has been attributed to the spin density wave (SDW) formation in these materials [19,20]. 40

The resistivity plot clearly shows superconductivity for x = 0.05, 0.10 at about

10.5 and 16 K respectively; similar to other one electron donor systems (i.e. for Co, Ir, and F dopants). For the x = 0.075, 0.15, and 0.20 samples superconductivity onset is observed; however, the transitions are more broad and a full superconducting state

onset is not realized above 5 K as it is with x=0.05 and x=0.1 samples. The observed Tc values are close to those reported for Co doping x = 0.05 (11.2 K) and x = 0.11 (14.3

onset K), and higher than the Tc for Ir doping, x = 0.05 (8 K) and x = 0.10 (11 K) [23,30].

While the reported critical temperature for NdFe0.9Rh0.1AsO is about 17 K [28]. It is interesting to note that the critical temperatures for LaFe0.9Rh0.1AsO (16 K) and

NdFe0.9Rh0.1AsO (17 K) are almost equal; however, the Tc for NdFeAsO0.89F0.11 (52K) is twice that of LaFeAsO0.89F0.11 (26K). 41

Figure 2.4 Normalized resistivity for the LaFe1-xRhxAsO solid solution. Compositions stated are based on the nominal stoichiometry of the reaction.

To a first approximation, the temperature dependence of a metal can be thought of as coming from a linear combination of various independent scattering mechanisms. The scattering by impurities, ρ0, is not generally temperature dependent, but for instance scattering from electron-electron interactions gives a T2 dependence, whereas scattering from localized magnetic moments can give a ln(T) dependence. In the high temperature limit for metals (T much greater than the 42

Debye temperature, where all phonon modes become activated) scattering by phonons is expected to cause a linear relationship between resistivity and T [32,33].

Figure 2.5. The normalized resistivity of LaFe1-xRhxAsO samples (where x=0-0.2) plotted as function of T2 and inset showing the normalized resistivity plotted as a function of T.

For NdFe1-xRhxAsO compositions (where x=0.0 - 0.2) it was observed that the resistivity at high temperatures (>200 K) follows a T2 relationship, which can be indicative of a Fermi liquid state (i.e. a Fermi gas with electron-electron interactions) 43

[28]. Similarly, the resistivity of LaFe1-xRhxAsO compositions (where x=0.0-0.2) at temperatures above 200 K appear to follow a linear relationship when plotted as a function of T2 which could suggest that electron-electron interactions play a significant role, see Figure 2.5. This is by no means a concrete conclusion though as fitting from 250 - 300 K appears to show little difference between T and T2 plots and

Fermi-liquid behavior (electron-electron scattering) is generally considered to be a low temperature phenomenon which is over shadowed by phonon-electron scattering near room temperature [33].

An electronic phase diagram as a function of nominal composition and temperature can be constructed, as shown in Figure 2.6, by considering the various changes observed in the resistivity (i.e. the sign of the temperature coefficient, presence of a SDW anomaly or superconductivity, and possible T2 dependence) for

LaFe1-xRhxAsO samples across the solid solution. 44

Figure 2.6 An electronic phase diagram for the LaFe1-xRhxAsO solid solution as determined from changes in the resistivity data. Compositions are based on nominal reaction stoichiometry.

The presence of superconductivity for the LaFe0.9Rh0.1AsO sample is further supported by AC susceptibility and DC magnetization measurements, Figure 2.7, which becomes negative below ~14 K and ~11 K respectively. The in-phase component of the AC susceptibility measurements is plotted as an inset with the DC magnetization data. It is clear that the diamagnetic transition is not complete at 5 K nevertheless an estimate of the superconducting volume fraction can be estimated using the basic relationships for magnetic inductance, magnetization, and magnetic field (in cgs units), B = H + 4πM [32,33]. If we assume that the entire sample is perfectly superconducting then any magnetic field inside of the sample should be 45 completely expelled, in other words B = 0 and therefore 4πM/H = -1. Since M/H is simply the volume susceptibility, χvol, this says that for a perfectly superconducting sample a plot of 4πχvol should result in a value of -1, once the material becomes superconducting. The plot of 4πχvol for LaFe0.9Rh0.1AsO, shown in Figure 2.7, for zero field conditions reveals the superconducting volume fraction at 5K to be approximately 11%. Given that the magnetic transition is still not complete at 5 K

(i.e. χvol has not yet leveled out), the estimated volume fraction of 11% indicates bulk superconductivity for LaFe0.9Rh0.1AsO.

Figure 2.7 DC magnetization volume susceptibility (field cooled and zero field cooled conditions) for LaFe0.9Rh0.1AsO. Inset shows the magnetic transition as determined using AC susceptibility measurements. 46

Figure 2.8 Resistivity normalized to 30 K for LaFe0.9Rh0.1AsO plotted as a function of magnetic field and with an inset detailing the change in temperature for 10%, 50%, and 90% resistance after application of the magnetic field.

Figure 2.8 shows the resistivity temperature dependence, ρ/ρ30 K, as a function of magnetic field for the x = 0.10 sample exhibiting maximum Tc. From this

onset plot it can be seen that Tc is rather insensitive to the applied field strength while

offset Tc is more quickly shifted to lower temperatures. The Werthhamer-Helfand-

Hohenberg (WHH) model can be used to estimate the upper critical magnetic field at zero Kelvin as Hc2(0) = -0.693(Tc)(dHc2/dT)T=Tc [34]. Using the (dHc2/dT)T=Tc value for

90% normal state resistance the upper critical field is estimated to be 55 T. This 47 value is almost double that reported for SmFe0.9Rh0.1AsO (25.6 T) and one half that found for NdFe0.9Rh0.1AsO (100 T) [28,29]. Ginzburg-Landau theory is known to relate the superconducting Cooper pair coherence length, ξ, to the upper critical field

2 through the relationship, Hc2 = Φ0 / π ξ , where Φ0 represents a quantum of magnetic flux, or one fluxon, and equals approximately 2.07 x 10-7 Oe cm2. Rearranging this to solve for the coherence length and using the Hc2(0) value determined by using the

WHH model a value of about 24 Å is realized for LaFe0.9Rh0.1AsO. Interestingly this value is very similar to that reported for F doped samples of LaFeAsO [35,36] even though the electron donor in this case is also introducing disorder to the FeAs conduction layer, but it should be reminded that this rough estimate using polycrystalline samples does not account for structural anisotropy.

The composition LaFe0.9Rh0.1AsO is found to exhibit a negative Hall coefficient at room temperature and below, as shown in Figure 2.9 along with the data for

LaFeAsO and LaFeAsO0.89F0.11 estimated from Refs. [21,37]. To a first approximation we can assume single band conduction (i.e. one carrier type) for LaFeAsO and related compositions and that the Hall-coefficient is inversely proportional to the density and charge of the carrier. The large increase in the Hall-coefficient for LaFeAsO observed at about 160 K therefore indicates a decrease in electron carriers corresponding to the structural (tetragonal to orthorhombic) and magnetic phase transitions (SDW formation) which also occur at this temperature. Doping suppresses these 48 transitions and this is evidenced by the Hall-coefficient data for LaFeAsO0.89F0.11 and well as LaFe0.9Rh0.1AsO which do not exhibit a large change in magnitude upon cooling. Using the single band approximation the density of charge carriers at 300 K

21 -3 is estimated to 1.6 x 10 cm for LaFe0.9Rh0.1AsO which is in good agreement with the reported values for LaFeAsO and LaFeAsO0.89F0.11 [21,37,38].

Figure 2.9 Hall coefficient for LaFe0.9Rh0.1AsO as well as LaFeAsO and LaFeAsO0.89F0.11 as determined from Refs. [21] and [37].

49

If these were truly single band metals however it would be expected that the

Hall-coefficient is temperature independent and we can clearly see that even for

LaFe0.9Rh0.1AsO and LaFeAsO0.89F0.11 (with no distinct change at 160 K) there is still some temperature dependence. This can be readily explained however as LaFeAsO and related compositions are known to be multi-band conductors with Fermi surfaces containing both hole and electron sheets [31]. For materials with both electron and hole carriers the Hall-coefficient depends on the mobility of each carrier type in addition to their respective densities [33]. The variation with temperature therefore arises from differences in the temperature dependence of electron and hole mobility, or in other words the scattering rates for carriers within each of the different bands crossing the Fermi energy have different temperature dependence.

2.3.3 Structural refinement of LaFe0.9Rh0.1AsO

Structural refinement of powder x-ray diffraction data for LaFe0.9Rh0.1AsO was carried out in order to further correlate atomic structure with properties, such as Tc.

Previous studies of Ln-1111 type and 122 type iron arsenide materials have demonstrated a high correlation between the As-Fe-As bond angle and critical temperature for these superconducting materials [39,40]. The main results (lattice parameters, Fe-As bond distance, and As-Fe-As tetrahedral angle) for the structural refinement of LaFe0.9Rh0.1AsO are shown in Figure 2.10; additional phases La2O3 and 50

FeAs were also included as part of the refinement and determined to amount to approximately 7 % by weight.

Figure 2.10 Rietveld refinement results LaFe0.9Rh0.1AsO detailing the average As-Fe- As angle found and its change relative to LaFeAsO, as reported in Ref. [1]. The tick marks from top to bottom represent the (hkl) reflections for LaFe0.9Rh0.1AsO, La2O3, and FeAs phases respectively.

51

Upon Rh insertion into the lattice of LaFeAsO it is found that the average

As-Fe-As angle moves farther away from the ideal tetrahedral value of 109.5°. This change in As-Fe-As angles is primarily due to the c-axis compression which occurs upon Rh substitution, as the a-lattice parameter and average pnictide height within the unit cell do not change much for LaFe0.9Rh0.1AsO (a = 4.0377(1) Å, zAs = 0.6510(1)) when compared to LaFeAsO (a = 4.0355 Å, zAs = 0.6512) from Ref. [1]. The As-Fe-As bond angles, α, for Ln-1111 type materials, adapted from Ref. [39], are shown plotted in Figure 2.11 along with the bond angle determined for LaFe0.9Rh0.1AsO as determined in this work and the bond angles for Nd and Sm analogs as determined from Refs. [28,29].

According to a survey of iron arsenide superconductors the Fe-As bond distance ranges from 2.39-2.41 Å, a seemingly small range for the relative difference in lattice parameters between large and small lanthanoid cations. More, specifically it has been reported that As-Fe-As angles determined by using lattice parameters and a fixed Fe-As bond length of 2.40 Å (independent of the lanthanoid cation present) yields angles with an error of approximately 1° [39]. It is by this means that the bond angles shown in Figure 3.11 for SmFe0.9Rh0.1AsO were determined, although only lattice parameters are explicitly reported in Ref. [29].

52

Figure 2.11 Critical temperature as a function of tetrahedral angle for Ln-1111 type materials, adapted from Ref. [39], as well as those determined for LaFe0.9Rh0.1AsO, NdFe0.9Rh0.1AsO (Ref. [28]), and SmFe0.9Rh0.1AsO (Ref. [29]).

It is immediately clear from Figure 2.11 that the critical temperatures for optimally doped Rh samples are all nearly the same and significantly less than those of their F-doped or oxygen deficient counterparts. It is also clear that the critical temperature of Rh doped samples fall well below the values expected based simply on tetrahedral angles. All systems demonstrated highest Tc at x = 0.1 and this observation suggests that while a threshold amount of electron doping is still needed 53 to achieve the maximum possible Tc it is the disorder imparted to the Fe layer through Rh doping which is limiting the Tc to less than 20 K.

Refined values for Fe-As, La-As, and La-O bond distances can also be used to calculate the expected cation valency through the use of tabulated bond valences

[39,41,42]. Tabulated bond valence values are determined from ionic compounds where the valence of each element is clearly established. If ionic bonding is in play, then their experimental bond valence sums (BVS) should match well with those expected from a purely ionic bonding standpoint. In the case of LaFeAsO these would be Fe2+ and La3+.

The bond distances found for LaFe0.9Rh0.1AsO are Fe-As = 2.4100(16) Å, La-As

= 3.3797(14) Å, and La-O = 2.3661(8) Å. has four bonds to oxygen which contribute 2.37 valence units (v.u.) and another four bonds to which contribute 0.83 v.u. to give a total La BVS of 3.20 v.u.. This value is in relatively good agreement with that expected from an ionic standpoint but the partial bond valences

(and bond distances) indicate that La is primarily bonded to oxygen, or in other words that the structural bonding is anisotropic. Iron has four bonds to arsenic which each contribute 0.85 v.u. to give a total BVS of 3.40 v.u., which is substantially above the expected value if purely Fe2+-As3- bonding were present. Bond lengths of approximately 2.60 Å would be needed to give a BVS of about 2.0 for Fe/Rh in 54

LaFe0.9Rh0.1AsO and therefore the Fe-As experimental bond lengths are too short for them to be completely ionic in nature. The Fe-Fe bond length of LaFe0.9Rh0.1AsO is determined to be 2.8551(4) Å. As it was pointed out in the first report of LaFeAsO

[8], and later supported through DFT calculations [31], this bond length is short enough that there is likely some Fe-Fe bonding at play in these materials and thus it is the interplay between Fe-As and Fe-Fe interactions which determines the observed bond lengths and angles for these materials.

2.4 Conclusion The synthesis of LaRhAsO as well as the solid solution between LaFeAsO and

LaRhAsO has been reported. This work confirms that doping with the 4d metal Rh is an effective means for inducing superconductivity within La-1111; similar to 3d Co and 5d Ir doping. A maximum critical temperature of 16 K is observed for

LaFe0.9Rh0.1AsO by resistivity measurements while magnetization and AC susceptibility data for the x = 0.10 sample show transitions at slightly lower

onset temperatures. The maximum Tc observed is close to what has been reported for

3d Co doping and for Rh doping within other LnFe1-xRhxAsO systems (where Ln = La,

Nd, Sm) but lower than the maximum found within the LaFeAsO1-xFx system. The critical temperatures of Rh doped samples cannot be explained through As-Fe-As bond angle considerations alone and doping Rh within the FeAs layer appears to limit 55 critical temperatures to less than 20 K. Furthermore, it has been shown that the metallic character increases with the addition of more Rh and that LaRhAsO remains metallic down to 5 K. The structures and properties of LnRhAsO and other LnIrAsO materials (where Ln = La, Ce, and Nd) will be discussed further in Chapter 3.

2.5 Materials and Methods Polycrystalline samples of LaFe1-xRhxAsO (x = 0.025, 0.05, 0.075, 0.10, 0.15,

0.20, 0.50, 0.75, 1.0) were prepared from La2O3 (Alfa Aesar 4N), La (Alfa Aesar 3N), Fe

(Sigma Aldrich 3N), Rh, and As (Alfa Aesar 2N). The Rh metal was prepared in two steps by oxidizing the metal and then reducing it under flowing 5% H2 / 95%

N2. The La2O3 was dried at 850°C for ~12 hrs before use. After weighing the reagents were ground ~15 minutes with agate mortar and pestle and pressed into pellets. The pellets were sealed inside fused silica tubes after evacuating to ~1 x 10-4 torr. All steps of the sample preparation, except annealing, were carried out in a high purity glove box. Samples were heated 400°C - 12 hrs, 850°C - 12 hrs, and 1090°C -

48 hrs with ramp rates of 300°C/hr between stages.

Lattice parameters were determined from powder x-ray diffraction data

(PXRD) obtained over 10 - 60° 2θ with a step size of 0.01° 2θ and Cu Kα radiation

(Rigaku MiniFlex II) with a graphite monochromator on the diffracted beam. A Le Bail 56 profile fitting method within the FullProf suite was utilized to refine the lattice parameters. GSAS and EXPGUI were used for Rietveld structural refinement of

LaFe0.9Rh0.1AsO over the range of 20 - 95° 2θ. Four-probe resistivity measurements were performed over a temperature range of 5 - 300 K. For the x = 0.10 sample resistivity as a function of magnetic field from 0 - 6T was examined in the 2-30 K transition region. Zero field and field cooled DC magnetization measurements were done on warming from 5 - 300 K in a 20 Oe field. Zero field cooled AC susceptibility measurements utilized a 750 Hz 10 Oe signal. Hall-coefficients were determined from the change in Hall resistance under magnetic fields from 0 - 6 Telsa using a

5 -probe wiring configuration. Electronic and magnetic measurements were both performed using a Quantum Design Physical Properties Measurement System

(PPMS).

57

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60

Chapter 3

Synthesis and Electronic Properties of LnRhAsO and LnIrAsO (Ln = La, Ce, Nd) Compositions

3.1 Abstract The synthesis and physical properties of LnRhAsO and LnIrAsO (where Ln = La,

Ce, and Nd) compositions are presented along with the electronic band structures of

LaRhAsO and LaIrAsO. The Ce based compositions demonstrate properties very similar to other magnetic Kondo materials and indicate that there is a cross over from

Kondo to Ruderman-Kittel-Kasuya-Yosida type interactions occurring in these materials as they are cooled. The resistivity curves for Ce compositions are nonlinear and display a sudden reduction in resistivity which coincides with AFM ordering. The

max Kondo temperatures, TK , observed through resistivity measurements and the magnetic ordering temperatures can be explained by considering that CeRhAsO should have an increased N(EF)J parameter relative to CeIrAsO. Electronic band structure calculations support this conclusion in that LaRhAsO has a larger N(EF) than

LaIrAsO and this difference, driven by the shorter a-parameter of Ir based compositions, should also ensure that CeRhAsO has a larger N(EF) than CeIrAsO.

61

3.2 Introduction Since the discovery of superconductivity for doped LaFeAsO compositions there has been a major resurgence of interest in layered oxypnictides which crystallize in the ZrCuSiAs structural archetype. Compared to first row transition metals compositions such as LnMPnO (where M = late 3d transition metal, including group 12) much less is known about the structure and properties of 4d and 5d transition metal based compositions. Figure 3.1 depicts the general structure of

LnMPnO materials.

Figure 3.1 The structure of LnMPnO compositions portrayed as a stack of alternating MPn layers (antifluorite like) and LnO layers (fluorite like) along the c-axis. 62

The oxyphosphide compositions LnRuPO, LnOsPO, ThAgPO, and LaCdPO

(where Ln = La - Nd, Sm), have previously been structurally investigated [1–3].

However, of these compositions only LnRuPO and LnOsPO (where Ln = La, Ce) have had their physical properties discussed within the literature. Interestingly CeRuPO is reported to develop a ferromagnetic (FM) Kondo lattice ground state at low temperature while CeOsPO adopts an anti-ferromagnetic (AFM) state [4,5].

Looking next at the oxyarsenides the only 4d metal examples known previous to the Rh based compositions discussed here were LnRuAsO (where Ln = La - Nd, Sm,

Gd-Dy) and LnCdAsO (where Ln = La - Nd). Previous to the Ir compositions discussed here there had been no fully substituted 5d metal containing oxyarsenides reported

[6,7].

Following the LnFeAsO discoveries, resistivity data for LnRuAsO (where

Ln = La, Ce) compositions have been reported and metallic behavior was observed down to 4 K even when fluorine doped [8]; however it was not until very recently that any additional physical properties, such as magnetic susceptibility and Hall coefficient were reported [9]. Aside from the structural properties originally reported for LnCdAsO compositions [2] there have been no reported investigations of the physical properties for Cd analogs. Previously we have reported on the solid solution LaFe1-xRhxAsO and demonstrated that Rh is an effective dopant for inducing 63 superconductivity in LaFeAsO (max. Tc ≈ 16 K for x = 0.1) as discussed in Chapter 2

[10]. In fact, some iron free pnictides, isostructural to BaFe2As2, which contain Rh and Ir are reported to become superconducting albeit at very low temperatures:

BaRh2P2 (Tc ≈ 1 K), BaIr2P2 (Tc ≈ 2.1 K), SrIr2As2 (Tc ≈ 2.9 K) [11,12]. Here we report on the 4d substituted compositions LnRhAsO, as well as the first 5d metal oxyarsenide compositions, LnIrAsO (where Ln = La, Ce, Nd) [10,13].

3.3 Results and Discussion 3.3.1 Crystal structure of LnRhAsO and LnIrAsO materials The primary phase in powder x-ray diffraction (PXRD) patterns for LnRhAsO as well as LnIrAsO compositions could be indexed as P4/nmm (#129); Figures 3.2 and

3.3 show the PXRD patterns for each series respectively. The lattice parameters for the LnRhAsO and LnIrAsO series, as determined through least squares refinement, are shown plotted in Figure 3.4. It is readily seen that there is a systematic decrease in a and c lattice parameters as smaller cations are substituted. Upon substitution of Rh for Co the a - parameter increases while the c-parameter remains relatively unchanged; however, upon substitution of Ir for Rh the a-parameter contracts while the c-parameter expands. Lattice parameters and standard deviations are shown listed in Table 3.1; for comparison the lattice parameters for Co compositions are also listed [14]. 64

Figure 3.2 Powder x-ray diffraction patterns for LnRhAsO compositions where Ln = La, Ce, Nd. Open diamonds (◊) indicate the presence of RhAs while closed diamonds () represent CeO2-x phases.

Figure 3.3 Powder x-ray diffraction patterns for LnIrAsO compositions where Ln = La, Ce, Nd. Open diamonds (◊) indicate the presence of IrAs2 while closed diamonds () represent Ir. 65

It is interesting that the a-parameters observed for LnIrAsO compositions are shorter than the corresponding LnRhAsO compositions. This is somewhat contrary to what might be expected from simple 4d and 5d orbital size considerations, for while it is well known that the lanthanoid contraction causes 5d metals to be only slightly larger than their 4d counterparts there is still generally an increase in lattice parameters when substituting Ir for Rh in solid state compositions. As it was discussed by Quebe et al., using bond lengths and simple electron counting for

LnFeAsO materials (and later evidenced through band structure calculations of

LaFeAsO) the distance between M cations in LnMAsO materials is relatively short and there are likely some M-M orbital interactions at play [14,15].

For these materials the M-M bond distance is simply related to the a-parameter via the relation, M-M = a / √2 and a decrease in the a-parameter therefore means the Ir-Ir bond is shorter than the Rh-Rh bond formed in these materials. The crystalline structure for LnRhAsO and LnIrAsO materials is shown in

Figure 3.5a; however the LnAs bonding is shown, rather than depicting the LnO and

MAs layers as distinct. Figure 3.5b highlights the MAs slab as viewed along the c-axis and demonstrates the relationship between the a-parameter and the M-M bonding distance, shown as a dashed line. 66

Figure 3.4 Lattice parameters for LnRhAsO and LnIrAsO compositions.

Table 3.1 Lattice parameters for LnCoAsO, LnRhAsO, and LnIrAsO compositions as determined by x-ray powder diffraction. Group 9 La Ce Nd Reference metal a (Å) c (Å) a (Å) c (Å) a (Å) c (Å) Co 4.054(1) 8.472(3) 4.015(1) 8.364(2) 3.982(1) 8.317(4) [14] Rh 4.123(1) 8.469(1) 4.089(1) 8.365(2) 4.063(1) 8.303(1) This work Ir 4.087(1) 8.557(1) 4.063(1) 8.444(2) 4.046(1 8.350(1) This work

67

Figure 3.5 The structure of LnRhAsO and LnIrAsO materials highlighting the LnAs bond distances (a) and the MAs layer as viewed along the c-axis (b), detailing the M-M distance and its relationship to the unit cell a- parameter.

Unit cell volumes, as well as M-M bond distances, for LnRhAsO and LnIrAsO compositions are shown plotted in Figure 3.6. The decrease in cell volume upon substitution for smaller rare earth elements follows a similar trend as the Fe, Co, and

Ru analogs previously reported [14] wherein the smaller cell volume observed for

CeCoAsO was rationalized by assuming some Ce to be in a mixed valent state, Ce3+/4+.

68

Figure 3.6 Unit cell volumes and metal-metal (M-M) distances for LnRhAsO and LnIrAsO compositions.

The same rationale might seem to apply to CeRhAsO and CeIrAsO, however further studies will be need to definitively prove the existence of mixed valent Ce3+/4+ and in the case of CeFeAsO and CeNiAsO X-ray photoelectron spectroscopy (XPS) and

X-ray absorption near-edge spectroscopy (XANES) have suggested that only Ce3+ is present [16].

69

Table 3.2 Primary results of powder x-ray diffraction Rietveld refinement for LnRhAsO and LnIrAsO compositions. LaRhAsO NdRhAsO LaIrAsO NdIrAsO P4/nmm (#129), a = b ≠ c, α = β = γ = 90° Ln (¼, ¼, zLn), M (¾ , ¼, ½), As (¼, ¼, zAs), and O (¾, ¼, 0) a - parameter, Å 4.12296(6) 4.06280(6) 4.08832(14) 4.05000(22) c - parameter, Å 8.46521(18) 8.29934(21) 8.56476(35) 8.3576(5) Unit cell volume, Å3 143.898(6) 136.992(4) 143.154(9) 137.084(23) Ln position, zLn 0.14121(11) 0.13819(16) 0.14240(18) 0.13842(17) As position, zAs 0.65439(22) 0.66304(31) 0.6604(4) 0.66793(34) Global Uiso 0.00375(19) 0.00164(6) 0.0014(4) 0.0061(4) M-M distance, Å 2.91537(4) 2.87283(4) 2.8909(1) 2.86377(16) M-As distance, Å 2.4409(10) 2.4408(14) 2.4630(18) 2.4638(16) Bond Valence, 0.826(1) 0.826(1) 1.23(1) 1.23(1) BVM-As As-M-As angle, ° 115.25(7) 112.66(10) 112.19(12) 110.55(11) Ln-As distance, Å 3.3902(9) 3.3128(13) 3.3480(17) 3.2894(15) Bond Valence, 0.203(1) 0.200(1) 0.227(1) 0.213(1) BVLn-As Ln-O distance, Å 2.3830(5) 2.3328(7) 2.3803(8) 2.3322(7) Bond Valence, 0.565(1) 0.648(1) 0.475(1) 0.541(1) BVLn-O Additional phases La2O3, RhAs Nd2O3 La2O3, IrAs2, Ir Nd2O3, IrAs2, included in fit Ir Estimated wt. % 95.9 98.2 93.8 94.6 of primary phase Rwp 15.27 14.07 12.81 12.77 χ2 2.505 2.451 3.357 3.679

70

Presumably the decrease in the M-M bond length is accompanied by an increase in the c-parameter as pnictide anions move toward the rare earth atoms, expanding the width of the M2Pn2 layer. To further investigate this structural change

Rietveld refinements were carried out for La and Nd based compositions and the primary results are shown in Table 3.2.

From the tabulated Rietveld refinement results it is clear that the average

Rh-As and Ir-As bond distances hardly change as the La is substituted for a smaller lanthanoid, Nd. This insensitivity of the M-As bond length is consistent with the results found for LnFeAsO compositions as well as Ru based compositions [9]. As the

M-M bond distances gets smaller the zAs position to shift toward higher values in order to maintain the ideal M-As bond distance. While the M-M bond distance may be shorter for Ir based materials the refinement results indicate that the Rh-As and

Ir-As bond distances follow the expected trend; namely that the Rh-As bond should be shorter than the Ir-As bond, because Ir is in fact larger than Rh.

A bond valence sum (BVS) analysis of the LnRhAsO and LnIrAsO compositions can be carried out using the experimental bond lengths along with tabulated values for M-As, Ln-As, and Ln-O bonds [17,18]; the calculated bond valences (BV) are shown in Table 3.2. As the M cations have four bonds to As anions the BVS is simply,

BVSM = 4 (BVM-As); for the Ln cations BVSLn = 4 (BVLn-As) + 4 (BVLn-O) owing to their four 71

As and four O bonds. The BVS analysis for these compositions reveals that the bond valences for Ln cations sum to approximately 3, as would be expected based on their

3+ Ln oxidation state; however, if the individual sums, BVSLn-As and BVSLn-O, are evaluated it is clear that the Ln are more strongly bonded to the O anions as BVSLn-O accounts for about 66% of the total BVSLn in all cases. This result clearly highlights the anisotropic bonding of these layered materials.

On the other hand, BVSM in all cases is much larger than the formal oxidation state of M2+ expected from a purely ionic standpoint. This indicates that the M-As bonds in these materials are much shorter than expected based on an ionic bonding scenario; specifically the bond distance for M-Rh would need to be about 2.62 Å, rather than 2.44 Å, and the distance M-Ir would need to be about 2.80 Å, rather than

2.46 Å to give a BVSM of approximately 2.0 in each case. The BVSM value obtained for

NdRhAsO (3.30) is nearly equal to that found for NdCoAsO (3.29), as mentioned in

Chapter 1, but the values BVSM obtained for NdIrAsO is significantly greater which seems to support the idea that M-M bonding is greater in the Ir compositions

The near constant M-As bond distance between La and Nd based compositions can actually be exploited along with the experimentally determined a-parameter to determine the As-M-As angles and zAs positions in CeRhAsO and

CeIrAsO compositions. The geometric relationships which allow these to be 72 determined are shown in Figure 3.7. The As-M-As angles, α and β, are equated by the equation shown only one of the tetrahedral angles must be defined as has previously been discussed in the literature [19].

Using these relationships we can then plot the As-M-As angles, α, and the As atomic position within the unit cell, zAs, for La, Ce, and Nd compositions as shown in

Figure 3.8. It is readily seen that the As atoms sit at a higher zAs position within the unit cell for LnIrAsO compositions when compared to their Rh counterparts which translates to a thicker MAs layer. It can also be seen from Figure 3.8 that the

As-M-As bond angles, α, for LnRhAsO and LnIrAsO compositions decrease as smaller lanthanoids are substituted into the structure and that in general the bond angles for

LnIrAsO compositions are close to the ideal value of 109.5°.

Figure 3.7 The relationships between M-As bond distance, a-parameter, As atomic position (zAs), and tetrahedral angles for LnMAsO materials. 73

Figure 3.8 Tetrahedral bond angle, As-M-As, and As atomic position, zAs, evolution as a function of lanthanoid element for LnRhAsO and LnIrAsO compositions.

3.3.2 Physical properties of LnRhAsO and LnIrAsO materials Moving next to look at the resistivity of LnRhAsO and LnIrAsO materials as a function of temperature we can see that all compositions are metallic, as shown in

Figures 3.9a and 3.9b respectively. 74

Figure 3.9 Normalized resistivity, ρ/ρ300K, for LnRhAsO compositions (a) and LnIrAsO compositions (b).

Immediately clear are the precipitous drops in resistivity beginning at ~ 7 K for

CeRhAsO and ~10 K for CeIrAsO. Another noteworthy observation is that the Ce based compositions display a generally bowed shape to their resistivity plots. This type of behavior as been observed before for other CeMAsO compositions such as,

CeNiAsO (transition at 9 K) and its normalized resistivity [8,20] is shown plotted with the data for CeRhAsO and CeIrAsO compositions in Figure 3.10.

The behavior of CeRhAsO, CeIrAsO, and CeNiAsO are all quite similar to that of other compositions within the general composition of CeM2Si2 (where M = Pd, Ag,

Au) and which contain the same antifluorite like transition metal layer [21–25].

These compositions have been described as a reduced moment antiferromagnets on 75 the border between Ruderman-Kittel-Kasuya-Yosida (RKKY) and Kondo lattice ground states [24].

In these materials the AFM ordering temperature observed is thus determined by two competing energy scales; on being dictated by the temperature at which an RKKY state forms, TRKKY, and the other being determined by the temperature at which the Kondo state is favored, TK.

Figure 3.10 Normalized resistivity, ρ/ρ300K, for the Ce based compositions CeRhAsO and CeIrAsO, as well as CeNiAsO as reported in Ref. [8].

76

The Kondo state is one in which scattering between the 4f moments and conduction electrons causes the 4f electrons to effectively undergo a spin flip interaction such that they align antiparallel to the conduction electrons [26–28].

When the Kondo scattering is incoherent (typically at higher temperatures) this interaction causes the temperature coefficient of resistance (TCR) to become negative (i.e. semiconducting) although the conduction electrons are still delocalized throughout the material; coherent Kondo scattering however still leads to a positive

(TCR) typical of metals. Additionally because the Ce electrons become aligned antiparallel to the conduction electrons the magnetic moment per Ce ion is effectively cancelled (or greatly reduced) in the Kondo state.

In the RKKY state 4f electrons and conducting electrons still interact, however in this case long range magnetic ordering arises because the localized 4f electrons partially polarize the conduction electrons [27,28]. As conduction electrons are delocalized through the material, the alignment of one Ce ion therefore influences the alignment of the other Ce ions and the result is long range magnetic ordering of

4f moments.

The competition between the RKKY and Kondo states has been well studied in the literature [24,25,29–32] and it has been suggested that a critical parameter exists, N(EF)J, below which the RKKY interaction dominates and above which the 77

Kondo singlet state dominates (where J represent the exchange parameter between the 4f electrons and d electrons and N(Ef) is the density of states for conducting electrons). The transition between the two states is in fact continuous and thus there are materials which lie in the cross over region where the Kondo and RKKY temperature scales are nearly equal and in which both states can be observed depending on the sample temperature. A qualitative description of this cross over from magnetic metal to Kondo state is shown in Figure 3.11, similar to those shown in Refs. [30] and [33].

Figure 3.11 A qualitative depiction of the Kondo lattice, TK, and RKKY temperature scales, TRKKY, as function of the parameter N(EF)J. Within the magnetic Kondo state (Mag. KS) the magnetic ordering temperature is labeled TM.

78

On the left hand side of Figure 3.11 the RKKY type interaction dominates and the magnetic ordering is observed to increase with increasing N(EF)J. Within the magnetic Kondo state (Mag. KS) the Kondo type interaction now dominates at high temperatures. However, as the sample is cooled further the RKKY interactions take over and magnetic ordering is eventually observed, although now at temperatures below the RKKY prediction; indicated by TM rather than TRKKY. At sufficiently high enough N(EF)J the RKKY interactions become further suppressed relative to the

Kondo interactions and magnetic ordering is no longer observed.

Generally speaking the strength of J depends on unit cell volume effects (the extent of orbital overlap in a material) as well as electronic effects (the density of states at the Fermi level N(EF)) [25]. Looking at the crystal structure parameters of

CeRhAsO vs CeIrAsO it can be seen that the c/a ratio is shorter for CeRhAsO (2.045) than for CeIrAsO (2.078) and this suggests that CeRhAsO might have a greater exchange constant, J, than CeIrAsO due to this decrease in Ce-M distance, even though their cell volumes are nearly equivalent. CeIrAsO would also be expected to have a decreased N(EF) relative to CeRhAsO simply due to the more diffuse 5d orbitals which are involved in bonding and which dominate the Fermi level. Both of these factors should ensure that CeRhAsO has a larger N(EF)J parameter than

CeIrAsO, thus qualitatively placing it to the right of CeIrAsO on the N(EF)J plot shown in Figure 3.11. 79

The nonlinear temperature dependence of resistivity exhibited by CeRhAsO and CeRhAsO at high temperatures is a hallmark of Kondo lattice materials and has been found in both 3d metal arsenides, such as CeFeAsO [34], and CeNiAsO [8,20], as well as phosphides like CeFePO [34], CeRuPO [4], and Ce3Cu4P4O2 [35]. To better understand the cause of this knee in the resistivity curves it is instructive to look at what is commonly called the 4f resistivity, ρ4f (or the magnetic resistivity, ρM). The 4f resistivity is simply the resistivity of the material having first subtracted any contributions due to lattice vibrations (phonons), i.e. ρ4f = ρNormal - ρPhonon. This can be readily done for CeRhAsO (or CeIrAsO) by assuming that the resistivity of LaRhAsO (or

LaIrAsO) will embody all the phonon contributions. As the only major difference between the two compositions is the addition of the 4f electron on the lanthanoid site we can then say, ρ4f = ρCeRhAsO - ρLaRhAsO. Figure 3.12 details the low temperature

4f resistivity (normalized to 200 K) for CeRhAsO and CeIrAsO compositions. 80

Figure 3.12 The 4f resistivity for CeRhAsO and CeIrAsO compositions as well as the max observed magnetic ordering temperatures, TN, and Kondo lattice temperature,TK .

If the Ce ions were fully isolated (free ions) their 4f orbitals would normally be degenerate; however within a solid this degeneracy is lifted due to the potential energy field caused by the other atoms, or what are simply called crystalline field (CF)

max effects. At high temperatures, T > TK , the negative 4f - TCR observed for CeRhAsO and CeIrAsO can be attributed to incoherent Kondo scattering of conducting electrons from many different thermally populated CF levels, i.e. excited Ce states, 81

min while the low temperature negative TCR region, T < TK , can be attributed to scattering from just the ground state CF levels of Ce [25, 36]. Figure 3.12

max max max demonstrates that TK is higher for CeRhAsO (TK ~ 99 K) than it is for CeIrAsO (TK

~ 72 K), thus indicating that Kondo lattice interactions are stronger in the former.

A detailed study of CeNiAsO has demonstrated that the low temperature drop in resistivity coincides with a small antiferromagnetic transition [20]; thus it is reduced carrier scattering upon magnetic ordering which causes this precipitous drop. Figure 3.13 and Figure 3.14 overlay the AC magnetic molar susceptibility for

CeRhAsO and CeIrAsO with the low temperature resistivity data.

Figure 3.13 Low temperature AC molar magnetic susceptibility and resistivity plots for CeRhAsO. 82

Figure 3.14 Low temperature AC molar magnetic susceptibility and resistivity plots for CeIrAsO.

These plots reveal that the low temperature changes in resistivity for

CeRhAsO and CeIrAsO are likewise related to an antiferromagnetic (AFM) transition and are not due to the presence of a small fraction of superconducting material; otherwise a negative magnetic susceptibility would be observed.

DC magnetic susceptibility measurements were also carried on both LnRhAsO and LnIrAsO compositions to better understand the low temperature transition as well as the long range paramagnetic interactions for these materials; see Figure 3.15 and Figure 3.16. The low temperature DC magnetic susceptibility for CeRhAsO clearly reveals an AFM transition at about 6.4 K, which agrees well with the TN observed by 83

AC susceptibility as well as resistivity measurements. The susceptibility transition observed for CeIrAsO is much less pronounced, similar to the AC magnetic susceptibility results collected at lower field strength, but nevertheless a small bump is observed centered at approximately 9.6 K. Given that these are layered materials and generally exhibit anisotropic magnetization, the weak signal for CeIrAsO may be simply the consequence of sample preferred orientation.

Figure 3.15 DC magnetic susceptibility of LnRhAsO compositions. Inset is a table listing the results of a Curie-Weiss fitting, 휒 푚표푙 = 퐶/(푇 − 휃), for Ce and Nd based compositions. 84

Figure 3.16 DC magnetic susceptibility of LnIrAsO compositions. Inset is a table listing the results of a Curie-Weiss fitting, 휒 푚표푙 = 퐶/(푇 − 휃), for Ce and Nd based compositions.

The DC magnetic susceptibility data from 175 - 300 K were initially fit using the modified Curie-Weiss model 휒 푚표푙 = 퐶/(푇 − 휃) + 휒 0, i.e. including a temperature independent term. In all cases except NdRhAsO the temperature independent terms,휒 0, proved unstable in that their refined standard deviations were larger than

휒 0 itself; therefore 휒 0 was subsequently fixed to be zero while refining the Curie and

Weiss constants. Refinements of LaRhAsO and LaIrAsO were not possible given their weak magnetization, which also demonstrates that the Rh and Ir ions contribute no 85 appreciable magnetic moment. It therefore makes sense that the observed Curie constants, C, and calculated effective moments, μEff, agree well with lanthanoid free ion values [37]; which are shown tabulated along with the experimental μEff values in

Figure 3.15 and Figure 3.16.

The negative Weiss constants seen for all compositions indicate long range

AFM exchange interactions between Ln moments within the paramagnetic regime.

NdRhAsO and NdIrAsO do not order above 2 K and this is similar to what was very recently reported for NdRuAsO, although an upturn in heat capacity at 4 K was observed for this composition indicating that Nd ions may order just below 2 K [9].

SmRuAsO and GdRuAsO exhibit AFM ordering at 4.5 K and 5 K respectively and show a corresponding drop in resistivity, but with a near linear resistivity-temperature relationship (i.e. no sign of Kondo scattering) [9]. 86

Figure 3.17 The positions of CeRhAsO, CeIrAsO, and CeNiAsO within the N(EF)J plot max based on their observed TK and TN values. Values for CeNiAsO were taken from Ref. [20].

Thinking back to Figure 3.11 we can now begin to understand the differences in physical properties for CeRhAsO and CeIrAsO. The increase in ordering temperature for CeIrAsO, lack of a second upturn in resistivity before AFM onset, and

max decrease in TK all indicate that the Kondo lattice state is not as energetically favored in CeIrAsO. As stated before, CeRhAsO should have a larger N(EF) and

max possibly J larger in comparison to CeIrAsO. The differences in TK and TN for these compositions suggest they must lie within the magnetic Kondo state (Mag. KS) 87 region; as shown in Figure 3.17, which now places these compositions qualitatively

max along the N(EF)J plot. The composition CeNiAsO is was reported to have a TK of approximately 92 K and a TN of approximately 9 K [20] and these values suggest that it falls in between the CeRhAsO and CeIrAsO along the N(EF)J plot.

Before moving on to talk about electronic structure it should be mentioned that Hall coefficient measurements indicate that hole type conduction dominates from 3 - 300 K. The Hall Effect for LaRhAsO is larger in magnitude and has stronger temperature dependence than for CeRhAsO. The temperature variation of Hall coefficient is understood to come from multi-band effects, i.e. different temperature dependence of scattering for holes and electrons and as will be discussed below

LaRhAsO is in fact a multi-band carrier. No clear evidence of AFM ordering is observed at low temperature for CeRhAsO; however Hall measurements performed under static magnetic field conditions are generally noisy due to a relatively small hall voltage signal and perhaps more careful measurements near this transition will reveal its presence. The Hall coefficient sign CeRhAsO agrees with room temperature

-1 thermopower measurements for CeRhAsO (α298 K = + 1.43 μV K ) and CeIrAsO

-1 (α298 K = + 7.49 μV K ). 88

Figure 3.18 The Hall coefficient temperature variation of LaRhAsO and CeRhAsO.

3.3.3 Electronic structure of LaRhAsO and LaIrAsO To further investigate the differences between Rh and Ir based compositions density functional theory (DFT) electronic band structure calculations were carried out for LaRhAsO and LaIrAsO using the Wien2K software package [38]. Calculations were done using experimentally reported lattice parameters and atomic positions, as well as with experimental lattice parameters and energy minimized internal parameters, similar to what has been done for LaFeAsO, CeFeAsO, CeFePO and 89

LaNiPO [15,32,39]. Specific details of the calculations are given in the Materials and

Methods section at the end of this chapter; however it should be mentioned here that the choice of local coordinates for the Rh (Ir) atoms was such that the x and y axis are directed at the near neighbor Rh (Ir) atoms and thus it is the 푑푥2−푦 2 orbitals which are responsible for M-M bonding. The band structure and density of states,

N(E), as results are shown in Figure 3.19(a-c) and Figure 3.20(a-c), where the black line represents is the band structure determined using energy minimized internal parameters, zAs and zLa, and the red dashed line represents that found using experimental zAs and zLa values. Additionally the Brillouin zone and its high symmetry points are shown overlaid in Figures 3.19(a) and 3.20(a).

The anisotropy of bonding in LaRhAsO and LaIrAsO is reflected in their band structures in that bands running along the z direction of the Brillouin zone, e.g. Γ to Z, show less dispersion than those moving along the x and y directions, e.g. Γ to X. The partial atomic density of states for each compound indicate that Rh(Ir) contribute the most to the DOS at the Fermi level. Experimental and energy minimized atomic positions for LaRhAsO match very well with one another; experimental zAs ~ 0.6544 and zLa ~ 0.1412, while minimized zAs ~ 0 .6589 and zLa ~ 0.1412. The presence of both electron and hole bands which cross the Fermi energy supports the observation of Hall coefficient temperature dependence. 90

Figure 3.19 Electron band structure of LaRhAsO (a), total and partial atomic density of states (b), partial density of states for d orbitals only (c), and magnified view near the Fermi level (d). Band structure calculated from energy minimized atomic positions is shown in black and red dashed lines indicate that determined using atomic positions from Rietveld refinement of PXRD data. 91

Figure 3.20 Electron band structure of LaIrAsO (a), total and partial atomic density of states (b), partial density of states for d orbitals only (c), and magnified view near the Fermi level (d). Band structure calculated from energy minimized atomic positions is shown in black and red dashed lines indicate that determined using atomic positions from Rietveld refinement of PXRD data. 92

Similarly to LaRhAsO, experimental and energy minimized atomic positions for LaIrAsO match well with one another; experimental zAs ~ 0.6604 and zLa ~ 0.1424, while minimized zAs ~ 0 .6637 and zLa ~ 0.1418.

The flat bands which lie just below the Fermi energy from Γ to Z and at the

Fermi energy from M to A are heavily 푑푥2−푦2 in character. The narrow band which lies just above the Fermi energy from Γ to Z is made from As p orbitals. By replacing

Rh with Ir two major changes appear to happen. First the 푑푥2−푦2 bands rise in energy so that they are now above the Fermi level and second the As p bands come down in energy so that they now either cross the Fermi level or lie just above it. These observations make sense in that LaIrAsO has a shorter a-parameter and should therefore have increased 푑푥2−푦2 overlap. This will destabilize these orbitals, but since

LaIrAsO has a longer M-As bond and wider MAs layer the p-orbital overlap should be slightly lowered, thus stabilizing the p-bands. In the case of LaIrAsO the As p band which now crosses the Fermi energy along Γ to Z creates a small hole pocket centered at the Γ point.

Looking next at the density of states, N(E), for LaRhAsO we can see that the

푑푥2−푦2 band which touches the Fermi energy at M and A causes a large density of states at the Fermi energy N(EF) to be observed, similar to what was observed for

LaCoAsO before including spin polarization effects [40]. This type of feature is 93 generally called a Van Hove singularity and can cause the spin up and spin down bands to split leading to an itinerant FM state, and indeed LaCoAsO is an itinerant FM

[41]. However, in the case of LaRhAsO a FM state is not observed and in fact the

Fermi level does not cut through the maximum in the N(E), but rather just below it.

Due to its shorter M-M distance the 푑푥2−푦2 bands of LaIrAsO are pushed above the

Fermi energy and a peak in N(E) is no longer observed anywhere near the Fermi level.

The N(E) plots show that N(EF) is greater for LaRhAsO than for LaIrAsO. While band structures of Ce analogs are not calculated here it has been shown for CeFeAsO [32] and CeRuPO [4] that including the Ce moment will lead to a 4f sub band forming 2-3 eV below the Fermi level, in addition to the Ln bands which are normally present 3-4 eV above the Fermi level. As the difference in N(EF) between Rh and Ir based is primarily driven the a-parameter reduction it is expected that the CeIrAsO will also have a reduced N(EF) relative to CeRhAsO. This conclusion supports the TK and TN temperature variation observed as discussed earlier in the context of the Kondo-

RKKY critical parameter N(EF)J.

The partial N(E) of Rh and Ir d orbitals reveals a complicated splitting where

푑푧2 and 푑푥2−푦2 states are highest in energy, which is contrary to the “3 over 2” d orbital arrangement typical of tetrahedral coordination. Interestingly the d orbital arrangement, with 푑푧2 and 푑푥2−푦2 states being highest in energy, looks much more 94 like a standard octahedral coordination. This result is in agreement with reports on

LaCoAsO [40,41] and essentially highlights the fact that M-M bonding is significant in these materials.

3.4 Conclusion New layered oxyarsenides containing the 4d and 5d metals Rh and Ir have been synthesized. Compositions from the series LnRhAsO and LnIrAsO show systematic decrease in lattice parameters and unit cell volumes as the rare earth ionic radius decreases. Upon substitution of Ir for Rh the a-parameter decreases which results in a shorter M-M interaction within the metal pnictide layer however the c-parameter elongates. CeRhAsO and CeIrAsO display a nonlinear temperature dependence of resistivity, indicating Kondo lattice behavior. Each shows a sudden drop in resistivity (7 K and 10 K respectively) coinciding with an AFM transition. This behavior is very similar to that reported for CeNiAsO, an AFM Kondo lattice material.

By comparing their TK and TN transition temperatures these compositions have been qualitatively arranged in order of increasing N(EF)J parameter it is suggested that they fall within the magnetic Kondo state regime where Kondo interactions dominate at high temperature and RKKY type interactions at low temperature. Band structure calculations reveal that LaRhAsO has an increased N(EF) relative to LaIrAsO and that its Fermi energy lies on the edge of sharp peak in the density of states. While further 95 measurements, such as heat capacity and high pressure resistivity, will help to better characterize these materials these findings demonstrate that exciting new ZrCuSiAs type compounds are still waiting to be explored.

3.5 Materials and Methods 3.5.1 Experimental Polycrystalline samples (0.5 - 1 g) were prepared in an Ar filled glove box using reagents of 3N (at. %) purity or better. Lanthanoid oxides (La2O3, CeO2, Nd2O3) were dehydrated overnight at 850 °C before use, and then reacted directly with lanthanoid metal (-40 mesh or smaller), Rh (or Ir) metal, and As metal. All reagents were finely ground in an agate mortar and pestle, pelletized, and placed in small

Al2O3 crucibles, which were then sealed inside evacuated fused silica ampoules (≈1 x

10-4 torr). A typical heating profile was: 20 K hr.-1 to 1273 K for 12-18 hr., then down

200 K hr.-1 to room temperature. Resulting products were reground, pressed, and sintered at 1173 K for an additional 12 hrs. to improve density before measuring electronic properties.

Room temperature lattice parameters were determined from powder x-ray diffraction data (PXRD) collected on a Rigaku MiniFlex II using a step size of 0.02° and a 2 second dwell time. Kα radiation was selected using a graphite 96 monochromator on the diffracted beam. Lattice parameters were initially obtained through least squares fitting via the software Powder4 [42]. Rietveld refinements of

LaMAsO and NdMAsO PXRD data was performed using GSAS via the expGUI interface

[43,44] over a 2θ range of ~ 15 - 100°. Four-probe resistivity (Cu wires with Ag epoxy), Hall coefficient (5 T static field), and magnetic measurements (AC 500 Hz 10

Oe, DC zero field cooled 0.5 T) were done from 3-300 K using a Quantum Design

Physical Properties Measurement System (PPMS). Room temperature thermopower was measured using a clamped electrode-heater system which capable of reaching a

ΔT≈10 K across the sample.

3.5.2 Computational Electronic band structure calculations were performed using the augmented plane wave plus local orbital method (APW+LO) as implemented Wien2K via the

W2Web interface [38] together with the generalized gradient approximation (GGA) for the exchange correlation functional in the Perdew-Burke-Ernzerhof (PBE) parameterization[45]. A cut off energy of -6 Ry. separating core and valence states was used. Linearized augmented plane wave (LAPW) muffin radii of 2.25a0 for La,

2.31a0 for Rh, 2.33 a0 for Ir, 2.05 a0 for As, and 2.0 a0 for O were used along with an

RmtKmax value of 7. The tetrahedron method [46] was employed to calculate the total and partial density of states on a 32 x 32 x 8 mesh (containing 544 k-points) inside 97 the Brillouin zone. All calculations were energy converged to 0.1 mRy and charge converged to 0.001 e or better.

98

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101

Chapter 4

The search for LaFeSbO and the synthesis and properties of La2PnO2 compositions (where Pn = Sb and Bi)

4.1 Abstract Density functional theory has been used to calculate the ground state heats of formation for LaFePnO compositions (Pn=P, As, Sb) as well as competing phases within the La-Fe-Pn-O systems. Results suggests that using standard high temperature synthesis methods will not yield LaFeSbO due to the high stability of competing phases, but that LaFeSbO is possibly metastable. The compound La2SbO2-x has been isolated and investigated in the course of searching for LaFeSbO and is found to be oxygen deficient via neutron powder diffraction with approximate composition of La2SbO1.7. Semiconducting behavior is observed and can be understood to come from large Sb displacement within the a-b plane. The nominal composition La2BiO1.5 was also investigated by neutron diffraction and is found to have an approximate formula of La2BiO1.9. Compared to La2SbO1.7 , La2BiO1.9 does not appear to show as much Pn disorder with the a-b plane.

102

4.2 Introduction The search for materials with useful and interesting properties often begins by investigating element substitution effects on the properties of a known material or structure. In fact the story of iron pnictide superconductivity highlights this discovery process on multiple levels. LnFePnO compositions (where Ln = Lanthanoid,

M = Late transition metal, and Pn = P, As) had been known to exist for some time

[1,2] before LaFePO was reinvestigated [3] due to its structural similarity with other wide band gap semiconductors having the composition LnCuChO (where Ln =

Lanthanoid and Ch = S or Se) [4]. Superconductivity was found for LaFePO at a critical temperature, Tc, of about 4 K and subsequently the substitution of As for P and partial substitution of F for O led to superconductivity, Tc = 26 K, for LaFeAsO0.89F0.11

[5] and a flurry of even higher Tc discoveries [6-11]. The obvious question to ask is then, can oxyantimonide and oxybismuthide analogs, such as LaFeSbO and LaFeBiO, be stabilized? or any other LnMSbO or LnMBiO compositions? If so, what properties might we expect? To answer this last question about LaFeSbO a number of density functional theory (DFT) calculations have been done [12–15] and these seem to indicate that Sb based compositions should have enhanced magnetic interactions relative to As compositions and thus, depending on the mechanism responsible for

Cooper pairing in these materials, an increased Tc may be found. It is has been a 103 while since LaFeAsO was first reported however and still there have been no experimental reports of LaFeSbO (or LaFeBiO).

The compositions LnMnSbO and LnZnSbO have been reported [16,17] and contain Sb3- anions as confirmed by 121Sb Mössbauer spectroscopy [18,19]. LaZnSbO and CeZnSbO are reported to show a metal to insulator transition upon cooling while the NdZnSbO is metallic [20]. The bismuthide LaNiBiO has also been reported

[21] and has been cited in the literature as one example of a layered oxy-bismuthide superconductor (Tc~4.4 K) [22–24]. However, the PXRD pattern reported for LaNiBiO

(P4/nmm #129) is missing a strong (112) reflection at approximately 36 degrees 2θ if it was in fact P4/nmm. Instead its pattern matches well with that of another layered oxypnictide, La2BiO2-x, for which we present structural and electronic properties, indicating it is likely a case of mistaken identity.

As will be discussed below we have attempted to synthesize the solid solution

LaFeAs1-xSbxO and there has been at least two other reports in the literature on the properties of this solid solution for x ~ 0.3-0.4 [25,26]. In addition we have investigated the LaZn1-xFexSbO solid solution, but it appears LaFeSbO cannot be stabilized. Thus the question begs to be asked, why won’t these heavier analogs form? As others have pointed out, it may simply be the high stability of iron antimonide impurity phases and the desire of Sb to form Sb-Sb bonds which prevents the formation of LaFeSbO [15]. Nevertheless, reports of compositions such as 104

CsFe2Sb2 [27] give hope to the search for an energetically stable iron antimonide composition with the same tetrahedral metal-pnictide coordination environment observed for LnFeAsO [14]. To gain further insight into the still elusive LaFeSbO, we used density functional theory to calculate formation energies for phases within the

La-Fe-Sb-O system.

Our efforts to realize LaFeSbO and other new LaMSbO and LaMBiO compositions (M = Fe - Ni) also led us to identify and characterize the layered oxypnictide compounds La2SbO2 and La2BiO2 [28]. These layered oxypnictides are peculiar in that the pnictide anions have an average charge state of Pn2- if all atomic positions are fully occupied; however the powder neutron diffraction studies presented here indicate these compositions can accommodate some disordered oxygen vacancies.

La2PnO2 compounds (where Ln = Sb and Bi) crystallize in a tetragonal ThCr2Si2 type structure (I4/mmm #139) with O coordinated to La in a fluorite like arrangement. These layers are stacked alternatively along the c-axis with a square mesh sub-lattice of Pn anions. Each Pn is in square prismatic coordination with

La atoms from the fluorite layers, see Figure 4.1. 105

Figure 4.1 Crystal structure of La2PnO2 compositions (where Pn = Sb and Bi). The unit cell (a) contains O in tetrahedral coordination to La cations creating fluorite like slabs while the Sb anions adopta square prismatic coordination to La ions between the fluorite slabs.

In fact the ThCu2Si2 archetype has anti-fluorite layers separated by cations and as such the La2PnO2 structure is sometimes referred to as the anti-ThCu2Si2 structural type. based compositions were originally reported by Benz [29]; however Nuss and Jansen have recently reinvestigated Ce2PnO2 as well as Pr2PnO2 single crystals and found evidence of displacement within the a-b plane, resulting in both short (3.2 Å) and long (4.8 Å) antimony distances within the plane 106

[30]. In the course of this research La2BiO2 and the Ln2BiO2 (where Ln = Y, La-Er) compositional series were also investigated by other researchers [31] and more recently the transport properties of Ln2SbO2 compositions (where Ln = La-Er ) have been reported [32]. Telluride analogs with the general formula Ln2TeO2 are also known [33] and appear to be indirect band gap semiconductors [34] but otherwise no other anti-ThCr2Si2 oxypnictide or oxychalcogenide compositions are known.

4.3 Results and discussion 4.3.1 The search for LaFeSbO

4.3.1.1 Investigation of LaFeSb1-xAsxO and LaFe1-xZnxSbO solid solutions Initial investigations of LaFeSbO focused on two solid solutions, one between

LaFeAsO and LaFeSbO and another between LaZnSbO and LaFeSbO. LaFeSb1-xAsxO compositions have Fe/Sb ratios greater than one, while LaFe1-xZnxSbO compositions have Fe/Sb ratios less than one and examination of these two solid solutions therefore complement one another. The lattice parameters for LaFeSb1-xAsxO and

LaFe1-xZnxSbO compositions are shown in Figure 4.2; however the Fe/Sb < 1 compositions are shown plotted with the x scale reversed in order to facilitate comparison of the two solid solutions. In addition, density functional theory (DFT) calculated lattice parameters for LaFeSbO from the work presented here [28] and 107 two others within the literature [12,13] are shown plotted in Figure 4.2, as well as the lattice parameters recently reported for LaFeAs1-xSbxO compositions by other researchers [26]. The a and c lattice parameters for compositions farthest from

LaFeSbO (i.e. x = 0.5 to 1) can be linearly fit in order to extrapolate the expected parameters of LaFeSbO and these are shown overlaid in Figure 4.2. In all cases the amount of FeSb impurity which is present increases steadily as the Fe/Sb ratio approaches one, indicating LaFeSbO is generally not favored and that the x values used to plot lattice parameters are by no means quantitative; they are simply the nominal values used during synthesis. Nevertheless there is a small decrease in the refined parameters, for LaFe1-xZnxSbO compositions, where x = 1-0.3 (i.e. starting at the Zn rich side) which indicates some Fe may be soluble in LaZnSbO. At x = 0.3 two phases can be identified (in addition to FeSb); the primary phase LaZnSbO having

P4/nmm symmetry and a secondary ‘pseudo P4/nmm’ phase. It is called ‘pseudo-

P4/nmm’ in that it appears to be P4/nmm but is missing a (112) hkl reflection at

2θ ≈ 36° if this were truly the case. Beyond y = -0.3 only FeSb and the ‘pseudo

P4/nmm’ phase are observed. When starting from LaFeAsO the lattice parameters for LaFeSb1-xAsxO compositions appear to shift upward incrementally (i.e. when moving from x = 1 to 0.3) and as Fe/Sb approaches one (x < 0.3) the primary phases observed are once again FeSb and the ‘pseudo P4/nmm’ phase. As discussed below, 108 this ‘pseudo P4/nmm’ phase turns out to be the composition La2SbO2, a layered oxypnictide with I4/mmm symmetry that was previously unreported.

Figure 4.2 The lattice parameters for LaFeAs1-xSbxO and LaFe1-xZnxSbO compositions.

4.3.1.2 Identification of La2SbO2 and a case of mistaken identity Figure 4.3 shows the PXRD pattern obtained when reacting La:Fe:Sb:O in a

1:1:1:1 stoichiometry within evacuated fused silica tubes. The term ‘pseudo

P4/nmm’ is used above because at first glance it appears that the primary phase can 109 be indexed as the composition LaFeSbO with P4/nmm symmetry; however a strong reflection is missing at 2θ ≈ 36° and therefore this space group assignment must be incorrect. By indexing all peaks peak positions other than those from FeSb and La2O3, using the programs DICVOL [35] and CMPR [36], this phase was identified as a having

I4/mmm symmetry. It was also determined that the c-parameter for this I4/mmm phase is about 1.5 times that calculated if the P4/nmm space group is assumed.

Using this information to search the ICSD database [37] revealed the ‘pseudo

P4/nmm’ phase to be La2SbO2 an unreported analog to Ce2SbO2. The calculated pattern for La2SbO2 is shown at the bottom of Figure 4.3.

It is important to highlight the ability to mistake the La2SbO2 pattern for

LaFeSbO because this is in fact appears to be what happened in the case of

LaNiBiO0.8. The composition LaNiBiO0.8 has been reported to crystallize in the

P4/nmm space group with a ~ 4.07 Å and c ~ 9.30 Å [21]; however the PXRD pattern shown within the paper is missing the (112) reflection predicted for P4/nmm symmetry and this discrepancy goes unaddressed by the authors. Instead the pattern matches that expected for La2BiO2, where the space group is I4/mmm, a ~ 4.07 Å, and c ~ 13.95 Å. Thus it appears that the only report of a ZrCuSiAs type layered oxybismuthide in the literature is a case of mistaken identity. 110

Figure 4.3 The experimentally observed PXRD pattern for La:Fe:Sb:O reactions along with the simulated patterns for LaFeSbO (P4/nmm) and La2SbO2 (I4/mmm).

4.3.1.3 Calculation of heats of formation within the La-Fe-Sb-O system To gain further insight into the La-Fe-Sb-O system and the relative stability of

LaFeSbO compared to compositions such as LaFePO and LaFeAsO density functional theory (DFT) calculations were carried out. Details of the calculations are described in the Materials and Methods section. Calculated zero temperature structural parameters based on DFT/GGA for stoichiometric La2SbO2 are a = 4.060(1) Å, 111 c = 13.97(1) Å, and zLa = 0.3392(1), which agree well with the experimental numbers as described below. Similarly, the calculated tetragonal lattice parameters for

LaFeAsO, a = 4.020(1) Å, c = 8.640(1) Å, and LaFePO, a = 3.930(1) Å, c = 8.500(1) Å, agree fairly well with reported room temperature lattice parameters. The energy minimized structural parameters found for LaFeSbO are a = 4.168(1) Å, c = 9.214(1) Å, zLa = 0.1283(1), zSb= 0.6595(1), which are in good agreement with other reports on LaFeSbO [12,13]. These DFT obtained lattice parameters are shown plotted in Figure 4.2. The heat of formation at zero temperature of LaFePnO

(Pn = P, As, Sb) relative to the standard state of the constituent elements (solid P, As,

Sb, La, ferromagnetic Fe, and gaseous molecular oxygen) are summarized in Table

4.1.

Table 4.1 Zero temperature heat of formation energy per formula unit of LaFePnO relative to the standard state of the constituent elements and relative to more common starting reagents.

Stability equations for LaFePO, LaFeAsO, and LaFeSbO ΔH0 K (eV)

La(s) + Fe(s) + P(s) + ½O2(g) → LaFePO -8.3

La(s) + Fe(s) + As(s) + ½O2(g) → LaFeAsO -7.7

La(s) + Fe(s) + Sb(s) + ½O2(g) → LaFeSbO -6.7

⅓La(s) + ⅓La2O3(s) + Fe(s) + P(s)  LaFePO -2.4

⅓La(s) + ⅓La2O3(s) + Fe(s) + As(s)  LaFeAsO -1.8

⅓La(s) + ⅓La2O3(s) + Fe(s) + Sb(s)  LaFeSbO -0.9

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Table 4.2 Zero temperature heat of formation energy per formula unit of LaFePnO relative to the most stable competing phases.

Stability equations for nearest competing phases to ΔH0 K (eV) LaFePO, LaFeAsO, and LaFeSbO

La2O3(s) + LaAs(s) + FeAs2(s) + 2Fe(s) = 3LaFeAsO -0.31

La2O3(s) + LaFe2P2(s) + FeP(s) = 3LaFePO(s) -0.13

La2SbO2(s) + FeSb(s) + Fe(s) = 2LaFeSbO(s) 0.27

We find a negative heat of formation for LaFeSbO but its magnitude is smaller than the heat of formations of LaFePO and LaFeAsO. In a search for competing phases in the space of the constituent elements La, Fe, O, and (P, As, Sb) we have calculated total energies for 68 compounds. In addition the heat of formation energies relative to the common starting reagents used in this study are shown in

Table 4.1.

In Table 4.2 the formation energies per formula unit of LaFePnO versus its nearest competing phases are summarized. We find LaFePO and LaFeAsO to be thermodynamically stable relative to their nearest competing phases, as they must, but LaFeSbO is thermodynamically unstable relative to the formation of La2SbO2 by

ΔH~0.1 eV per formula unit of LaFeSbO. Two questions remain: is LaFeSbO a potentially metastable compound and if so, can it be made? To answer the first question, we have calculated the phonon spectrum of LaFeSbO, shown Figure 4.4, and found no imaginary phonon frequencies. Hence there are no lattice instabilities 113 and the formation energy of LaFeSbO is indeed a local minimum. To answer the second question would require as a first step the determination of the depth of the energy well around the local minimum for the LaFeSbO compound. This is a formidable task, which would require the determination of an unknown number of transition states and we must leave our second question unanswered. We conclude that LaFeSbO is thermodynamically unstable relative to the formation of La2SbO2, but can potentially exist in a metastable state. Our calculations allow no prediction of the stability range of this metastable state

Figure 4.4 Phonon spectrum for LaFeSbO as found using the DFT energy minimized structural parameters, a = 4.168Å c = 9.214Å zLa= 0.1283 and zSb = 0.6595. . 114

4.3.2 Further investigation of La2PnO2 compositions (where Pn = Sb and Bi)

4.3.2.1 Synthesis and initial x-ray diffraction studies In the course of our efforts to realize LaMPnO (M = Fe - Ni and Pn = Sb, Bi) compositions, using standard solid state methods, the La2PnO2 phases have consistently been the primary products which form along with transition metal pnictide phases, such as FeSb. The general structure of La2PnO2 compositions is introduced above. First efforts to synthesize a pure phase sample of La2SbO2 for characterization always showed La2O3 products despite careful dehydration of all reagents and ampoules as well as the use of a Zr foil oxygen getter and alumina reaction crucible inside the ampoule. High quality samples were only obtained when the nominal oxygen content was reduced to approximately La2SbO1.5-1.6. The oxygen content was controlled by limiting the amount of La2O3 and therefore compensating with La metal to maintain a 2:1 ratio of La:Pn. Similar to La2SbO2, high quality samples of the analog could only be produced for nominal compositions around La2BiO1.5. The need to reduce oxygen content in order to produce high quality samples implies that oxygen vacancies may exist within the La2PnO2 structure.

In fact the oxygen deficient formulas La2PnO2-x (x ≥ 0.5) can be charge balanced assuming an average pnictide charge greater that Pn2- . The intergrowth composition Y2SrFeCuO6.5 is known to contain 25% oxygen vacancies within a 115 and oxygen fluorite type layer[38]. In that case however the compensating yttrium and surrounding oxygen atomic displacements cause the symmetry to shift from tetragonal to orthorhombic. The PXRD patterns observed for the nominal compositions La2SbO1.5 and La2BiO1.5 are shown in Figure 4.5. As shown in Figure 4.5 we have not observed any visible splitting of (hkl) reflections as would be expected if oxygen vacancies were ordered and the surrounding atoms displaced.

Figure 4.5 The PXRD patterns for La2SbO2-x and La2BiO2-x compositions where nominally x = 0.5. Also shown are the general atomic positions for La2PnO2 materials and refined lattice parameters for each composition.

116

Lattice parameters and cell volume determined for La2SbO2-x, where nominally x = 0.5, are refined to be a = 4.05653(32) Å and c = 13.8559(13) Å.

Comparing these parameters to those found for LaFeSbO via DFT (discussed below) it is readily observed that the Sb coordination environment in LaFeSbO is truncated compared to that of La2SbO2 because of Fe-Sb bonding. As expected from atomic size considerations, there is an increase in the unit cell lattice parameters of La2BiO2-x samples, a = 4.07858(17) Å and c = 13.9114(7) Å. Table 4.3 summarizes the room temperature lattice parameters determined for these compositions and those found in other recent studies, as well as those determined via powder neutron diffraction in this study. Further details of the neutron diffraction studies are discussed below.

The lattice parameters for La2Bi0.75Sb0.25O1.5 were found to be a = 4.0785(15) Å and c = 13.8982(6) Å while those of La1.9Sr0.1SbO1.5 were found to be a = 4.06079(34) Å, c = 13.9773(15) Å.

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Table 4.3 Lattice parameters determined for La2SbO2 and La2BiO2 compositions Nominal Experimental a-parameter c-parameter Reference Composition Composition (Å) (Å)

La2SbO2 single crystal - 4.067(1) 13.70(1) [32] x-ray

La2SbO2 powder x-ray mixed phase 4.063(1) 13.96(1) This work

La2SbO1.5 powder x-ray single phase 4.056(1) 13.86(1) This work

La2SbO1.5 powder La2SbO1.7 4.063(1) 13.81(1) This work neutron + 6 wt% LaSb

La2BiO2 powder x-ray 4.085(1) 13.99(1) [31]

La2BiO2 powder x-ray mixed phase 4.081(1) 13.93(1) This work

La2BiO1.5 powder x-ray single phase 4.079(1) 13.91(1) This work

La2BiO1.5 powder La2BiO1.9 4.087(1) 13.96(1) This work neutron

4.3.2.2 Neutron diffraction studies of the nominal compositions

La2SbO1.5 and La2BiO1.5 As mentioned above, single phase samples were only produced when the nominal oxygen content was reduced suggesting that either some oxygen vacancies are stabilized within the structure or reducing the oxygen content simply offsets the stray oxygen present in the synthesis process. The oxygen content in these materials is actually quite important in that when all atomic sites are fully occupied the charge state of the pnictide anions is effectively Pn2-; if the oxygen site is less than fully occupied then the average charge state of the pnictide anions may be greater than

Pn2-. The Pn2- charge state is relatively uncommon, or unusual as described in Ref.

[31], an oxygen vacancies may help to shift the average charge state back towards 118

Pn3-. Neutron diffraction is quite useful in characterizing these materials therefore in that it allows for the precise refinement of oxygen occupancies (and thermal parameters) even in the presence of the relatively heavy elements (such as La, Sb, and Bi).

The major results of Rietveld refinement are shown tabulate in Table 4.4.

Similar to the reports on Ce2SbO2 [30] a better fit to the La2SbO1.5 data is obtained by using a split site model, otherwise a large Sb thermal parameter is observed within the a-b plane as shown in Table 4.4. The split site model allows the Sb atoms to displace within the a-b plane and more realistic thermal parameters are observed, see Table 4.4. For La2BiO1.5 on the other hand, the difference between the non-split and split site models is much less, with the non-split site model giving a slightly better fit. The propensity for Sb-Sb square nets to distort is well studied in the literature

[39]. In fact a 4 X 4 superstructure of ordered Sb dimers has been reported for single crystal Pr2SbO2 samples thereby giving both long and short Sb-Sb distances within the a-b plane [30]. The Sb offset observed using the split site model amounts to a distance of about 0.40 Å from the ideal positions and results in short and long Sb-Sb distances within the a-b plane of 3.26 Å and 4.86 Å.

Refinement of atomic site occupancies reveals that some oxygen vacancies are present in these materials and that the oxygen site is more deficient in the

La2SbO1.5 sample than in the La2BiO1.5 sample; however, no super structure peaks are 119 observed indicating they are disordered. The refined compositions are La2SbO1.7 and

La2BiO1.9 indicating that despite best attempts to control the atmosphere there is likely stray oxygen present in the synthesis process. Given that the La ions will adopt

3+ a La state. The reduced oxygen content for La2SbO1.7 samples implies that the average charge state of the pnictides anions should be Sb2.6- a value much closer to the typical charge state of Sb3- as found in LaMnSbO [18]; however, Mössbauer studies of these compositions will help to directly probe the charge state of the antimony anions. The average charge state in the bismuth sample on the other hand is Bi2.2-. It is interesting to note that the Sb sample has more oxygen vacancies and shows the greater Pn displacement within the plane. The Sb displacement for the

La2SbO1.7 sample examined here is also greater than the displacement parameter recently reported for La2SbO2 by other researchers using x-ray diffraction and a completely different synthetic approach[32]. This brings up the question of how strongly are the oxygen vacancies and Sb displacement coupled. Future studies of these materials should therefore attempt to link these structural features and check to see if fully stoichiometric La2SbO2 samples show strong displacement. If not they do not then it may be oxygen vacancies which are driving the displacement to occur at all in these materials. However, as stated before Sb square net distortion is a common phenomenon across a wide range of material types and thus the more likely explanation is that the displacement is simply enhanced by introducing more oxygen 120 vacancies, but there is still some Sb displacement present even in the fully oxygenated samples.

Table 4.4 Powder neutron diffraction Rietveld refinement results for La2SbO1.5 and La2BiO1.5 nominal compositions. La SbO La BiO Nominal La SbO 2 1.5 La BiO 2 1.5 2 1.5 split model 2 1.5 split model Composition (300 K) (300 K) (300 K) (300 K) a (Å) 4.06265(8) 4.06321(14) 4.08738(8) 4.08740(8) c (Å) 13.8052(5) 13.8031(8) 13.9641(4) 13.9646(5) V (Å3) 227.857(12) 227.884(22) 233.293(12) 233.303(13) zLa 0.34001(10) 0.33965(12) 0.33664(10) 0.33658(10) x Sb 0.0984(14) 0.0501(10) (split model) O site 0.848(6) 0.872(7) 0.957(5) 0.958(5) occupancy La 0.01023(32) 0.0088(4) 0.00561(22) 0.00543(23) U11=U22=U33 O 0.0071(6) 0.0101(8) 0.0062(4) 0.0054(4) U11=U22=U33 0.1055(28), 0.0184(23) 0.0270(6) 0.0144(5) Sb 0.1074(28) 0.0184(23) 0.0270(6) 0.0144(5) U ,U ,U 11 22 33 0.0159(22) 0.0184(23) 0.0079(8) 0.0144(5)

Rwp (%) 5.81 5.43 5.05 5.21 GOF value 2.35 2.21 1.75 1.80 (χ) 2θ Range (°), 15-145, 0.05 15-145, 0.05 15-145, 0.05 15-145, 0.05 step size (°)

Pn - Pn (Å) 4.06265(8) 4.06321(14) 4.08738(8) 4.08740(8) Pn - La (Å) 3.6237(9) 3.642(4) 3.6820(8) 3.6864(9) La - O (Å) 2.3812(8) 2.3788(9) 2.3750(7) 2.3746(7) La - O - La (°) 117.09(6) 117.31(7) 118.75(6) 118.78(6)

121

Figure 4.6 The powder neutron diffraction pattern for the nominal composition La2SbO1.5. Refinement of atomic site occupancy results in a composition of La2SbO1.7.

Figure 4.7 The powder neutron diffraction pattern for the nominal composition La2BiO1.5. Refinement of atomic site occupancy results in a composition of La2BiO1.9. 122

4.3.2.3 Electronic properties of the nominal compositions La2SbO1.5 and La2BiO1.5 La2SbO1.5 samples and La2BiO1.5 display semiconducting behavior with respect to temperature. Typical thermally activated semiconductors obey the relationship,

푇 0 휌 = 휌0푒 푇 and thus a plot of natural log resistivity versus reciprocal temperature should be linear. The natural log of resistivity as a function of reciprocal temperature is shown for La2SbO1.5 and La2BiO1.5 samples in Figure 4.8 as well as La2Bi0.75Sb0.25O1.5 and La1.9Sr0.1SbO1.5. The slope from this plot, T0, relates to the semiconductor band gap via T0 = Eg/2kB (kB, Boltzmann constant)

Figure 4.8. Natural log of resistivity for La2SbO1.5,La2BiO1.5, La2Bi0.75Sb0.25O1.5, and La1.9Sr0.1SbO1.5 compositions. The slope, T0, is related to the band gag, T0 = Eg/2kB, and the calculated Eg values at both high and low temperatures are shown. 123

In fact the resistivity of La2SbO1.5, La2Bi0.75Sb0.25O1.5, and La2BiO1.5 appears

푇 1/4 0 more linear when plotted according to 휌 = 휌0푒 푇 and this suggests that three dimensional variable range hopping (VRH) may be at play. VRH hopping is a generally associated with conduction in disordered materials [40] so it makes some sense that it would be observed in these materials given their large disorder of the Sb site and that conduction in these materials is expected to be dominated by Sb states [31,32].

However, VRH conduction is generally a low temperature phenomenon and the resistivity of these materials appears to show linear behavior up to 300 K when plotted on a T-1/4 scale. It is also interesting to note that even though the materials are clearly two dimensional in nature, it appears that the three dimensional variable range hopping model gives a slightly better fit than the two dimensional model,

푇 1/3 0 휌 = 휌0푒 푇 [40].

Within the VRH model the value T0 is related to the amount disorder present and therefore the energy needed for a carrier to tunnel from one site to another.

More specifically it is inversely proportional to the density of states at the Fermi level, N(EF), and the localization radius of the carriers, rloc. The localization radius effectively represents the distance that the carriers can tunnel. The plot of natural

-1/4 log of resistivity versus T is shown in Figure 4.9 along with the values of T0 obtained. Regardless of what model is used, thermally activated or VRH, the 124 increased slope for La2SbO1.5 samples indicates that it takes more energy for conduction to occur in La2SbO1.5 than it does for La2BiO1.5. This is expected within the

VRH model, as the La2BiO1.5 samples do not show as significant Pn site disorder according to structural refinement.

Figure 4.9. Natural log of resistivity for La2SbO1.5,La2BiO1.5, and La2Bi0.75Sb0.25O1.5 -1/4 plotted as function of T . The value T0 is obtained by evaluating the slope of the linear best fit line.

125

Actually band structure calculations for La2SbO2 and La2BiO2 [31] predict metallic conduction, as indicated in Figure 4.10 by a Fermi level which lies midway through the Sb px and py bands.

Figure 4.10 The band structure for La2SbO2. A primitive cell was used for calculation cell rather than I-centered due to computational ease. The Brillouin zone is shown in the upper corner.

The dispersive band which crosses the Fermi energy forms a hole pocket centered at the gamma point of the primitive Brillouin zone and arises from the overlap of pnictide px and py orbitals. The observation of non-metallic behavior for

La2BiO2 based materials has been explained to arise from their relatively large a-parameters (Bi-Bi distance) which causes these materials will have poor Bi orbital 126 overlap [31]. Reduced carrier mobility due to the narrower px and py band can give a poorly metallic/insulating state and indeed using smaller rare earths like Er or Y for La causes the a-parameter to contract and a metallic state to emerge [31]; supporting the conclusion that px and py orbital overlap is a major factor in determining transport properties.

For Sb compositions a clear semiconductor behavior is seen even when small rare earths such as Er are used [32]. The behavior of Ln2SbO2 analogs is perhaps

4- 2- better explained by considering that either Sb2 dimer formation or random Sb disorder should cause semiconducting behavior [30, 40]. In the case of dimer formation, the energy separation between bonding and antibonding states of Sb dimers introduces a pseudogap at the vicinity of the Fermi level, while significant Sb atomic disorder can lead to the localization of states at the top and bottom of bands around the pseudo gap which is created[40]; this behavior is called Anderson localization. The semiconductor behavior (with small activation energy) then naturally arises from the energy needed to populate carriers across this gap and beyond the localized states near the pseudo gap.

In fact a distortion driven metal to insulator transition is exactly what has been observed for SrZnSb2 and SrZnBi2 which both contain square nets of Sb and Bi anions [39]. In SrZnSb2 the Sb square net distorts to form zig-zag chains of Sb atoms however in SrZnBi2 the Bi anions remain undistorted and it was shown that the 127 electronegativity difference between the Pn square mesh and the other portion of the structure (a [SrZnPn] layer for SrZnPn2 materials or a [La2O2] for La2PnO2 materials) is what determines if a distortion of the Pn square mesh occurs [39]. The key difference between these layered oxypnictides and the SrZnPn2 materials is that the formation of ordered zig-zag chains in the latter gives rise to a tetragonal to orthorhombic transition. In case of La2SbO1.5 no orthorhombic transition is observed; however, the Sb displacement gives rise to Sb-Sb bond distances of around 3.26 Å, which are near those found in elemental antimony [41]. Both the observed structural and electronic properties therefore support the idea that disordered Sb-Sb dimers may be forming within the a-b plane. Further experimental and band structure studies will be needed however to examine the extent to which introducing additional electron carriers and structural disorder, through oxygen vacancies, affects

Sb displacement and physical properties.

4.4 Conclusion While many researchers have attempted the synthesis of LaFeSbO thus far there have still been no reports of LaFeSbO. Density functional theory has been used to calculate the ground state heats of formation for LaFePnO compositions (Pn=P, As,

Sb) and competing phases within the La-Fe-Pn-O systems. The and arsenic analogs are found to be stable with respect to their elemental constituents as 128 well as nearest competing phases, while LaFeSbO is found to be stable only with respect to the elemental constituents. Our work suggests that using standard high temperature synthesis methods will not yield LaFeSbO due to the high stability of competing phases such as La2SbO2 and FeSb. The ground state phonon spectrum of

LaFeSbO reveals no lattice instabilities and suggests that LaFeSbO is possibly metastable. Whether or not LaFeSbO (or similar compositions containing Fe2Sb2 layers separated by a 2+ charge donor layer) can be formed remains an open question but sequential layering of Fe, Sb, and charge donor constituents may still allow these thermodynamically favored phases to be sidestepped and for LaFeSbO to be experimentally characterized.

The compound La2SbO2 has been isolated and investigated in the course of searching for LaFeSbO. La2SbO1.5 nominal compositions are found to be oxygen deficient according to neutron powder diffraction and to have a composition of

3- approximately La2SbO1.7 which should shift the average pnictide charge closer to Pn .

Semiconducting behavior is observed similar to the previously reported compositions

Ce2SbO2 and Pr2SbO2 and can be understood to come from the Sb displacement within a-b plane causing a gap to open at the Fermi surface and thereby destroying the metallic state predicted from band structure calculations.

The nominal composition La2BiO1.5 was also investigated by neutron diffraction and is found to have an approximate formula of La2BiO1.9. Compared to 129

La2SbO1.7 using a split site model for La2BiO1.9 does not result in a better overall refinement indicating that Bi displacement is not as significant or in this material.

While weakly insulating electronic behavior is observed in La2BiO2 when metallic conduction is expected, in this case the reduced conductivity is understood to arise from the poor Bi px and py orbital overlap within the a-b plane, not from Bi-Bi dimer formation.

4.5 Materials and Methods 4.5.1 Experimental

Reagents used throughout this study include La2O3 (99.9 at%), La (99.9 at%),

Sb (99.5 at%), and Bi (99.9 at%). The composition La1.9Sr0.1SbO2 was investigated by substitution of La2O3 by SrO (99 at%). All starting materials were weighed, ground, and pressed into pellets inside of an Ar filled glove box before sealing in fused silica

-4 ampoules, evacuated to approximately 1 x 10 torr. La2O3 was heated overnight at

1023 K while SrO was heated at 1473 K before transferring to the glove box. First efforts to synthesize a pure phase sample of La2SbO2 for characterization always showed La2O3 products despite careful dehydration of all reagents and ampoules as well as the use of a Zr foil oxygen getter and alumina reaction crucible inside the ampoule. High quality samples were only obtained when the nominal oxygen content was reduced to approximately La2SbO1.5-1.6. The oxygen content was 130 controlled by limiting the amount of La2O3 and therefore compensating with La metal to maintain a 2:1 ratio of La:Bi. Similar to La2SbO2, high quality samples of the bismuth analog could only be produced for nominal compositions around La2BiO1.5.

The sealed ampoules were typically heated at 1223 K for 12-18 hrs. in a programmable box furnace. To achieve a pellet suitable for property measurement it was necessary to regrind, press, and heat again at 1173-1223 K for 24 hrs. All

La2SbO2-x compositions examined are air sensitive and therefore stored in an Argon glove box or under vacuum.

Resistivity was determined using a Quantum Design Physical Properties

Measurement System (PPMS) using a four probe method with copper wires and paint to make electrode contact. Samples were prepared in air immediately before measuring to minimize exposure and sample decomposition is practically unnoticeable by XRD before and after the measurement. La2SbO2 appears black with a slightly reddish hue. Crystal structures at room temperature were characterized by powder X-ray diffraction (PXRD) using a Rigaku MiniFlex2 scanned from 2θ = 20−68° with a step size of 0.02° and a graphite monochromator on the diffracted beam

(Cu-Kα). Powder neutron diffraction patterns were collected at the Australian

Nuclear Science and Technology Organization (ANSTO) Bragg institute on the Echidna instrument [42] via their mail in proposal service. Data was collected at both 3 K and

300 K from approx. 15 - 145° 2θ using a radiation wavelength of 1.62150 Å selected 131 with a Ge(335) monochromator. Rietveld refinement of PXRD and PND patterns was carried out using GSAS [43] via the ExpGUI software[44]. The PXRD peaks corresponding to La2SbO2 were originally indentified and indexed using the program

DICVOL [35] via CMPR [36] from multi phase products obtained while attempting to synthesize LaFeSbO.

4.5.2 Calculations Lattice parameters and heat of formation energies at zero temperature were calculated using ab initio density functional theory (DFT). The projector augmented plane wave method as implemented in the Vienna ab-initio simulation package

(VASP) [45][46] was used together with the generalized gradient approximation

(GGA) for the exchange correlation functional in the Perdew-Burke-Ernzerhof (PBE) parameterization[47]. Total energies were converged to less than 1 meV/atom by controlling all calculation parameters (planewave cutoff, Fermi broadening and k- point mesh density). The band structure presented for La2SbO2 was calculated using

Wien2K via the W2Web interface [48]. Lattice parameters as determined from neutron diffraction experiments were used during the calculation and the Sb position was fixed at (0, 0, 0) (i.e. the Sb site was not split). PBE-GGA parameterization was used with a cut off energy of -6 Ry. separating core and valence states. Linearized augmented plane wave (LAPW) muffin tin radii of 2.21a0 for La, 1.96 a0 for O, and 132

2.50 a0 for Sb were used along with an RmtKmax value of 7. All calculations were energy converged to 0.1 mRy and charge converged to 0.001 e or better.

133

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136

General Conclusions and Future Work

The structure-property relationships of new layered oxypnictides have been explored. A full solid solution between LaFeAsO and LaRhAsO has been synthesized and superconductivity is observed with a maximum critical temperature of 16 K is observed for 10 at% Rh. Correlations between As-Fe-As bond angle and Tc do not to hold when doping group 9 transition metals onto the Fe site most likely due to the disorder introduced into the FeAs layer, which is responsible for conduction in these materials. The fact that superconductivity is observed at all when doping onto the Fe site is an indication that the mechanism responsible for electron pairing in LaFeAsO materials is quite robust and only partially sensitive to the effects of conduction layer disorder. By examining the temperature coefficient of resistivity an electronic phase diagram for the LaFe1-xRhxAsO system has been constructed. LnRhAsO and LnIrAsO

(where Ln = La, Ce, Nd) have been synthesized and are found to be metallic. The Ce based compositions display a resistivity temperature dependence which is characteristic of a magnetic Kondo state. At high temperature the scattering is dominated by Kondo scattering between Ce4f1 electrons and conduction electrons while at low temperature RKKY interactions dominate and AFM is observed. Their behavior is quite similar to other Ce based compositions which have been reported such as CeNiAsO. The Kondo scattering temperatures and magnetic transition temperatures observed for CeRhAsO and CeIrAsO can be rationalized by qualitatively 137 considering the relative difference in Fermi level density of states (NEF) and magnetic exchange parameter, J, for these compositions. Future work of interest would be to examine the solid solutions CeFe1-xRhxAsO and CeFe1-xIrxAsO and thereby possibly explore the evolution from a superconducting ground state to a magnetic Kondo state. Additionally it would be fundamentally interesting to investigate the change in properties for CeRhAsO and CeIrAsO under high pressure. If their placement along the N(EF)J scale is correct then as J increases with pressure a lowering of the magnetic ordering temperature should be observed and eventually it should be suppressed altogether. On the other hand CePd2Si2 is known to exhibit similar behavior to these compositions and at high enough pressure it changes from a magnetic Kondo to superconducting ground state. Given the structural similarity (each has the same anti-fluorite transition metal layer) and that the PdSi layer effectively has the same count as a RhAs (or IrAs) layer (PdSi is a group 10 metal with a group

14 anion, while RhAs is a group 9 metal with a group 15 anion) it would interesting to see if CeRhAsO or CeIrAsO displayed this transition as well. In the very near term band structure calculations of CeRhAsO and CeIrAsO should be compared and contrasted with those of CePd2Si2 and used to quantitatively calculate the expected difference in N(EF)J for these compositions.

In addition to exploring LnMAsO compositions (where M = Rh and Ir), the composition LaFeSbO has been investigated both experimentally and through the use 138 of density functional theory calculations. While partial solid solutions with LaFeAsO and LaZnSbO have been looked at LaFeSbO has not yet been stabilized. Density functional theory calculations of the heats of formation for competing phases within the La:Fe:Sb:O system reveal that it LaFeSbO is unstable with respect to the phases

La2SbO2 and FeSb, which are all that have been observed in experiment; however it does appear stable with respect to bare elements and common starting reagent. In addition the phonon dispersion spectrum also displays no indication of lattice instability and thus LaFeSbO appears to be metastable. The phases La2SbO2 and

La2BiO2, which were uncovered in the search for LaFeSbO and LaFeBiO compositions, have been explored further primarily through resistivity and neutron diffraction experiments. Powder neutron diffraction suggests that the samples prepared in this study had the approximate compositions, La2SbO1.7 and La2BiO1.9. Significant Sb site disorder is observed in these materials and a split site model was used to accommodate the Sb displacement within the a-b plane. For La2BiO1.9 the split site model does not appear to be as necessary. Both La2SbO1.7 and La2BiO1.9 display non metallic behavior with a small activation energy, contrary to the metallic behavior predicted by band theory. The behavior of Sb compositions can generally be understood to come from a pseudo gap opening up at the Fermi level due to significant disorder. Further neutron diffraction studies should attempt to better correlate, oxygen content, Sb site disorder, and physical properties across multiple 139 samples. Near term future studies should couple neutron diffraction, Sb Mössbauer spectroscopy, and magnetism of various Ln2SbO2-x materials to examine changes in the charge state of Sb; as their average charge state should be above Sb2- owing to the oxygen vacancy electron doping.

In all, these materials explorations have shown that new materials with interesting properties are waiting to be discovered. These studies comprise just a small fraction of work towards a much larger goal of deeply understanding the interactions of layered oxypnictides and solid state matter in general. It is hoped that the observations presented here will aid in future quests to design new materials, which in turn expand our knowledge and improve our lives. 140

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APPENDICES

148

Appendix A Rapid Microwave Synthesis of the Iron Arsenides NdFeAsO and NdFe0.9Co0.1AsO

A.1 Abstract The future of iron pnictide superconductors in technology is still undecided.

While these materials are now known to possess relatively high critical temperatures and critical magnetic fields, processing methods for these superconductors are still in the development stage. Recently we have been investigating possible ways to speed up the synthetic process for obtaining polycrystalline iron arsenide superconductors and other transition metal pnictides. Here we report the synthesis of NdFeAsO and

NdFe0.9Co0.1AsO in less than one hour total exposure to microwave radiation using an additional microwave susceptor to surround the reaction ampoule. Structure and property measurements reveal the samples to be of high quality and superconducting when Co doped.

149

A.2 Introduction It has been almost 4 years now since layered iron arsenide materials, such as

LaFeAsO and BaFe2As2, were reborn and a new surge of superconductor research began [1],[2]. At 4 years old, the iron arsenide superconductors are still young; however, they are off to a good start and proving themselves worthy of industrial application by virtue of possessing high critical temperatures and large critical magnetic fields [3]. The doped BaFe2As2 materials have been recently examined with bi-crystals and it was shown that grain boundaries are not as detrimental to superconductivity when compared with the cuprates [4]. Less demanding grain boundary limitations for arsenide superconductors may translate into lower cost wires and tapes, as less stringent manufacturing requirements increases the number of commercially viable synthetic pathways.

Microwave energy has been used to rapidly make polycrystalline samples of many different functional materials; including the high Tc superconductors

YBa2Cu3O7-x and Bi2Sr2CaCu2O8+x [5, 6]. However, there has yet to be a report on microwave synthesis of iron pnictide materials. Recently we have been working to extend the well known field of solid state microwave synthesis to transition metal pnictides such as skutterudites [7]. Here we report the synthesis of the iron arsenide compounds NdFeAsO and NdFe0.9Co0.1AsO, with microwave radiation and copper 150 , CuO, as an additional microwave susceptor. This is the first report of a microwave synthetic method for iron pnictides and one of the fastest to date.

We have been interested in rapid solid state microwave reactions primarily for two reasons: 1) the speed at which samples can be produced allows for more rapid trial and error analysis of potentially new and interesting compositions within the iron pnictide family 2) if high quality polycrystalline materials are needed, for sputter target manufacturing or powder in tube wire production, then a rapid microwave synthesis method may prove economical.

Many inorganic metal powders and metal oxides will heat extremely fast when exposed to 2.45 GHz radiation, the common frequency of domestic and research grade multimode microwave ovens. The thermal energy dissipated by these materials upon exposure can be harnessed to drive atomic diffusion and solid state reactions. If none of the reactants is a microwave susceptor than an additional susceptor may be used to surround the reaction vessel or may be placed in direct contact with the reactants, as long as there are no possible side reactions with the susceptor. For more general information on these techniques see the reviews and references by Rao et al. or Mingos and Baghurst [8] [9].

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A.3 Materials and Methods

Elemental Fe and As were used along with Nd and Nd2O3 as starting reagents; all were of 3N purity or better from Alfa Aesar. Stoichiometric amounts of the starting materials were ground and pressed into pellets (0.5 gramsin an Ar glove box.

The pellets were subsequently sealed inside fused silica ampoules (OD 11 mm, ID 9 mm, approx. 90 mm length) evacuated to approximately 1 x 10-4 torr. The sealed ampoule was then placed inside an alumina combustion boat crucible and buried within 110-115 grams of CuO. The reaction vessel was then surrounded by microwave transparent high-temperature fire bricks and placed in the microwave oven cavity, see Figure A.1.

Figure A.1 Schematic of the setup used to synthesize NdFeAsO and NdFe0.9Co0.1AsO issuing reaction ampoules surrounded by an additional susceptor within a 2.45 GHz multi-mode microwave cavity. 152

The microwave used was a CEM MDS-81D (2.45 GHz multi-mode oven). If continuously exposed to microwaves the temperature of the CuO can reach its melting point (1474 K) and cause softening of fused silica ampoules, so care must be taken to limit the maximum temperature. For this reason microwave exposure was limited to 5 minute intervals at 100% power (750 Watts). The temperature was estimated to be approximately 1170-1220 K by placing a thermocouple in the surrounding CuO immediately after exposure. Ten 5 minute heating cycles was sufficient to produce high quality samples of NdFeAsO and NdFe0.9Co0.1AsO. The total microwave exposure time was thus only 50 minutes. After microwave synthesis, powder x-ray diffraction (PXRD) patterns were collected from 20-60° 2θ, using a Rigaku MiniFlex 2, with stepsize of 0.02° 2θ, and 2 second dwell time.

Magnetization and resistivity data were collected with a Quantum Design PPMS.

A.4 Results and Discussion Recently a 20 minute low-temperature (1173 K) synthesis of iron arsenide superconductors such as SmFeAsO0.85F0.15 has been reported; however the reactants were ball-milled for 4 hours to enhance reaction kinetics [10]. It is interesting to note that even without ball milling the speed of microwave synthesis appears to be on par with that of the ball milled samples. Cobalt and iron powders are both excellent microwave susceptors and have been reported to achieve temperatures around 1000 153

K after about 3 - 7 minutes exposure [8]. It is possible that this coupling to the microwave energy allowed for enhanced interdiffusion of the Fe and Co; however, the small batch size used here meant that the energy dissipated by the Co and Fe reactants alone was not enough to drive the reaction. By using an additional susceptor we are in essence creating a miniature furnace around our sample with an extremely high ramp rate. The radiative heat from this susceptor material then becomes the main driving force for the reaction.

Due to their stoichiometry LnFePnO materials are often called the 1111-type iron arsenides, where Ln=Lanthanoid and Pn=Pnictide. At room temperature these materials adopt a tetragonal structure with P4/nmm (#129) symmetry. The structure can be described as alternating Nd2O2 and Fe2As2 layers stacked along the c-axis, wherein the former adopts a fluorite type arrangement (anion in tetrahedral coordination) and the later adopts an anti-fluorite coordination (metal cation in tetrahedral coordination). Figure A.2 details the 1111-type structure and c-axis stacking. Least squares refinement [11] of the PXRD data within the P4/nmm space group reveals the lattice parameters of NdFeAsO to be a=3.9612(9) Å and c=8.583(2)

Å. The lattice parameters of NdFe0.9Co0.1AsO are found to be a=3.9656(6) Å and c=8.554(1) Å. PXRD patterns for each are shown in Figure A.2; very small peaks identified as FeAs are labeled, (*). The refined unit cell lengths for both compositions agree well with previously reported lattice parameters [12]. The a-parameter of 154

NdFe0.9Co0.1AsO undergoes a slight expansion as Fe is replaced with Co while the c-parameter contracts more sharply. Structural refinements of the NdFe1-xCoxAsO system have shown that Co substitution is accompanied by a shift in the As z-position towards the Fe/Co layer and a decrease in the average M-As bond length [12].

Figure A.2 X-ray powder diffraction patterns for NdFeAsO and NdFe0.9Co0.1AsO synthesized using microwave radiation and an additional susceptor, CuO, surrounding the reaction ampoule. Shown right is the structure of Nd(Fe/Co)AsO materials depicted as layers of Nd2O2 and Fe2As2 stacked along the c-axis.

155

Magnetic susceptibility for the composition NdFe0.9Co0.1AsO, plotted as 4πχvol, indicates the onset critical temperature, Tc(χonset), is 16 K while χvol ≥ 1 at 2 K demonstrates that the superconductivity is of a bulk nature, Figure A.3.

Figure A.3 Magnetic susceptibility of NdFe0.9Co0.1AsO collected under zero field cooled conditions with a magnetic field of 20 Oe; shown plotted as 4πχvol. The critical temperature, Tc, as determined by the dχvol/dT plot is 16 K; shown overlaid for comparison

156

The critical temperature as determined by resistivity, Tc(ρonset) is found to be

18.4 K, when using the 90% normal state resistivity criterion. Both critical temperature values are in excellent agreement with those previously reported [12].

Figure A.4 shows the normalized resistivity, (ρ/ρ25K), for NdFe0.9Co0.1AsO from 5-35 K at various magnetic field strengths.

Figure A.4 Normalized resistivity (ρ /ρ25K) of NdFe0.9Co0.1AsO collected under magnetic fields from 0 - 6 Tesla (T). The inset details how the critical temperatures for 10%, 50%, and 90% of the normal state resistivity change with increasing magnetic field; the rate of the change for each is listed below the label. 157

The inset of Figure A.4 details the temperature of 10%, 50%, and 90% normal state resistance as a function of magnetic field strength. The slope, (dHc2/dT)T=Tc, is found to be -1.7, -2.7, -7.3 T/K for the 10%, 50%, and 90% normal state resistances respectively. The Werthhamer-Helfand-Hohenberg model can then be used to estimate the upper critical field at zero Kelvin as Hc2(0)=-0.693(Tc)(dHc2/dT)T=Tc. Using the 50% normal resistance value for (dHc2/dT)T=Tc the upper critical field is calculated to be 32 T.

A.5 Conclusion

In summary we have synthesized NdFeAsO and NdFe0.9Co0.1AsO with microwave radiation in conjunction with CuO as an additional microwave susceptor for the first time. The total exposure time needed to produce the samples is 50 minutes. Properties such as lattice parameters, Tc, and Hc2(0) agree well with those reported in the literature and show this method to be capable of producing high quality samples of iron pnictide superconductors. The ability for iron arsenides to make inroads to the industrial market place will depend on a large number of factors, one of them being ease of manufacturing. We believe microwaves offer a rapid and efficient means of producing bulk polycrystalline iron arsenide materials and as such will also aid in efforts to discover new analogs.

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A.6 References [1] Y. Kamihara, T. Watanabe, M. Hirano, H. Hosono, J. Am. Chem. Soc. 130 (2008) 3296-3297.

[2] M. Rotter, M. Tegel, D. Johrendt, Phys. Rev. Lett. 101 (2008) 107006-4.

[3] A. Gurevich, Nat. Mater. 10 (2011) 255-259.

[4] T. Katase, Y. Ishimaru, A. Tsukamoto, H. Hiramatsu, T. Kamiya, K. Tanabe, H. Hosono, Nat. Commun. 2 (2011) 409.

[5] D.R. Baghurst, A.M. Chippendale, D.M.P. Mingos, Nature 332 (1988) 311.

[6] M. Kato, K. Sakakibara, Y. Koike, Jpn. J. Appl. Phys. 38 (1999) 5867-5868.

[7] K. Biswas, S. Muir, M.A. Subramanian, Mat. Res. Bull. 46 (2011) 2288-2290.

[8] K.J. Rao, B. Vaidhyanathan, M. Ganguli, P.A. Ramakrishnan, Chem. Mater. 11 (1999) 882-895.

[9] D.M.P. Mingos, D.R. Baghurst, Chem. Soc. Rev. 20 (1991) 1.

[10] A.-H. Fang, F.-Q. Huang, X.-M. Xie, M.-H. Jiang, J. Am. Chem. Soc. 132 (2010) 3260-3261.

[11] N. Dragoe, J. Appl. Crystallogr. 34 (2001) 535-535.

[12] A. Marcinkova, D.A.M. Grist, I. Margiolaki, T.C. Hansen, S. Margadonna, J- W.G. Bos, Phys. Rev. B 81 (2010) 064511-10.

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Appendix B Celebrating superconductivity: A case study for developing high-impact science education programs

B.1 Abstract To celebrate the 100th anniversary of superconductivity and the 2011

International Year of Chemistry we created an educational program that involved classroom outreach, summer internships, free magnetic levitation demonstration kits for educators, and the development of a website for students and educators interested in learning more about superconductor materials. As a result of the demonstration kit delivery program alone it is estimated that over 2,000 students will be shown the superconductor demonstration. Through our experience we have identified four key components in the development of any high-impact science education program.

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B.2 Introduction 2011 has been a big year for chemistry. Throughout the year we celebrated a number of centennial anniversaries including the 100th anniversary of Mme. Curie’s

Nobel Prize in chemistry, and the 100th anniversary of the founding of the

International Association of Chemical Societies (now IUPAC). This year has also marked another important 100th anniversary; that of superconductivity, a physical phenomenon where electrons travel through a material without any resistance [1].

Superconductors have advanced numerous technologies, the most common are perhaps magnetic imaging (MRI) devices, in which superconducting wires are used to generate the magnetic field. In addition, 2011 marks the 25th anniversary of high temperature superconductivity and the discovery of the ceramic cuprate materials which have enabled superconductivity to be brought into the public and made tangible [2,3].

Over the course of 100 years, huge strides have been made; the breadth of known superconducting materials, the temperature range of properties, and comprehensive theories describing the phenomenon of superconductivity have all been greatly expanded. Chemists have contributed much to our knowledge in all of these areas (full disclosure, this article is written entirely by chemists); however, superconductivity is undoubtedly a highly collaborative field and heavily steeped in low-temperature physics. 161

It is superconductivity’s multifaceted nature which makes it an exciting and relevant tool for introducing many materials chemistry and science topics. Some of the stated goals for the 2011 International Year of Chemistry were to increase public appreciation of chemistry in meeting world needs, increase the interest of young people in chemistry, and to generate enthusiasm for the creative future of chemistry

[4]. The superconductor-magnetic-levitation demonstration can achieve all these goals. Basically, anytime is a good time to talk about superconductors; however, the

100th anniversary and the IYC 2011 added impetus and fortunate context.

Unfortunately, it seems educational opportunities involving the sciences of superconductors are often limited to undergraduate laboratories or even graduate level studies [5-13].

Early engagement with exciting science is particularly important when it comes to the discussion of America’s science education needs *14,15+. Students must be shown that being a materials scientist is an exciting and rewarding career choice before they begin to choose colleges and graduate from high school. The demonstration of magnetic levitation using high temperature superconductors is great way to excite students about the possibilities of materials science and chemistry [13]. Through superconductors topics such as stoichiometry and crystalline structure (chemistry), electricity and thermal energy transfer (physics), and materials and system design (engineering) can be easily introduced to middle 162 and high school students. We should not wait to inform students about one of the most amazing discoveries in human history and the importance of materials science and materials discovery today!

Figure B.1 A schematic detailing the four primary components of our superconductor outreach activities as part of the 2011 centennial anniversary of superconductivity.

With this in mind, we set out to design a high-impact superconductor education program but with small and easily achievable goals. It should also be stated that by high-impact we do not simply mean reaching the largest possible number of students, it is also about the quality of the educational experience. We therefore chose to develop a program that would involve one-on-one mentoring, 163 classroom teaching, as well as broad dissemination of information through demonstration kits distribution and website construction. A general overview of the program and its four primary components is shown in Figure B.2.

It should also be mentioned that plans for this education program were initially born from discussions within the Department of Chemistry Teaching Mentor

Program at Oregon State University (OSU). This mentor program offers graduate student interested in teaching the opportunity to discuss teaching philosophy and pedagogy with experienced educators of all types and most importantly to gain lecturing experience within the university classroom. Further details of the OSU mentoring program can be found at www.chemistry.oregonstate.edu/mentor_program.

B.3 Results and discussion The most substantial part of the program involved getting outside of the research lab and bringing superconductor and superconductor-related demonstrations to local schools. Before the end of the 2011 school year, we visited a local middle school and talked with two separate 8th grade science classes. Students at this level in Oregon often begin to study electricity and so superconductivity is an excellent addition to their ongoing scientific discussions. Based on student responses it was a definite hit which the students thought was informative and 164

“cool”. Some of the comments solicited anonymously from students after the presentation included the following statements.

“It is interesting to imagine the possibilities of a room-temperature superconductor”

“I started thinking about how they (superconductors) will affect the future”

“I enjoyed seeing it (superconductors) in person instead of reading from a text book.”

“I never expected to be able to ever have the experience of seeing a superconductor in action”

“The levitating magnet was really awesome because I had never seen anything like that before. It was exciting to think about a future with floating transportation.”

However, it is interesting to note that when asked for a constructive criticism, a number of students responded with comments like “how about a PowerPoint?” or

“could you give more time to play with the superconductors?” These types of responses are anecdotal evidence that these students were expecting a multimedia demonstration and that they are certainly hungry for more hands-on interaction with exciting scientific demonstrations. In addition, we also presented superconductor demonstrations to 6 freshman high school chemistry classes. Once again, according to the commentary of students and the teacher, the talk on materials chemistry, thermal energy, electronic conduction, and magnetic levitation was very well 165 received. In total, the demonstrations were given to approximately 240 middle and high school students by us directly.

As part of the centennial celebration, we also wanted to make a number of free superconductor demonstration kits available to educators who did not already have a superconductor levitation demonstration as part of their repertoire. We decided that rather than simply making the well-known YBa2Cu3O7 (YBCO) high- temperature-superconducting pellets ourselves, we should utilize the laboratory experience to enable high school interns and undergraduate researchers to make the

YBCO pellets. The established procedure offers opportunities to teach solid-state chemical processing, and strong lessons about stoichiometry due to the oxygen sensitivity of the structure. Students are also given a higher purpose to their activity in the means of returning science education to American children. Early on, we realized that the most efficient way to achieve this goal would be to work with existing summer research programs that match students with laboratories. Over the course of the summer Sean Muir and Mas Subramanian worked as mentors to a high school intern participating in the Saturday Academy - Apprenticeships in Science and

Engineering [16].

Through this internship, the student learned the basics of solid state chemistry and materials characterization. Specific materials explored include superconductors, phosphors, and pigments. The high school A.S.E. intern, along with 166 another undergraduate worked to produce a number of YBCO superconductor pellets, which were then used to create demonstration kits. In addition S.M. worked as a superconductor advisor to the National Science Foundation Rock Camp (a summer research program on solid state chemistry) [17]. As part of this program, undergraduate students and college educators investigate a number of interesting materials and the synthesis of YBCO serves a great platform for teaching traditional solid state synthesis techniques. As part of the 2011 celebration the YBCO pellets generated by Rock Camp students were then contributed to the demonstration kit program. Knowing that their efforts were helping to educator others was not only rewarding for interns and Rock Camp students, it allowed educational value to be fed back into the educational system.

The superconducting demonstration kits were designed to be as perpetually sustainable as possible. Each kit contains everything needed to do the magnetic levitation demo except for the liquid needed to cool the YBCO pellet. In total 15 kits have been mailed out and have been distributed to high schools and undergraduate educators. Specifically, the kit contained a 7 gram YBCO pellet, 2 small Nd:Fe:B magnets (one cylinder, one cube), a pair of non-magnetic nylon tweezers, a thermally- insulating Styrofoam dish to hold liquid nitrogen, and stickers to hand out to students containing a web address (www.solidchem.net, discussed below) linking them to more information, Figure B.2. 167

Figure B.2 Details of the free superconductor demonstration kits which were made available to educators as part of the 2011 centennial anniversary of superconductivity and the 25th anniversary of high temperature superconductor materials.

The schools where receiving educators are located range from large universities in urban areas such as Portland, OR, with class sizes of 200 students or more, to small high schools in rural Washington and Arizona counties with only 10-20 students per class. Educators answered a short list of questions about the classes for whom they would be doing the superconductor-levitation. These questions asked about the official science subject taught by the educator, the number of class periods they teach, and what their average class size is. Estimates based on the educators’ responses suggest that approximately 2,149 students (excluding our home institution) will be shown the superconductor demo. Perhaps the most exciting part is that every single educator answered that they had no previous experience with 168 doing the superconductor demonstration. This means that by offering free demonstration kits this program was able to reach students who would otherwise probably not have seen the demonstration and been introduced to the science of superconductors at school.

However, this last point brings up another issue. If the educators receiving kits have never done the superconductor demonstrations, presumably, they have little background knowledge of superconductor science. We wanted to offer educators a resource for learning about superconductors while also giving them a direction to guide students for more information about superconductors. To achieve this goal a website was built and designed to serve as a companion resource to educators who received the kits. The websites contains enough information and external links that an educator can put together notes/slides effectively and lead a discussion on superconductivity within the classroom, but it can also to serve as a calling card their students can later access and use to remember key concepts or learn more. The additional benefit of this medium is that the information is readily available to anyone with internet access, even if they did not receive a demo kit directly from us. This website can now be viewed at www.solidchem.net

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B.4 Conclusions While the research program presented here deals with the aptly-timed subject of superconductors, we feel that this is a program model which is well suited for many areas of science. We plan to expand our efforts in the near future to encompass materials such as phosphors, pigments, and even material-science- dependent devices such as thin-film transistors and thermoelectrics. The basic framework for helping any scientist develop a high impact outreach program has been laid out. The four defining characteristics are:

1) The program should have a component which involves public and/or student outreach. Scientists should use this opportunity to reflect on the fundamental concepts behind their research and how can these be presented in the most publically- engaging manner possible. It is important to engage students with the marvels of modern science in order to inspire them to dream larger.

2) The program should give students the opportunity to learn within the laboratory setting and gain hands on experience with science. This opportunity should be directly promoted to high school students during classroom or public demonstrations. However, it is important not to reinvent the wheel. If possible, work with internship programs that already exist. In this case, we worked with programs such as the Saturday Academy [16] and the NSF-Rock Camp [17]. 170

3) Part of the student researchers’ experience should come from creating something which can subsequently be used to educate more students; in this case, a superconductor pellet. It is the demonstration kits that we delivered to other educators which allowed a factor of ten more students outside of our campus community to be reached by our activities. Student interns take pride in the fact that the materials produced through their hard work and understanding of chemistry are used to teach others.

4) Students and educators are spending an ever increasing amount of time on the web. Therefore, creating a web presence for science is very important. The internet should be used for disseminating information among educators and to engage students with multiple forms of media including diagrams, videos, and animations. When going out to classrooms to do demonstrations, a website can serve as a great calling card and allow students to later access and recall bits of information. Additionally, a website is something that students can share with peers and family even if they did not actually see the demo or presentation. The website acts as a method of sharing an experience and cementing the memory in their minds.

B.5 References [1] H.K. Onnes, Comm. Phys. Lab. Leiden (1911).

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[17] Saturday Academy - Apprenticeships in Science and Engineering, http://www.saturdayacademy.org/ase/ (accessed Oct 2011).

[18] NSF Summer Research Program in Solid State and Materials Chemistry, http://ssmchem.uoregon.edu/ (accessed Oct 2011).