Development of High-efficiency Metal Chalcogenide ...398523/s42210359_final... · 1 Development of...

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Development of High-efficiency Metal Chalcogenide

Thermoelectric Nanomaterials by Nanostructure Engineering

Lei Yang

Master of Philosophy in the Field of Nanotechnology

A thesis submitted for the degree of Doctor of Philosophy at

The University of Queensland in 2016

School of Mechanical and Mining Engineering

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Abstract

The development of high performance thermoelectric materials, which can

directly convert heat into electricity, is becoming an alternative to overcome

the global energy shortage. The efficiency of thermoelectric materials is

determined by the dimensionless figure-of-merit ZT=S2σT/κ, where S is the

Seebeck coefficient, σ is the electrical conductivity, T is the absolute

temperature, and κ is the thermal conductivity includes the contribution from

the electron (κe) and lattice (κL) components. High efficient thermoelectric

materials can be extensively applied as alternative energy sources in many

fields such as waste heat recycling, solid state power generation, and

refrigeration.

The challenge of developing high performance thermoelectric materials is how

to enhance the ZT value by optimizing the conflict and interdependent

parameters (S, σ and κ). Metal chalcogenides have been targeted in this

thesis because they intrinsically have good electrical transport properties

(high S and σ) and low κ. Among them, Bi2Te3, PbTe, and Cu2Se have the

highest ZT in room temperature, intermediate (500-800 K) and higher

temperature (~1000 K) range, respectively. To further enhance their

thermoelectric performances, nanostructure engineering has been applied in

this thesis. Cu2Se, Bi2Te3 and PbTe nanomaterials have been synthesized via

facile and controllable solvothermal methods; their structures and

thermoelectric properties were extensively investigated.

Nanostructured β-phase Cu2Se materials were synthesized and sintered

using spark plasma sintering process. The nano-sized grains were preserved

after the sintering, leading to high density of small angle grain boundaries

accommodated by defects, which significantly reduced the κL of as-prepared

samples but did not affect the electrical transport properties, resulting an

outstanding ZT of 1.82 at around 850 K. Via a controlled synthesis approach,

Cu-deficient Cu2-xSe nanomaterials were obtained. The high degree of Cu

deficiency was found to trigger the phase transition from β- to α-phase,

leading to small amount of α-phase in the Cu1.95Se sample. The Cu deficiency

was proved to harm the thermoelectric performance of Cu2-xSe nanomaterials

via increasing the carrier concentration, and leading to a significantly reduced

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S. Tellurium was doped into Cu2Se nanomaterials to modify the electrical

transport properties. The effects of Te doping to the Cu2Se nanomaterials

were carefully studied, the Cu2Se0.99Te0.01 sample was found to have the

highest S among all the Te-doped samples and the ZT of Cu2Se0.98Te0.02 has

the highest peak ZT ~1.76.

The developed nanostructure engineering was approved to be effective on

Bi2Te3 and PbTe nanomaterials. Pure Bi2Te3 hexagonal nanoplates were

synthesized and sintered. High density of small angle grain boundaries

accommodated by defects were also found in the sintered Bi2Te3

nanomaterials, which significantly reduced the κ and resulted in an improved

ZT ~0.88 at 400 K. The Bi-doped PbTe nanocubes were obtained, and the

doping of Bi was confirmed via multiple technologies. The high density of

grain boundaries and the Bi dopant effectively reduced the κ. Also, the Bi

dopants improved the electrical transport properties of PbTe, finally leading to

enhanced ZT.

In this thesis, reliable, facile and controllable solvothermal methods were

developed to obtain metal chalcogenides-based nanomaterials. By applying

nanostructure engineering, the enhancement of thermoelectric performances

for metal chalcogenides-based nanomaterials has been achieved.

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Declaration by author

This thesis is composed of my original work, and contains no material

previously published or written by another person except where due reference

has been made in the text. I have clearly stated the contribution by others to

jointly-authored works that I have included in my thesis.

I have clearly stated the contribution of others to my thesis as a whole,

including statistical assistance, survey design, data analysis, significant

technical procedures, professional editorial advice, and any other original

research work used or reported in my thesis. The content of my thesis is the

result of work I have carried out since the commencement of my research

higher degree candidature and does not include a substantial part of work that

has been submitted to qualify for the award of any other degree or diploma in

any university or other tertiary institution. I have clearly stated which parts of

my thesis, if any, have been submitted to qualify for another award.

I acknowledge that an electronic copy of my thesis must be lodged with the

University Library and, subject to the policy and procedures of The University

of Queensland, the thesis be made available for research and study in

accordance with the Copyright Act 1968 unless a period of embargo has been

approved by the Dean of the Graduate School.

I acknowledge that copyright of all material contained in my thesis resides

with the copyright holder(s) of that material. Where appropriate I have

obtained copyright permission from the copyright holder to reproduce material

in this thesis.

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Publications during candidature

Peer-reviewed papers:

1. Yang, L.; Chen, Z. G.; Han, G.; Hong, M.; Huang, L.; Zou, J., Te-doped

Cu2Se nanoplates with high average thermoelectric figure of merit.

Journal of Materials Chemistry A, 2016, 4, 9213-9219.

2. Yang, L.; Chen, Z.-G.; Han, G.; Hong, M.; Zou, J., Impacts of Cu

deficiency on the thermoelectric properties of Cu2−XSe nanoplates.

Acta Materialia, 2016, 113, 140-146.

3. Yang, L.; Chen, Z.-G.; Nie, T.; Han, G.; Zhang, Z.; Hong, M.; Wang,

KL.; Zou, J., Co-doped Sb2Te3 Paramagnetic Nanoplates. Journal of

Materials Chemistry C, 2016, 4, 521-525.

4. Yang, L.; Chen, Z.-G.; Hong, M.; Han, G.; Zou, J., Enhanced

Thermoelectric Performance of Nanostructured Bi2Te3 through

Significant Phonon Scattering. ACS Applied Materials & Interfaces,

2015, 7, 23694-23699.

5. Yang, L.; Chen, Z.-G.; Han, G.; Hong, M.; Zou, Y.; Zou, J., High-

Performance Thermoelectric Cu2Se Nanoplates through Nanostructure

Engineering, Nano Energy, 2015, 16, 367-374.

6. Yang, L.; Chen, Z.-G.; Han, G.; Cheng, L.; Xu, H.; Zou, J., Trifold

Tellurium One-Dimensional Nanostructures and Their Formation

Mechanism. Crystal Growth & Design, 2013, 4796–4802.

7. Hong, M.; Chen, Z. G.; Yang, L.; Zou, J., Enhancing Thermoelectric

Performance of Bi2Te3-based Nanostructures through Rational

Structure Design. Nanoscale, 2016. DOI: 10.1039/C6NR00719H.

8. Hong, M.; Chen, Z. G.; Yang, L.; Zou, J., BixSb2−xTe3 nanoplates with

enhanced thermoelectric performance due to sufficiently decoupled

electronic transport properties and strong wide-frequency phonon

scatterings. Nano Energy, 2016, 20, 144-155.

9. Chen, Z.-G.; Zhang, C.; Zou, Y.; Zhang, E.; Yang, L.; Hong, M.; Xiu, F.;

Zou, J., Scalable Growth of High Mobility Dirac Semimetal Cd3As2

Microbelts. Nano Letters, 2015, 15, 5830−5834.

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10. Hong, M.; Chen, Z.-G.; Yang, L.; Han, G.; Zou, J., Enhanced

Thermoelectric Performance of Ultrathin Bi2Se3 Nanosheets through

Thickness Control. Advanced Electronic Materials, 2015, 1500025, 1-9.

11. Han, G.; Chen, Z.-G.; Ye, D.; Wang, B.; Yang, L.; Zou, Y.; Wang, L.;

Drennan, J.; Zou, J., In3Se4 and S-doped In3Se4 nano/micro-structures

as new anode materials for Li-ion batteries. The Journal of Materials

Chemistry A, 2015, 3, 7560–7567.

12. Chen, Z.-G.; Yang, L.; Ma, S.; Cheng, L.; Han, G.; Zhang, Z.-d.; Zou, J.,

Paramagnetic Cu-doped Bi2Te3 nanoplates. Applied Physics Letters,

2014, 104, 053105.

13. Han, G.; Chen, Z.-G.; Yang, L.; Hong, M.; Drennan, J.; Zou, J.,

Rational Design of Bi2Te3 Polycrystalline Whiskers for Thermoelectric

Applications. ACS Applied Materials & Interfaces, 2014, 7, 989-995.

14. Zou, Y.; Chen, Z.-G.; Huang, Y.; Yang, L.; Drennan, J.; Zou, J.,

Anisotropic Electrical Properties from Vapor–Solid–Solid Grown Bi2Se3

Nanoribbons and Nanowires. The Journal of Physical Chemistry C,

2014, 20620–20626.

15. Han, G.; Chen, Z. G.; Ye, D.; Yang, L.; Wang, L.; Drennan, J.; Zou, J.,

In-Doped Bi2Se3 Hierarchical Nanostructures as Anode Materials for Li-

Ion Batteries. The Journal of Materials Chemistry A, 2014, 7109-7116.

16. Liao, Z.-M.; Chen, Z.-G.; Lu, Z.-Y.; Xu, H.-Y.; Guo, Y.-N.; Sun, W.;

Zhang, Z.; Yang, L.; Chen, P.-P.; Lu, W.; Zou, J., Au impact on GaAs

epitaxial growth on GaAs (111)(B) substrates in molecular beam

epitaxy. Applied Physics Letters, 2013, 102 (6).

17. Han, G.; Chen, Z.-G.; Yang, L.; Cheng, L.; Jack, K.; Drennan, J.; Zou,

J., Thermal stability and oxidation of layer-structured rhombohedral

In3Se4 nanostructures. Applied Physics Letters, 2013, 103 (26).

18. Han, G.; Chen, Z.-G.; Yang, L.; Cheng, L.; Drennan, J.; Zou, J., Phase

Control and Formation Mechanism of New-Phase Layer-Structured

Rhombohedral In3Se4 Hierarchical Nanostructures. Crystal Growth &

Design, 2013, 5092–5099.

19. Han, G.; Chen, Z. G.; Sun, c.; Yang, L.; Cheng, L.; Li, Z.; Lu, W.; Gibbs,

Z. M.; Snyder, J.; Jack, K. S.; Drennan, J.; Zou, J., A New Crystal:

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Layer-Structured Rhombohedral In3Se4. Crystal Engineering

Communications, 2013, 393-398.

20. Cheng, L.; Chen, Z. G.; Yang, L.; Han, G.; Xu, H. Y.; Snyder, G. J.; Lu,

G. Q.; Zou, J., T-Shaped Bi2Te3–Te Heteronanojunctions: Epitaxial

Growth, Structural Modeling, and Thermoelectric Properties. The

Journal of Physical Chemistry C, 2013, 12458-12464.

21. Cheng, L.; Chen, Z. G.; Ma, S.; Zhang, Z. D.; Wang, Y.; Xu, H. Y.;

Yang, L.; Han, G.; Jack, K.; Lu, G. Q.; Zou, J., High Curie Temperature

Bi1.85Mn0.15Te3 Nanoplates. Journal of the American Chemical Society,

2012, 134 (46), 18920-18923.

22. Chen, Z. G.; Han, G.; Yang, L.; Cheng, L.; Zou, J., Nanostructured

thermoelectric materials: Current research and future challenge.

Progress in Natural Science: Materials International, 2012, 22 (6), 535-

549.

Conference Papers:

1. Liao, Z. M.; Xu, H. Y.; Sun, W.; Guo, Y. N.; Yang, L.; Chen, Z. G.;

Zou, J.; Lu, Z. Y.; Chen, P. P.; Lu, W., Effects of Au catalyst on GaAs

(111)B surface during annealing, on Optoelectronic and

Microelectronic Materials & Devices. COMMAD, 2012, 7-8.

Conference proceedings:

1. Yang, L., Chen, Z.G., Han, G., Cheng, L.N., Zou, j., Synthesis and

morphological modification of Te nanowires via a simple solvothermal

method. ACMM 22 / APMC 10 / ICONN 2012 Perth. 2012. (Oral

presentation)

2. Chen, Z.G., Yang, L., Han, G., Cheng, L.N., Zou, j., Development of

Thermoelectric Nanomaterials by a facile Solvothermal Method.

ACMM 22/ APMC 10 /ICONN 2012 Perth. 2012 (Oral presentation)

3. Han, G., Chen, Z.G., Yang, L., Cheng, L.N., Drennan, J., Zou, J., The

Solvothermal Synthesis of Indium Selenide Flowerlike Nanostructures.

ACMM 22/ APMC 10 /ICONN 2012 Perth. 2012. (Oral presentation)

4. Yang, L., Chen, Z.G., Han, G., Cheng, L.N., Zou, j., Understanding of

The Growth Mechanism of Tri-fold Tellurium Nanowires. UQ EAIT

Postgraduate Conference. Australia. 2012. (Oral presentation)

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5. Yang, L., Chen, Z.G., Han, G., Zou, j., Bi-doped PbTe Trifold

Nanostructures with improved thermoelectric performance. ICAMP8.

Gold Coast, Australia. 2014. (Oral presentation)

6. Yang, L., Chen, Z.G., Han, G., Zou, j., High-Performance

Thermoelectric Cu2Se Nanoplates through Nanostructure Engineering.

C-MRS. Guiyang, China. 2015. (Oral presentation)

7. Yang, L., Chen, Z.G., Zou, j., Nanostructure Engineering on Cu2Se-

based Thermoelectric Materials. The 35th International Conference &

The 1st Asian Conference on Thermoelectrics in Wuhan, China, 29th

May- 2nd June, 2016. (Oral presentation)

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Publications included in this thesis

Yang, L.; Chen, Z.-G.; Han, G.; Hong, M.; Zou, J., Impacts of Cu deficiency

on the thermoelectric properties of Cu2−XSe nanoplates. Acta Materialia 2016,

113, 140-146. – incorporated as Chapter 4, part 4.2.1

Contributor Statement of contribution

Lei Yang (Candidate) Carried out sample synthesis (90%)

Designed experiments (60%)

Wrote the paper (60%)

Carried out characterization (90%)

Carried out data analysis (60%)

Zhi-Gang Chen Designed experiments (40%)

Wrote and edited paper (10%)

Supervised the project (60%)


Guang Han Carried out sample synthesis (10%)

Min Hong Carried out characterization (10%)

Jin Zou Wrote and edited paper (30%)



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Yang, L.; Chen, Z. G.; Han, G.; Hong, M.; Huang, L.; Zou, J., Te-doped

Cu2Se nanoplates with high average thermoelectric figure of merit. Journal of

Materials Chemistry A 2016, 4, 9213-9219. – incorporated as Chapter 4, part

4.2.2













Liqing Huang Carried out sample synthesis (10%)





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Yang, L.; Chen, Z.-G.; Han, G.; Hong, M.; Zou, Y.; Zou, J., High-performance

thermoelectric Cu2Se nanoplates through nanostructure engineering. Nano

Energy 2015, 16, 367-374. – incorporated as Chapter 5, part 5.2.1













Yichao Zou Carried out data analysis (5%)




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Yang, L.; Chen, Z.-G.; Hong, M.; Han, G.; Zou, J., Enhanced Thermoelectric

Performance of Nanostructured Bi2Te3 through Significant Phonon Scattering.

ACS Appl. Mat. Interfaces 2015, 7, 23694-23699. – incorporated as Chapter

5, part 5.2.2











Min Hong Carried out sample synthesis (10%)

Guang Han Carried out characterization (10%)




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Contributions by others to the thesis

No contributions by others.

Statement of parts of the thesis submitted to qualify for the award of

another degree

None

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Acknowledgements

Firstly, I sincerely acknowledge my supervisors, Dr Zhi-Gang Chen and

Professor Jin Zou for their careful and kind guidance to me in every aspect of

my PhD study in the University of Queensland. I admire their impressing

excellence as scientific researchers: well-rounded and profound knowledge,

determination, great patience, passion and the seriousness. They devote

themselves to pursue and share the scientific truth, also help many students

like me to develop the experimental skills and the ability to accomplish

scientific research projects. They have spent huge effort to teach me how to

design experiments and write scientific papers. It is a great honour for me to

join this family-like group under their supervisions.

Also, I would like to thank my colleagues, Dr Yong Wang, Mrs Ya Wang, Dr

Yang Huang, Dr Jing Lin, Dr Yanan Guo, Dr Lina Cheng, Dr Hongyi (Justin)

Xu, Dr Guang Han, Dr Wen Sun, Dr Zhi Zhang, Dr Kun Zheng, Dr Lihua

Wang, Zhiming Liao, Min Hong, Yichao Zou, Mun Soo, Chen Zhou, Liqing

Huang, in our research group at The University of Queensland. Thank you

very much for sharing your knowledge and skills, providing great help and

support in both of my study and life in Brisbane.

I acknowledge all the staff of Centre for Microscopy and Microanalysis (CMM)

at the University of Queensland for their technical support. I have learned

many useful analysing skills from them. And I acknowledge the financial

support from the China Scholarship Council for providing my PhD stipend.

Last but not least, I want to give my great thanks to my family: my dad, my

mum, and my elder sister, my wife Zilin Zhang, and all my friends in China

and Australia. Their endless love and encouragement are always the great

power in my life.

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Keywords

Thermoelectric materials, high performance, metal chalcogenides,

nanostructure engineering, solvothermal synthesis.

Australian and New Zealand Standard Research Classifications

(ANZSRC)

ANZSRC code: 091205, Functional Materials, 50%

ANZSRC code: 100706, Nanofabrication, Growth and Self Assembly, 40%

ANZSRC code: 100712, Nanoscale Characterisation, 10%

Fields of Research (FoR) Classification

FoR code: 0912, Materials Engineering, 50%

FoR code: 1007, Nanotechnology, 50%

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Table of Contents Chapter 1: Introduction ........................................................................................... - 1 -

1.1 Background .................................................................................................... - 1 -

1.2 Objective and Scopes ..................................................................................... - 3 -

1.3 Thesis outline ................................................................................................. - 4 -

Chapter 2: Literature Review .................................................................................. - 6 -

2.1 Thermoelectric Effects .................................................................................... - 6 -

2.1.1 Seebeck Effect........................................................................................ - 7 -

2.1.2 Peltier Effect ........................................................................................... - 7 -

2.1.3 Thomson Effect....................................................................................... - 7 -

2.1.4 Thermoelectric Generation and Refrigeration and the Figure-of-Merit .... - 8 -

2.1.5 Effective Factors ................................................................................... - 10 -

2.2 Development of Thermoelectric Materials .............................................. - 12 -

2.2.1 Thermoelectric Alloys ........................................................................... - 13 -

2.2.2 Materials of Complex Structures ........................................................... - 14 -

2.2.3 Metal Chalcogenides Thermoelectric Materials .................................... - 15 -

2.3 Principles and Methodologies to Achieve High ZT ....................................... - 26 -

2.3.1 Optimize the Carrier Concentration ...................................................... - 26 -

2.3.2 Band Engineering ................................................................................. - 28 -

2.3.3 Nanostructure Engineering ................................................................... - 30 -

2.4 Unsolved Issues and Opportunities .............................................................. - 33 -

2.5 Summary ...................................................................................................... - 34 -

Chapter 3 Methodologies ........................................................................................................ - 51 -

3.1 Synthesis Methods ....................................................................................... - 51 -

3.2 Characterization Methods ............................................................................. - 54 -

3.2.1 X-Ray Diffraction (XRD) ........................................................................ - 54 -

3.2.2 Scanning Electron Microscopy (SEM) .................................................. - 55 -

3.2.3 Transmission Electron Microscopy (TEM) ............................................ - 57 -

3.3 Thermoelectric Measurements ..................................................................... - 59 -

3.3.1 Thermal Properties ............................................................................... - 59 -

3.3.2 Seebeck Coefficient .............................................................................. - 60 -

3.3.3 Electrical properties .............................................................................. - 61 -

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Chapter 4 Controllable Synthesis of Metal Chalcogenides Nanostructures and Their

Thermoelectric Performances ........................................................................... - 66 -

4.1 Introduction ................................................................................................... - 66 -

4.2 Manuscripts for Publication........................................................................... - 66 -

4.2.1 Effects of Cu Deficiency on the Thermoelectric Properties of Cu2-XSe

Nanostructures .............................................................................................. - 67 -

4.2.2 Te-induced Phase Transition of Cu2SexTe1-x Nanomaterials and Their

Thermoelectric Properties .............................................................................. - 87 -

Chapter 5 Enhanced Thermoelectric Performances of Metal Chalcogenides via

Nanostructure Engineering ................................................................ - 107 -

5.1 Introduction ................................................................................................. - 107 -

5.2 Journal Publications and Manuscript .......................................................... - 107 -

5.2.1 High-Performance Thermoelectric Cu 2Se Nanoplates through

Nanostructure Engineering .......................................................................... - 108 -

5.2.2 Enhanced Thermoelectric Performance of Nanostructured Bi2Te3 through

Significant Phonon Scattering ...................................................................... - 130 -

5.2.3 Manuscript .......................................................................................... - 149 -

Chapter 6 Conclusions and Recommendations ................................................. - 171 -

- 1 -

List of Figures

Figure 1.1 Global energy supply (http://aleklett.wordpress.com).

Figure 1.2 ZT values of different materials as a function of temperature.

Figure 2.1 (a) Schematic diagram shows the thermoelectric process; (b) typical

thermoelectric module.

Figure 2.2 Efficiency of thermoelectric devices as a function of ΔT.

Figure 2.3 The relation between carrier concentration and the value of ZT.

Figure 2.4 Figure of merit (ZT) as a function of temperature for several high-

efficiency bulk thermoelectric materials

(http://chemgroups.northwestern.edu/kanatzidis/greatthermo.html).

Figure 2.5.(a-c) Crystal structure of CoSb3 revealing the large voids with rattlers, the

type-I clathrate Na8Si46, and β-Zn4Sb3; (d) ZT as a function of temperature for

skutterudites as thermoelectric materials; (e) Variable temperature ZT of clathrates,

and β-Zn4Sb3.

Figure 2.6 Schematics shows the structures and phase transition of Cu2-xSe

between α- and β-phase.

Figure 2.7 Rhombohedral crystal structure of Bi2Te3.

Figure 2.8 Valence band structure of PbTe1-xSex. (a) Brillouin zone showing the low

degeneracy hole pockets (orange) centred at the L point, and the high degeneracy

hole pockets (blue) along the Σ line. (b) Relative energy of the valence bands in

PbTe0.85Se0.15. At 500K the two valence bands converge, resulting in transport

contributions from both the L and Σ bands.

Figure 2.9 (a), (b) Experimental power factors as a function of Hall carrier

concentration; (c), (d) calculated power factors as a function of Hall carrier

concentration and (e), (f) calculated power factors as a function of reduced Fermi

level for Pb1-xLaxTe and PbIxTe1-x, respectively.

Figure 2.10 Schematic illustrates the resonant level on the electronic density of

states (DOS).

Figure 2.11 Schematic diagram of phonon scattering mechaism and electronic

transport within a thermalelectric material.

Figure 2.12 Calculated results for n-type Si80Ge20 show how the carriers with

different energy contribute to the Seebeck coefficient.

http://aleklett.wordpress.com/

http://chemgroups.northwestern.edu/kanatzidis/greatthermo.html

- 2 -

Figure 3.1 An autoclave used in solvothermal synthesis.

Figure 3.2 The process of a typical solvothermal route.

Figure 3.3 (a) The principle of XRD (http://www.pic2fly.com); (b) A XRD

spectrophotometer (http://analyticalinstrumentengineer.com).

Figure 3.4 (a) Electron beam penetrates into the sample and generate different

signals with different information (http://mee-inc.com); (b) Changes in the interaction

volume with topography (http://www.ammrf.org.au).

Figure 3.5 (a) The schematic outline of a TEM (http://www.hk-phy.org); (b) TECNAI

F20 TEM (http://www.bo.imm.cnr.it).

Figure 3.6 The principle of imaging and diffraction in TEM

(http://www.microscopy.ethz.ch).

Figure 3.7 Schematic diagram shows the working principle of the Netzsch DSC 404

F3: (a) the furnace and (b) the calculation of Cp using obtained signals.

Figure 3.8 Schematic diagram shows the working principle of the Netzsch LFA 457

system (from http://ap.netzschcdn.com).

Figure 3.9 Schematic diagram shows the working principle of Seebeck coefficient

measurement.

Figure 3.10 Schematics of Van der Pauw technique (http://www-

personal.umich.edu).

Figure 3.11 The measurement of Hall coefficient using Van der Pauw technique

under a reversible magnetic field

(http://en.wikipedia.org/wiki/File:Van_der_Pauw_Method_-_Hall_Effect.png).

http://www.pic2fly.com/

http://www.ammrf.org.au/

http://www.hk-phy.org/

http://www.bo.imm.cnr.it/

http://www.microscopy.ethz.ch/

http://ap.netzschcdn.com/

http://en.wikipedia.org/wiki/File:Van_der_Pauw_Method_-_Hall_Effect.png

- 3 -

List of Tables

Table 1 Thermoelectric properties of advanced Bi2Te3-based materials

Table 2 Advanced thermoelectric materials in Pb-Te system

Table 3 Common synthesis methods of low dimensional materials

- 4 -

List of Abbreviations RTGs: radioisotope thermoelectric generators

PGEC: phonon-glass/electron-crystal

FCC: face-center-cubic

COP: coefficient of performance

MS: melt spinning

HP: hot pressing

HEBM: high energy ball milling

SPS: spark plasma sintering

SP: solution phase method

MW: microwave synthesis

CC: cold compaction

NASA: the national aeronautics and space administration

MBE: molecular beam epitaxy

CVD: chemical vapour deposition

XRD: X-ray diffraction

SEM: scanning electron microscope

TEM: transmission electron microscope

EDS: energy dispersive X-ray spectroscopy

JCPDS: the joint committee on powder diffraction standards

HRTEM: high-resolution TEM

SAED: selected area electron diffraction

FFT: fast Fourier transform

EPMA: electron probe micro analysis

FIB: focused ion beam

XPS: X-ray photoelectron spectroscopy

- 1 -

Chapter 1: Introduction

1.1 Background

As the exhausting of fossil fuels, most countries are facing the shortage of energy.

Meanwhile, the environmental deterioration leads the requirement to clean energy.

Much attention has been paid on searching alternative energies to get through the

energy crisis. Solar cells, wind driven generators, and nuclear power plants have

been developed for several decades, which show the reliable quality and a

sustainable future of power generation. However, as it can be seen in Figure 1.1,

heat engines which consume fossil fuels to generate power, still provide about 90%

of the power requirement.1 But even the most efficient engines or factory systems

waste almost 70% of the primary energy,2 most of the energy was emitted with

exhausted gas (500-800K) or taken by the cooling systems. According to the total

consumption of energy in the world, 15TW1 of energy was wasted in the whole world

just to heat the environment. Fortunately, thermoelectric generation systems offer us

an appealing option that thermal energy could be transformed to electrical power.

The most attractive factor of the application of thermoelectric generators is that the

sources of heat exist at everywhere around the earth. By using thermoelectric

materials, most of these waste heats are expected to be converted into electricity.

- 2 -

Figure 1.1 Global energy supply (http://aleklett.wordpress.com).

Thermoelectric materials can be used as power generators and refrigerators.

Thermoelectric generators and refrigerators are silent and reliable because they are

solid-state devices without any moving parts. Actually, thermoelectric is not a novel

field. Radioisotope thermoelectric generators (RTGs) have been used to supply

power for many space missions such as Apollo lunar mission.3 Thermoelectric

materials has also been assembled to form devices which can convert waste heat

from exhaust gas of automobiles and factories, or be used as refrigerators for

cooling computers, infrared detectors, electronics and other equipment.4

With the increasing interest in thermoelectric applications, scientists have paid more

attention to find novel high-performance thermoelectric materials. Thermoelectric

materials has rapidly developed after the establishing of the basic science of

thermoelectric at 1950s.5 One of the most commercial TE material would behave as

a “phonon-glass/electron-crystal” (PGEC),6 that is, it would have the electrical

properties of a crystalline material and the thermal properties of an amorphous or

glass-like material.7 Bulk materials with the crystal structures of Skutterudites,

clathrates, complex alloys, chalcogenides, and oxides are identified as good

thermoelectric materials (shown in Figure 1.2).2, 8 It is worth to develop these

materials as well as to find new material systems. With the development of

nanotechnology in 1990s, an increasing number of researchers realized that

- 3 -

nanostructures can boost the thermoelectric properties of many kinds of materials,5,

9, 10 which offered a new approach to achieve higher thermoelectric performance. It is

believed that current thermoelectric science could be further developed and

advanced high efficient thermoelectric materials will play a crucial role of power

generation and refrigeration in near future.

Figure 1.2 ZT values of different materials as a function of temperature.7

1.2 Objective and Scopes

To carry out this project, general understanding of the theories and principles of

thermolelectrics should be established to find out the vital factors which affect the

thermoelectric performance of materials. The main goal of this project is to develop a

low-cost, efficient, reliable and environmental friendly synthesis method of as-

designed high-efficiency thermoelectric materials. Specifically, in this project, metal

chalcogenides-based system with relatively superior thermoelectric properties will be

synthesized, characterized and analysed by appropriate methods.

- 4 -

1.3 Thesis outline

To develop high performance metal chalcogenides-based thermoelectric

nanomaterials, Cu2Se, Bi2Te3 and PbTe-based nanomaterials have been

synthesized via facile and controllable solvothermal methods. The structures of the

as-prepared products were carefully investigated via advanced electron microscopy

and other methods. The products were then sintered and their thermoelectric

properties were analysed.

Chapter 1 is the introduction of this project.

Chapter 2 provides the literature review of previous studies about thermoelectric

materials. Some basic thermoelectric effects are introduced to help establishing the

understanding of effective factors for high performance thermoelectric materials. The

development of thermoelectric materials has been summarized. Then metal

chalcogenides thermoelectric candidates are focused due to their intrinsic high

thermoelectric performances and great potential on commercialization.

Chapter 3 introduces the relative experimental technologies used in this project.

Several popular synthesizing methods for nanomaterials were listed and compared

with each other. Solvothermal method is introduced in detail as the synthesis method

used in this project. Then some basic working principles of these experimental

technologies are introduced.

Chapter 4 demonstrates the controllable synthesis of metal chalcogenides

nanostructures and their thermoelectric performances. In this chapter, compositional

control of Cu2-xSe nanostructures and their thermoelectric properties, and phase

control of Cu2-xSe nanostructures triggered by Te-doping and their thermoelectric

properties have been carefully studied. This part includes two drafted manuscripts,

which have been submitted.

Chapter 5 focuses on enhanced thermoelectric performances of metal

chalcogenides via nanostructure engineering.

In the first part, nanostructure engineering was applied to enhance the thermoelectric

performance of Cu2Se nanomaterials. This part is included as the Nano Energy,

2015, 16, 367-374.

- 5 -

The second part confirms the effectiveness of nanostructure engineering on Bi2Te3

nanomaterials. This part is included as ACS Applied Materials & Interfaces, 2015, 7

(42), 23694-23699.

The third part combines the doping and nanostructure engineering on PbTe,

carefully analysed the growth mechanism and the thermoelectric properties, which

has been drafted and ready to be submitted.

Chapter 6 gives the conclusion of this thesis with recommendations for future

development of thermoelectric materials.

- 6 -

Chapter 2: Literature Review

In this chapter, fundamental thermoelectric effects and phenomena are presented to

address the basic factors which determine the thermoelectric performance of

materials. The conflict factors drawback the development of thermoelectric materials.

The approaches which can optimize the thermoelectric properties are then

demonstrated: to control the effective factors and choose appropriate materials

systems. Tellurium-based materials are focused due to their promising properties;

tellurium and lead telluride are studied in this project. The state-of-art researches

about tellurium-based materials are shown in this chapter.

2.1 Thermoelectric Effects

The conversion between temperature gradient and electricity was named as

thermoelectric effects. A series of theories have been developed by scientists to

describe and study thermoelectrics since the first thermoelectric effect was found by

Thomas Seebeck. The general principles of thermoelectrics are followed.

- 7 -

2.1.1 Seebeck Effect

In 1821-1823, Seebeck discovered the deflection of a compass needle which had

been placed in the vicinity of a close loop formed from two dissimilar conductors

when he heated one of the junctions,11 in other words, the difference of the

temperature leads the generating of current flow (Figure 2.1a). This phenomenon

was described as S=V/ΔT, which combined the Seebeck coefficient (S, sometimes

α), the voltage (V) and the temperature deference (ΔT).11 The Seebeck coefficient is

measured in V/K or μV/K. Thermoelectric power-generation devices which convert

thermal energy directly into electricity have been developed based on the discovery

of Seebeck effect.

2.1.2 Peltier Effect

Peltier observed temperature changes of the junction between two dissimilar

conductors when he input a current (I) through the circuit.11, 12 the Peltier effect was

explained by Lenz that heat (Q) is generated or absorbed at the junction between

two dissimilar conductors depending on the direction of the current, which is the

basis for thermoelectric refrigeration and could be described as:

Q=ΠI (1)

where Π is named as the Peltier coefficient, Q is the heat flow and I is the current.

2.1.3 Thomson Effect

Thomson (Lord Kelvin) established the relationship between Seebeck coefficient and

Peltier coefficient as: ST=Π, and predicted the third thermoelectric effect, Thomson

effect, that heat is absorbed or generated when a current passes along a single

homogeneous conductor with temperature gradient.

- 8 -

Figure 2.1 (a) Schematic diagram shows the thermoelectric process; (b) a typical

thermoelectric module.

2.1.4 Thermoelectric Generation and Refrigeration and the Figure-of-Merit

Thermoelectric couples which composed of n-type (electron carriers) and p-type

(hole carriers) semiconductors are illustrated in Figure 2.1. A thermoelectric device is

built up of these couples (Figure 2.1b). When the temperature gradient form across

the device, which induces the differences of carrier concentration, a voltage is

generated and drives a current pass through the device or an electric load. In the

case of refrigeration, a power supply drives the current and heat flow to cool the

heat-absorb part of the device.

The efficiency of power generation is9, 13, 14

=

√

√

⁄

(2)

- 9 -

where is the average temperature of the hot-end temperature (Th) and the cold-

end temperature (Tc).

The coefficient of performance (COP) for a refrigeration mode is9, 15

=

√

⁄

√ (3)

To judge the thermoelectric performance, figure-of-merit can be defined:

= (4)

Where Z is the figure-of-merit, the unit of Z is 1/K. T is the absolute temperature, S is

the Seebeck coefficient, σ is the electrical conductivity and κ is the thermal

conductivity. S2σ is the electrical power factor.

From these equations above, it is clear that higher thermoelectric performance relate

on higher ZT value of materials. Figure 2.2 demonstrate the relation between the

efficiency of thermoelectric devices and the temperature difference, which involved

the value of ZT. For a given operating temperature or a temperature difference, a

higher ZT leads a higher efficiency in power generation or refrigeration. As can be

seen from the Figure 2.2, when the ZT of thermoelectric devices close to or higher

than 2, the efficiency could achieve even higher than 15%, which could significantly

reduce the cost of power generation or refrigeration, therefore, make the application

of thermoelectric devices commercially available. For large scale and efficiency

commercial applications of thermoelectric materials, higher ZT value means more

competitive.9 To achieve high ZT value, high electrical power factor (S2σ) and low

thermal conductivity should be achieved.

- 10 -

Figure 2.2 Efficiency of thermoelectric devices as a function of ΔT.7

2.1.5 Effective Factors

The principles of the achievement high figure-of-merit or high thermoelectric

performance are complicated. By analysing the thermoelectric principles, the

effective factors of thermoelectric properties can be targeted to efficiently optimize

the thermoelectric performance of materials. Basically, a material which has a high

thermoelectric performance should have an appropriate carrier concentration and a

low thermal conductivity. Tuning the carrier concentration and thermal conductivity

through various methods is an efficient way to adjust the thermoelectric properties of

materials.

2.1.5.1 Carrier Concentration

To a great extent, electrical power factor (S2σ) determines the thermal function of

materials. So we should ensure that high Seebeck coefficient and high electrical

conductivity could be achieved synchronously. Both Seebeck coefficient (S) and

electrical conductivity (σ) relate to the carrier concentration of materials.16

=

⁄ (5)

- 11 -

= ⁄ = (6)

Where n is the carrier concentration of materials, is the effective mass of the

carrier, ρ is the electric resistivity, e is the electrical charge of an electron, and kB is

the Boltzmann constant and μ is the carrier mobility. Specifically, μ is the function of

E (the magnitude of the electric field applied to a material) and vd (the magnitude of

the electron drift velocity caused by the electric field). These equations show us that

lower carrier concentration leads lower electric conductivity but higher Seebeck

coefficient.

Figure 2.3 The relation between carrier concentration and the value of ZT.16

It is hard to achieve a high Seebeck coefficient with a high electrical conductivity due

to the complex connection of S, σ and n. Figure 2.3 shows that maximizing the ZT of

Bi2Te3 thermoelectric materials involves a compromise of the conflicting properties

(Seebeck coefficient, electronic conductivity and thermal conductivity). As the

thermal conductivity is also the function of carrier concentration, optimizing the

carrier concentration is crucial for maximizing the value of ZT. For an ideal

thermoelectric material which has been typically heavily doped, the carrier

concentration is between 1019 and 1021.

2.1.5.2 Thermal Conductivity

Thermal conductivity of materials comes from two sources16: electrons and holes

transporting heat (κe) and lattice thermal conductivity (κL):

http://en.wikipedia.org/wiki/Euclidean_vector

http://en.wikipedia.org/wiki/Electric_field

http://en.wikipedia.org/wiki/Euclidean_vector

- 12 -

= (7)

= = (8)

where L is Lorenz factor, 2.4×10-8 J2K-2C-2 for free electrons. As the electronic

thermal conductivity is directly related to the electric conductivity, κe should be

maintained at a high level while the thermal conductivity should be reduced for

achieving higher figure-of-merit.

2.1.5.3 Achieve High ZT

Theoretically, thermoelectrics require materials with high electrical power factor (S2σ)

and low thermal conductivity; furthermore, high electronic properties and low lattice

thermal conductivity17 should be achieved. These kind of special materials are called

a phonon-glass electron-crystal.6 Many crystalline semiconductors are identified as

good thermoelectric materials.18-23 In recent years, high ZT (>1) was extensively

achieved in heavily doped semiconductors,8, 16, 24 superlattice structures24-27 as well

as nanostructures.1, 5, 10, 24

State-of-the-art high ZT materials have been achieved by the using of nano-

technology. Both nano-miniaturization and nanocomposites are proved as efficient

approaches for the enhancement of thermoelectric properties, which are suggested

to increase the Seebeck coefficient and decrease the lattice thermal conductivity due

to the quantum confinement and the strong scattering to phonons. The theory and

principle will be discussed in the part of 2.3.

2.2 Development of Thermoelectric Materials

Since the thermoelectric effects were discovered, a large number of materials have

already been tested as thermoelectric candidates. Bi-Sb alloy system,28-32 Bi-Te

system,31, 33-39 Pb-Te system,40-48 Si-Ge system,49-52 some complex oxides2, 53 and

complex minerals with phonon-glass electron-crystal structure such as

Skutterudites8, 54-56 were considered as good thermoelectric materials, their ZT

values were experimentally measured as shown in Figure 2.4.

- 13 -

Figure 2.4 ZT as a function of temperature for several high-efficiency bulk

thermoelectric materials

(http://chemgroups.northwestern.edu/kanatzidis/greatthermo.html).

2.2.1 Thermoelectric Alloys

Figure 2.4 shows four kinds of high-performance thermoelectric materials with the

peaks of ZT located in different temperature extents. BiSb alloy has the highest ZT at

0-200 K. Bi2Te3 and its alloy have the highest ZT at room temperature, and have

maximum operating temperature11 at about 450 K. PbTe and SiGe have higher

operating temperature around 650 and 1200 K, respectively.

In 1962, Smith and Wolfe28, 32 reported that un-doped, Bi-rich, n-type Bi-Sb alloys

showed unique high ZT between 20-220 K while other materials’ ZT had significant

decrease in this temperature range, and their thermoelectric properties could be

improved with the presence of a magnetic field. Bi-Sb alloys are easy to prepare, a

high ZT of 0.88 could be reached at 80K in a 0.13 T magnetic field.29 The brittleness

is a drawback for the wide application of Bi-Sb alloys,30 which has been intensively

studied.

http://chemgroups.northwestern.edu/kanatzidis/greatthermo.html

- 14 -

Bi2Te3 is the best thermoelectric materials in the range of 300-500 K and had been

commercially applied in refrigeration devices for many years.57 Bi2Te3-based

thermoelectric materials can not only be used in room temperature but also can be

used for lower temperature applications by being doped. In 2000, Chung58 and co-

workers developed a new Bi-Te system material: CsBi4Te6 which achieved a

maximum ZT of ~0.8 at 225 K. This work pointed out a new approach to expand the

application temperature of thermoelectric materials. The new enhancement of bulk

materials’ thermoelectric properties could be expected.

The PbTe-based alloys have been extensively used for power generation, small-

scale cooling even power supplies for space exploration43, 59 due to their outstading

thermoelectric properties in middle temperature range. PbTe has a FCC crystal

structure with the band gap of ~0.32 eV at room temperature. PbTe could reach a ZT

of 0.8 near 550K, which could be significantly increased by doping or forming alloys

or soild solution. Poudeu60 and co-workers reported a ZT of 1.5 in Pb0.96Sb0.2Te10-

xSex system at 800K when x=7. In Poudeu’s another work,61 they also achieved very

high ZT in the system of Na1-xPbmSbyTem+2.

Si-Ge alloys are the best thermoelectric materials in high temperature arange (over

700 K), with a ZT of ~0.8 at around 1100 K (Figure 2.4). The thermoelectric

properties of Si-Ge alloys could be adjusted through changing the ratio of Si and Ge,

the conventional ratio is around Si80Ge20. Recntly, nanostructural SiGe62 has been

widely studied as well as bulk SiGe.

In additon, nowadays, the field of bulk thermoelectric materials are not limited on

these simple alloys or solid solution systems, more people have paid attention to

alloys with multiconponents, such as Ag-Sn-Sb-Te system63, (GeTe)(x)(AgSbTe2)(100-

x),64 Zn4Sb3/Bi0.5Sb1.5Te3

65 and complex oxides like La1−xSrxCoO366, NaCo2O4.

67

Hsu68 and co-workers achieved a ZT of 2.2 in AgPbmSbTe2+m , which showed us a

bright future of bulk thermoelectric materials. As high ZT of these materials have

been achieved in these materials, their extensive thermoelectric application in future

could be expected.

2.2.2 Materials of Complex Structures

Skutterudite (CoSb3) has a high Seebeck coefficient and high lattice thermal

conductivity. When the crystal ‘cage’ (Figure 2.5) was filled by guest atoms which

- 15 -

could bring the thermal conductivity down due to their fremitus in lattice, the

thermoelectric properties will be significantly improved. Figure 2.5d shows that the

highest ZT value of Co4Sb12 is 0.2 at 500 K. The thermoelectric property of Co4Sb12

was boosted to a very high level (>1) in a larger temperature extent after it was

doped by some other elements.

There are two types of Clathrates (I and II), both of them have low thermal

conductivities and open framework which could incorporate large electropositive

atoms. The principle of improving the ZT of Clathrates is similar as in the case of

Skutterudites.

Figure 2.5.(a-c) Crystal structure of CoSb3 revealing the large voids with rattlers, the

type-I Clathrate Na8Si46, and β-Zn4Sb3; (d) ZT as a function of temperature for

skutterudites as thermoelectric materials; (e) Variable temperature ZT of Clathrates,

and β-Zn4Sb3.69

2.2.3 Metal Chalcogenides Thermoelectric Materials

Metal chalcogenides materials are considered as ideal thermoelectric materials

which have shown promising thermoelectric properties in a wide range of applying

temperature. Among them, the best-performed thermoelectric material in a broad

- 16 -

temperature range, namely, copper selenide (800-1000 K), lead telluride (500-800 K)

and bismuth telluride (300-400 K) are extensively studied.

2.2.3.1 Cu2Se-based Thermoelectric Materials

Copper chalcogenides Cu2-xX (X= S, Se or Te) have been realized as a group of

promising thermoelectric materials due to their unique properties70-76 in last several

years, especially Cu2Se, which can be applied in photovoltaics,77 thermoelectrics,71,

78 photocatalyst,79 gas sensoring,80 electrode81 and superionic conductors.82

Generally, Cu2Se shows an α-phase when the temperature is lower than 400 K, it

has a monoclinic crystal structure with lattice parameters of a = 0.7138 nm, b =

1.238 nm, c = 2.739 nm, β = 94.308°, in which Cu ions are fixed in 12 positions

within the Se frame.83 When the temperature is increased and reaches to above 400

K, α-Cu2Se transfers to a β-phase with a lattice parameter a = 0.58 nm and the

space group ,71, 73, 84, 85 which can be demonstrated in Figure 2.6. Such a

phase transformation is reversible through cooling or heating processes. During the

phase transformation, Se ions formed a face-centre-cubic (FCC) frame and Cu+ ions

partially occupied86, 87 the 8(c) and 32(f) interstitial sites,86-90 these Cu+ ions are

highly mobile and exhibiting super-ionic liquid-like behaviour, which is crucial for its

intrinsically low κL because the phonons will be strongly scattered by such liquid-like

ions and finally lead to a high ZT of Cu2Se.71, 78

Interestingly, kinetically favoured β-Cu2Se was found as the preferred phase for

nanostructured Cu2Se product at room temperature rather than α-Cu2Se.78, 86 The

reason could be that the nanostructured Cu2Se has different surface energy state

compared with bulk materials.86 However, the phase transition from β- to α-phase

can be triggered by doping other elements such as Sb86 under certain condition. For

the Cu2Se-based materials, the stoichiometry is crucial for their structures and TE

performance.71, 91, 92 According to theoretical calculation,91 the stoichiometric Cu2Se

is a zero-gap material, significant Cu deficiency in β-Cu2-xSe can be allowed71, 90 to

form the non-stoichiometric Cu2-xSe (x=0-0.25)87 materials, which become intrinsic p-

type semiconductors with modified band structure. The existence of Cu deficiency

was found to affect the phase transition temperature88 and cause structural and

phase change87-90 of bulk Cu2-xSe compared to the stoichiometric Cu2Se. Low

temperature β-phase can be found in the composition range of 0.15≤x≤0.25 in Cu2-

3Fm m

- 17 -

xSe.88, 90 Theoretically, the existence of Cu deficiency in Cu2-xSe is expected to

change the electrical transport properties71, 92 by modifying the carrier concentration

but can also accelerate the cation exchange reaction93 and provide extra phonon

scattering by vacancies.94

Figure 2.6 Schematics shows the structures and phase transition of Cu2-xSe

between α- and β-phase.

Up-to-now, Cu2Se-based materials have been fabricated by various methods

including solid state reaction method,71, 73, 91, 95-97 ultrasonic chemical method,98

solvothermal method,78, 86 wet chemistry method,99, 100 Schlenk line techniques,70, 93

and self-propagating high-temperature synthesis.74 Among them, bulk Cu2Se (non-

stoichiometric) is an p-type semiconductor with a band gap of ~ 1.23 eV,71-74, 97, 101

and recently demonstrated a peak ZT of 1.5 with a high S2σ up to 12 µW cm-1 K-1 at

1000 K.71 Furthermore, the phase transition from α-Cu2Se to β-Cu2Se has resulted in

a high ZT (> 2) in I-doped Cu2Se.73 Zhong et al101 reported an peak ZT ~2.6 of bulk

Cu1.94Al0.02Se at around 1000 K due to the aligned lamellae structure.

2.2.3.2 Bi2Te3-based Thermoelectric Materials

As it has been discussed before, Bi2Te3 has the best thermoelectric performance at

room temperature,53, 102 which makes it the prior candidate for room temperature

- 18 -

thermoelectric power generation and Peltier cooling devices. Bi2Te3 is a narrow band

gap (~ 0.15 eV) semiconductor at room temperature,103 it belongs to a rhombohedral

crystal system but usually can be observed by an hexagonal primitive cell for

convenience (Figure 2.7).The hexagonal cells are stacked by the quintuple layers

(Figure 2.7a and c) collocated along c axis and bonded by van der Waals

interactions.104 Within the quintuple layer, there are ionic and covalent bands exist

between Bi and Te atoms, which are much stronger than the interlayer van der

Waals interactions and makes the Bi2Te3-based materials layered materials, which

have very anisotropic properties and can be easily cleaved along the planes

perpendicular to the c-axis.102 The band structure of Bi2Te3 has been extensively

studied.103, 105 For Bi2Te3, there are six valleys for the valence band maximums are

located in the mirror planes of the Brillouin Zone and the effective mass are highly

anisotropic.105 Due to the narrow band gap, high valley degeneracy and the

anisotropic effective mass, a high electrical conductivity and high Seebeck coefficient

of Bi2Te3 can be predicted. The carrier concentration, the conductivity type and the

electrical conductivity of Bi2Te3 should be easily adjusted by self-doping or doping

with other elements. By inducing mass contrast, defects and grain boundaries, the

thermal conductivity of Bi2Te3 can be efficiently reduced, thus, the overall

thermoelectric performance can be enhanced. Stoichiometrically crystalized Bi2Te3

ingots have intrinsic p-type conductivity102 due to the excess of Bi atoms and the

formation of antisite defects.102 The extra Bi atoms can occupy the Te site and

contribute as single electron acceptors. When an excess of Te was achieved, there

could be the Bi vacuums and antisite Te which replaced Bi and the conductivity type

of Bi2Te3 can be tuned to an n-type. Furthermore, the carrier concentration can also

be adjusted while the stoichiometric ratio of Bi and Te was modified. Although Bi2Te3

has the best thermoelectric performance (ZT~0.8) at room temperature, its

thermoelectric properties have to be further improved for achieving higher power

generation or cooling efficiency to make it more competitive in commercial

application.

- 19 -

Figure 2.7 Rhombohedral crystal structure of Bi2Te3.69

Different elements have been doped into Bi2Te3 crystal matrix to modify the

thermoelectric performance of Bi2Te3. In bulk materials, Sb and Se have been

successfully doped into Bi2Te3. A peak ZT of 1.4 has been achieved in p-type

(BiSb)2Te3 material at 373 K106 and a peak ZT of 1.56 has been obtained in

Bi0.52Sb1.48Te3 at 300 K.107 The highest ZT of (BiSb)2Te3 material had been reached

to as high as 1.8 at 316 K.108 In n-type bulk materials, a peak ZT of 1.04 was

observed in Bi2Te2.7Se0.3 at about 400 K and the ZT can be further increased to

about 1.1 by Cu doping.109, 110 The Seebeck Coefficients and the electrical

conductivities of those doped Bi2Te3-based materials have been obviously enhanced

while their thermal conductivities were significantly reduced. However, it is getting

harder to improve the thermoelectric performance of bulk Bi2Te3-based materials due

to the conflict relation between their thermoelectric properties. More and more

attention has been paid on nano-structured Bi2Te3 materials. Low-dimensional

Bi2Te3 materials have been intensively studied in recent years and many promising

- 20 -

results have been achieved.111-115 The main approach for improving the

thermoelectric performance of Bi2Te3 still lies on the increase of its electrical

conductivity and Seebeck coefficient while decrease the thermal conductivity. The

obtained ZT values of Bi2Te3 nanomaterials are still lower than that of doped bulk

materials, but the enhancement of ZT compared with pure Bi2Te3 are encouraging,

and the ZT can be expected to achieve a higher level in the future.

Table 1 Thermoelectric properties of advanced Bi2Te3-based materials.

Material Carrier

type

ZT κL/κ [W m-1 K-1] T

(K)

Synthetic

method*

Ref.

Bi2Te3-based bulk materials with nanocomposites

Bi0.4Sb1.6Te3 p 1.8 – 316 MS+HP 108

Bi0.4Sb1.6Te3 p 1.5 0.16 300 MS+HP 108

Bi0.52Sb1.48Te3 p 1.56 0.26 300 MS+SPS 116

Bi2Te2.7Se0.3 n 1.04 – 498 HEBM+HP 109

Bi2Te3-based quantum well or superlattice

Bi2Te3/Sb2Te3 p 2.4 0.22 300 – 117

Bi2Te3/Bi2Te2.83Se0.17 n 1.4 0.58 300 – 117

Bi2Te3-based nanomaterials

Ultrathin Bi2Te3

nanowire

n 0.96 0.92 380 SP+SPS 111

Te- Bi2Te3 nano-

barbell

p 0.24 0.309 400 SP+SPS 112

Bi2Te3-xSex n 0.54 0.6 300 SP+SPS 113

Ultrathin Bi2Te3

nanoplates

n 0.62 0.4 400 SP+SPS 114

Sulphur-doped

Bi2Te3 nanosheets

p 1.1 0.2-0.5 300 MW+CC 115

- 21 -

*The abbreviations used in the column of the synthetic method represent the

following meanings: MS= melt spinning; HP= hot pressing; HEBM= high energy ball

milling; SPS= spark plasma sintering; SP= solution phase method; MW= microwave

synthesis; CC= cold compaction.

2.2.3.3 PbTe-based Thermoelectric Materials

PbTe is a well-known IV-VI semiconductor with a narrow band gap (~0.3 eV) at room

temperature and large average excitonic Bohr radius of ~46 nm.118, 119 The crystal

structure of PbTe is face-centre-cubic. Because of the narrow band gap and the

large Bohr radius, a strong quantum confinement could exist in a large size range

(depend on the Bohr radius). In addition, PbTe has a lot of good physical and

chemical properties. It has high melting point, low vapour pressure, good chemical

stability and high ZT.119, 120 Furthermore, lead telluride can be easily doped to form

both n-type and p-type semiconductors. The self-doping of PbTe could be achieved

by varying the stoichiometry of lead and tellurium to modify the semi-conductive

property. The band gap of PbTe can also be adjusted by doping or alloying with

other elements. For instance, the band gap of PbTe will be reduced if it is alloyed

with Sn, or be increased by alloying it with Eu.118

The history that PbTe was used as practical thermoelectric materials can be tracked

back to 1930s.121, 122 In 1960s, both p- and n-type PbTe were used to assemble the

RTG for powering the spacecraft launched by The National Aeronautics and Space

Administration (NASA).121 Although the concentration had been shifted to Si-Ge

thermoelectric materials duo to their higher working temperature (above 1000 K) for

several years, the interest of lead telluride-based materials has recently been

reinvigorated because of their stability and promising thermoelectric efficiency. In

NASA’s latest Mars rover mission, PbTe-based materials were again chosen for

power supply.122

Lead telluride is a kind of covalent intermetallic compound with intrinsic low thermal

conductivity and good electrical properties.123 The FCC PbTe can be easily doped to

obtain n-type or p-type PbTe-based semiconductors with the implement of optimizing

their carry concentration, even the mechanical strength enhancement could be

achieved124 by doping or alloying. Both p- and n-type PbTe show the highest ZT in

the intermediate temperature range (500-900K) in all candidates of thermoelectric

- 22 -

materials,53 this operational temperature range is far below the melting point of PbTe

(1195 K).

The lightly doped p-type (by sodium) and n-type (by iodine) PbTe were

systematically studied and the peak ZT value of 0.7 for p-type and 0.8 for n-type

PbTe were found by Fritts123 in 1960. Limited by the approach of thermal conductivity

measurement, Fritts used the thermal conductivity measured at room temperature to

calculate the ZT values in the whole investigated temperature range. He realized that

the thermal conductivities were overestimated for high temperature and suggested

that the true ZT values should be higher.123 The reinvestigation of Na-doped PbTe

and I-doped PbTe had been done;125, 126 the laser flash method was introduced to

measure the thermal conductivity. The results showed the peak ZT value as 1.4

rather than 0.7 or 0.8. The doping level of PbTe was also proved could be further

optimized.

The outstanding thermoelectric performance of PbTe and PbTe-based materials

comes from the unique band structure of PbTe. With a high symmetry face-centre-

cubic crystal structure, convergence of many valleys can arise in PbTe.127 The high

valley degeneracy (Nv) lead a high Seebeck coefficient, thus, a high ZT.127

A two valence band model (Figure 2.8) of PbTe was suggested in 1960s.128, 129

When the band gap of PbTe was studied, it was found that ⁄ ( is the band

gap while T is the absolute temperature) was approximately zero at high temperature

range (above 450 K) while the ⁄ value was 4.1×10-4 eV/°K at low

temperature.128 This indicated that a principle valence band (L band, located at the L

point of the Brillouin zone) active at low temperature and a secondary valence band

(Σ band, located at the Σ point of the Brillouin zone) dominate at temperature higher

than 450 K.130 There is an energy separation of about 0.2 eV of these two valence

bands and the L band and Σ band were found to have valley degeneracies of 4 and

12, respectively.121, 127, 131 With the temperature increasing, the L band moves

converging with Σ band at about 450 K and providing an overall valley degeneracy of

16, which will greatly benefit to the thermoelectric performance of PbTe. When the L

band keeps moving below the Σ band, the dominating Σ band will still provide a

valley degeneracy of 12. The convergence temperature of these two valence band

- 23 -

was found could be adjusted by alloying or doping PbTe with another elements such

as Se,127 to broaden the applying temperature range of PbTe-based materials.

Figure 2.8 Valence band structure of PbTe1-xSex. (a) Brillouin zone showing the low

degeneracy hole pockets (orange) centred at the L point, and the high degeneracy

hole pockets (blue) along the Σ line. (b) Relative energy of the valence bands in

PbTe0.85Se0.15. At 500K the two valence bands converge, resulting in transport

contributions from both the L and Σ bands.47

The intrinsic high valley degeneracy and tuneable band structure makes PbTe as an

ideal candidate of high performance thermoelectric materials. Extensive efforts have

been made to obtain PbTe-based thermoelectric materials with high ZT value. Table

1 shows several advanced thermoelectric materials based on PbTe, high ZT values

were achieved in these materials by various methods in a large temperature range

(room temperature to 800 K), which makes the PbTe-based materials more

competitive compared with other potential thermoelectric materials used in the

intermediate temperature range (500-900 K).

So far, most studies of the thermoelectric performance of PbTe-based materials

have been focused on bulk materials because bulk materials are easy to be

prepared and their properties can be accurately measured compared with low-

dimensional materials. However, the peak ZT value of bulk materials is very hard to

achieve 2 or greater due to the conflict electrical and thermal properties which

determined the thermoelectric performance. With the development nano-science,

both theoretical25 and experimental26 studies suggested that it is expected to achieve

- 24 -

substantial ZT in low-dimensional materials. In last two decades, PbTe

nanomaterials have been paid much attention in synthesis and the improvement of

thermoelectric performance.

Up to now, various morphologies of PbTe such as nanospheres,132 nanowires,118, 133

nanosheets,134, 135 nanorods,136 nanoboxes,21 nanotubes,132, 137 nanoparticles,135, 138

hierarchical nanostructures139 and thicket-like119 nanostructures have been

fabricated by many methods. Molecular Beam Epitaxy (MBE),118 chemical vapour

deposition (CVD),140, 141 electrochemical deposition,119, 142 solvothermal or

hydrothermal43, 143, 144 methods are intensively used to prepare nanostructured PbTe

and its alloys or doped PbTe. Besides to form nanostructures such as quantum dots

to change the electronic density of states and reduce the lattice thermal conductivity

and increase the electrical power factor, one of other efficient approaches to

enhance the thermoelectric properties of lead telluride is to dope PbTe with some

other elements. The doping atoms can cause the distortion of not only the crystal

lattice but also the electronic density of states,46 which could also decrease the

thermal conductivity of the materials.

Solvothermal method has been identified as an efficient approach for synthesizing

PbTe with highly controllable processes. Well-crystallized lead telluride

nanostructures with variety of morphologies have been extensively fabricated and

reported and very high thermoelectric performances are expected to be achieved on

these PbTe nanostructures.

This project mainly concentrates on the understanding of experimental methodology,

the synthesis of nanostructured thermoelectric materials and to improve the

thermoelectric performance of PbTe-based nanomaterials. PbTe nanostructures

have been synthesized through a solvothermal route, and have been doped to

modify the carrier concentration and band structure. The samples have been

characterized by X-ray diffraction (XRD), scanning electron microscope (SEM) and

transmission electron microscope (TEM) to analyse their structures.

Table 2 Advanced thermoelectric materials in Pb-Te system

Materials Type of

carriers

ZT Synthesis

method

Temperature Ref.

- 25 -

Pb0.98Na0.02Te1-xSex - 1.8 Hot-

pressing

800K 47

PbSnSeTe superlattice n 2.0 MBE 300K 45

Tl-doped PbTe p 1.5 Mechanical

alloying

773K 46

PbTe–PbS

pseudo-binary

p 2.3 Solid

solution

923 K 145

AgPbm SbTe2+m n 2.2 Solid

solution

800K 68

PbTe-1% CdTe0.055% Pbl2 - 1.2 Melting 720K 146

PbTe/Pb1-xEuxTe - 2.0 MBE Room

temperature

26

Ag(Pb1-ySny)mSbTe2+m p 1.4 Melting 630K 147

(Pb0.95Sn0.05Te)0.92(PbS)0.08 - 1.5 - 642K 148

AgPbmSbTem+2 - 1.5 Mechanical

alloying

700K 20

La-doped PbTe-Ag2Te n >1.5 775K 42

Na-doped PbTe-Ag2Te >1.5 650K 149

Hierarchical architecture

SrTe-doped PbTe

p 2.2 SPS 915K 150

Na2Te-doped PbTe-SrTe p 1.7 SPS 815K 151

PbTe-CdTe - 1.7 Melting and

hot pressing

775K 152

phase-separated

PbTe0.7S0.3

p >2 Solid

solution

673 - 923 K 153

MgxPb1-xTe - 1.7 Melting and 750K 154

- 26 -

hot pressing

2.3 Principles and Methodologies to Achieve

High ZT

Based on the introduction of materials which intrinsically have high thermoelectric

performances, some necessary requirement for high performance thermoelectric

materials can be revealed: good σ, high S for electronic transport and low κ for

maintaining the temperature difference. Some other materials with such nature can

also be applied in thermoelectrics. However, pristine materials rarely achieve the ZT

above unity,7 which limits the application of thermoelectric materials and raises the

urgent demand to enhance their thermoelectric performances and implement the

commercialization.

2.3.1 Optimize the Carrier Concentration

The ideal thermoelectric materials could be described as ‘phonon-glass electron-

crystal’,155 which requires the κ as low as glass-like materials and the electrical

transport properties as good as crystalline materials.7 According to this principle,

several strategies can be applied to improve ZT of bulk materials, in which the most

effective approaches are to improve the electrical transport properties via optimizing

the carrier concentration and band engineering, which can decouple the σ and S,

therefore, achieve high power factor.

Among all the parameters which affect the thermoelectric performance of materials,

the S, σ and κe directly relate with n and are coupled by n as shown in equations (5),

(6) and (8), which makes controlling the carrier concentration a crucial strategy to

obtain high ZT. It is important to realize that the optimal carrier concentration (n*) is

strongly temperature dependent for most thermoelectric candidate as they are

semiconductors,156 so that the n* for different materials should greatly depend on

their application temperature.156 To provide a simplified prediction of the n*, the

relation n*~ (m*T)1.5 was suggested157 based on single band model and classic

statistics equations, where m* is the effective mass of carriers. According to this

- 27 -

relation, thermoelectric materials applied in lower temperature range should have

lower n* compared to materials which can be applied in higher temperature range.

For example, Bi2Te3158 has n ~1019 at room temperature while Cu2Se has the room

temperature n over 1020.71 Also, band engineering on the m* was applied to stabilize

the n* in PbTe,159 showed a good agreement with this relation. Based on the study of

La- and I- doped PbTe, Pei et. al156 obtained a simple estimation for n* as n*=3.25(T

/ 300)2.25×1018 cm-3. As it can be seen in Figure 2.9, such prediction provided fairly

accurate results compared to the experimental results and the n*~ (m*T)1.5 relation

(Figure 2.9a-d). However, the La-doped PbTe showed different experimental and

calculated power factor compared to the I-doped PbTe even when they had the

same n, which is because that I-doping can only tune the position of Fermi energy of

PbTe while La-doping can change the m*.156 They also introduced the reduced

Fermi level (ξ) as a guide to achieve the optimized ZT (Figure 2.9e and f) when the

ξ=0.3.156 In fact, the doping of La into PbTe not only tuned the n but also modified

the band structure of PbTe, which can also be realized as band engineering.

- 28 -

Figure 2.9 (a), (b) Experimental power factors as a function of Hall carrier

concentration; (c), (d) calculated power factors as a function of Hall carrier

concentration and (e), (f) calculated power factors as a function of reduced Fermi

level for Pb1-xLaxTe and PbIxTe1-x, respectively.156

2.3.2 Band Engineering

The main purposes of band engineering on thermoelectric materials are 1) tuning the

band gap160 to optimize the n, therefore, the σ in application temperature range and

2) tuning the m* (refer to equation 5) or achieve resonant state near the Fermi level

to obtain larger S46, 125, 161 via various doping.

- 29 -

2.3.2.1 Tuning the Band Gap

For the thermoelectric materials which can be applied at intermediate or high

temperature, bipolar conductivity is one of the main reasons to limit the peak ZT.162,

163 With the increased temperature, minority carriers164 will be significantly increased

due to the thermal excitation, increasing the electrical thermal conductivity, therefore,

the total thermal conductivity. This finding inspired the strategy which reduces the

thermal excitation of minority carriers at high temperature via increase the band gap

of materials. According to the theoretical165-166 and experimental160, 163 studies, the

band structures if semiconductors can be efficiently tuned by doping with other

elements. For example, by doping with Ag167 or Cd,152 the band gap of PbTe can be

enlarged, so the carrier concentration of the as-doped PbTe were stabilized at high

temperature range, which lead to slower degradation of power factor152 so that a

high ZT can be achieved.

2.3.2.2 Tuning the m* and Resonant States

According to equation 5, for a given carrier concentration, a high m* will lead to a

high S. The principle of increase the m* via doping is that it can change the flatness

of bands, result in an increased band effective mass (mb*). However, the high

effective mass will reduce the carrier mobility,164 therefore, reduce the power factor,

so that the high S value does not always lead to a high overall ZT.156 In the example

in part 2.3.1, the La-doped PbTe has higher S due to heavier m* compared to the I-

doped PbTe with similar doping concentration, but the I-doped PbTe shows slightly

higher net power factor (Figure 2.9), which will lead to a higher ZT.156 As a

consequence, practically, to enhance the electrical transport properties of

thermoelectric materials, it is more important to obtain an optimized power factor

rather than only focus on the S or σ individually.

Resonant doping46 can also effectively increase the m*. In some situation, the

impurity energy level lies in the conduction or valance bands (depend on the n- or p-

type conductivity), which creates a “resonant” density of state due to a distortion on

the host band (Figure 2.10), and when the Fermi level close to the resonant state,

the S will be significantly increased.46, 164, 168 Such effect has been found in Tl-doped

PbTe46 and Al-doped PbSe,169 which could be promising on developing high

performance thermoelectric materials.

- 30 -

Figure 2.10 Schematic illustrates the resonant level on the electronic density of

states (DOS).168

2.3.2.3 Tuning the Number of Valley Degeneracy

For some materials which have high symmetry crystal structures such as PbTe47

(Figure 2.8), achieving high valley degeneracy (NV) via band convergence is an

effective approach to enhance their thermoelectrical performance because it can

simultaneously achieve high electrical conductivity and high Seebeck coefficient.47

When the energy separation of bands is small compared with kBT (kB is the

Boltzmann constant), the bands could be converged, leading to increased NV, thus,

the increased m* without harming the carrier mobility (μ).47 Another important benefit

of achieving high NV is that the σtotal and Stotal will be increased because they contain

the contribution from the converged bands. As a consequence, the ZT will be

significantly improved by the increased NV.

2.3.3 Nanostructure Engineering

With the development of advanced technology, nanostructure engineering have

been realized as effective method to ontain high performance thermoelectric

materials.7, 168, 170-172 The nanostructure engineering for the thermoelectric materials

includes the development of low dimensional and nanocomposites, both of them

benefit from reducing the size and dimensions of materials, but with different

mechanisms. The enhancement of ZT on low dimensiomal thermoelectric materials

mainly due to the quantum confinement effect,25 while the nanocomposites involve

- 31 -

complex phonon scattering150 and low energy carrier filtering168 by the nano-scaled

substructures.

2.3.3.1 Low Dimensional Thermoelectric Materials

Nanostructures are defined as having one or more dimension between 1 and 100

nm.173 In 1993, Hicks and Dresselhaus25 suggested that low-dimensional or

nanocrystalline materials could achieve a high ZT due to the quantum size effect

which lead a higher density of states and a decreased lattice thermal conductivity

caused by the increase of phonon scattering.117, 174, 175 The dimensional decrease

causes the intensification of the quantum confinement as well as the decreasing

electron energy bands in nanostructure, which produce large Seebeck coefficient.

After that, a large number of theoretical calculations25, 17, 176, 177 and experimental

investigations26, 45, 178 have been done. High thermoelectric performance have been

extensively achieved in nanomaterials117 and a very high ZT are expected to be

obtained ultimately.

Remarkable enhancement of ZT of materials can be achieved through the

dimensionality decrease. In 2001, Venkatasubramanian117 reported the highest ZT

of ~2.4 in Bi2Te3/Sb2Te3 superlattice film, which boosted the thermoelectric property

compared to the bulk materials. Lin179 and Dresselhaus predicted that

heterostructure nanowires could have better thermoelectric performances than

superlattice films or conventional nanowires, which attract much attention to study

nanowires. However, the thermoelectric device based on this principle has not been

well-demonstrated yet.

2.3.3.2 Nanocomposites Thermoelectric Materials

Compared to nano-sized thermoelectric materials, nanocomposites174 materials are

easier to be fabricated and applied. Through appropriate procedures,150, 180-182 bulk

thermoelectric materials with nano-sized substructures can be synthesized. The

introduce of nanostructures into bulk materials will create high density of interfaces

(grain boundaries), lattice distortion and defects,172 which can significantly scatter

phonons which contribute to the lattice thermal conductivity, therefore, reduce the

thermal conductivity. Figure 2.11 shows the principle of how nanoparticles that

dispersed in bulk materials can scatter phonons with different wavelengths without

block the transport of electrons. The phonons with medium and long wavelength

- 32 -

could be strongly scattered by the nano-sized grains, grain boundaries, while short

wavelength phonons can be scattered by the point defects (such as atomic defects).

Such strong scattering to the phonons can reduce the lattice thermal conductivity at

most 50% and lead to a 2-fold increase in ZT.183 Meanwhile, electrons which have

much shorter wavelength than phonon would not be scattered so strongly, which

means that the electric conductivity will not be significantly influenced by the

nanocomposites.

Figure 2.11 Schematic diagram of phonon scattering mechaism and electronic

transport within a thermalelectric material.177

Another advantage of nanocomposites materials is that the boundary between the

nano-participates and the host materials can filter lower-energy charge carriers due

to the existing of potential barrier.168 As it can be seen form equation 5, the S value

strongly related with the carrier behaviour. From figure 2.12, low-energy carriers

negatively contribute to the Seebeck coefficient, while the carriers have energy

between 0.05 and 0.1 eV contribute the most. When the carriers transport through

the boundary potential barrier between grains, lower-energy charge carriers will be

- 33 -

filtered, therefore, the average carrier energy can be increased, and the |S| can be

increased as a consequence while the electrical conductivity will not be significantly

reduced.

Figure 2.12 Calculated results for n-type Si80Ge20 show how the carriers with

different energy contribute to the Seebeck coefficient.168

2. 4 Unsolved Issues and Opportunities

Due to the interesting physical and chemical properties, metal chalcogenides

materials are always considered as promising thermoelectric materials with huge

potential of wide application. With the increasing attention being paid, great efforts

have been made to improve the thermoelectric performance of metal chalcogenides-

based materials. The thermoelectric properties of Cu2Se, PbTe and Bi2Te3-based

materials have been significantly enhanced in last two decades, but there are some

issues have been left in which the opportunities can be addressed:

1. An efficient synthesis approach is required. For the requirement of the industrial

applications of thermoelectric materials, environmental friendly, high efficient and

low cost methods which can produce high-quality thermoelectric materials are

necessarily studied.

2. There is still a large gap of ZT value between the experimental results and

theoretical calculation, which means for all the current thermoelectric materials,

- 34 -

improvement of ZT is possible. For example, the calculated value of Seebeck

coefficient of p-type Bi2Te3 is 313 µVK-1 at 300 K ,105 which is superior than any

existing experimental value. If the high Seebeck coefficient can be achieve while

the thermal conductivity of Bi2Te3 approaching its amorphous limit, the ZT could

reach a very high value.

3. The mechanism of improving the ZT of Cu2Se, PbTe and Bi2Te3-based

nanomaterials by nano engineering and doping can be further studied and

explained systemically.

The goal of this project is to understand thermoelectric materials from a scientific

view, establish my own systemic research to understand master the synthesis

methodology to obtain Cu2Se, PbTe and Bi2Te3-based nanomaterials, and improve

their thermoelectric performance via nanostructure engineering and doping.

2.5 Summary

This Chapter has generally introduced the background for the project, provided a

literature review on the thermoelectric effects, the development of thermoelectric

materials, and the development of some major metal chalcogenides thermoelectric

materials, namely, Cu2Se, Bi2Te3 and PbTe. Based on the detailed review, although

extensive studies have been done on metal chalcogenides-based thermoelectric

materials, there are still urgent desires to further improve their thermoelectric

performance and understand the mechanisms for commercializing applications.

- 35 -

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170. Boukai, A. I.; Bunimovich, Y.; Tahir-Kheli, J.; Yu, J.-K.; Goddard, W. A., III;

Heath, J. R., Silicon nanowires as efficient thermoelectric materials. Nature 2008,

451, 168-171.

171. Heremans, J. P.; Dresselhaus, M. S.; Bell, L. E.; Morelli, D. T., When

thermoelectrics reached the nanoscale. Nat. Nano. 2013, 8, 471-473.

172. Kim, W.; Wang, R.; Majumdar, A., Nanostructuring expands thermal limits.

Nano Today 2007, 2, 40-47.

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173. Xia, Y. N.; Yang, P. D.; Sun, Y. G.; Wu, Y. Y.; Mayers, B.; Gates, B.; Yin, Y.

D.; Kim, F.; Yan, Y. Q., One-dimensional nanostructures: Synthesis, characterization,

and applications. Adv. Mater. 2003, 15, 353-389.

174. Fan, X. A.; Yang, J. Y.; Xie, Z.; Li, K.; Zhu, W.; Duan, X. K.; Xiao, C. J.; Zhang,

Q. Q., Bi2Te3 hexagonal nanoplates and thermoelectric properties of n-type Bi2Te3

nanocomposites. J. Phys. D Appl. Phys. 2007, 40, 5975-5979.

175. Lu, W. G.; Ding, Y.; Chen, Y. X.; Wang, Z. L.; Fang, J. Y., Bismuth telluride

hexagonal nanoplatelets and their two-step epitaxial growth. J. Am. Chem. Soc.

2005, 127, 10112-10116.

176. Pichanusakorn, P.; Bandaru, P., Nanostructured thermoelectrics. Mat. Sci.

Eng. R 2010, 67, 19-63.

177. Vineis, C. J.; Shakouri, A.; Majumdar, A.; Kanatzidis, M. G., Nanostructured

thermoelectrics: big efficiency gains from small features. Adv. Mater. 2010, 22, 3970-

3980.

178. Grabecki, G.; Wrobel, J.; Zagrajek, P.; Fronc, K.; Aleszkiewcz, M.; Dietl, T.;

Papis, E.; Kaminska, E.; Piotrowska, A.; Springholz, G.; Bauer, G., Quantum

nanostructures of paraelectric PbTe. Physica E 2006, 35, 332-337.

179. Lin, Y. M.; Dresselhaus, M. S., Thermoelectric properties of superlattice

nanowires. Phys. Rev. B 2003, 68, 075304.

180. Yang, L.; Chen, Z.-G.; Hong, M.; Han, G.; Zou, J., Enhanced thermoelectric

performance of nanostructured Bi2Te3 through significant phonon scattering. ACS

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Paramagnetic Cu-doped Bi2Te3 nanoplates. Appl. Phys. Lett. 2014, 104, 053105.

182. Cheng, L.; Chen, Z. G.; Yang, L.; Han, G.; Xu, H. Y.; Snyder, G. J.; Lu, G. Q.;

Zou, J., T-Shaped Bi2Te3–Te heteronanojunctions: epitaxial growth, structural

modeling, and thermoelectric properties. J. Phys. Chem. C 2013, 117, 12458-12464.

183. Kim, W.; Zide, J.; Gossard, A.; Klenov, D.; Stemmer, S.; Shakouri, A.;

Majumdar, A., Thermal conductivity reduction and thermoelectric figure of merit

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2006, 96. 045901.

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Chapter 3 Methodologies

In this project, solvothermal method is used to synthesize nanostructured tellurium-

based materials which are the characterized by using Scanning Electron Microscopy

(SEM), Transmission Electron Microscopy (TEM) and X-ray diffraction (XRD) to

investigate the morphologies, crystal structures, chemical composites of synthesized

materials. The thermoelectric properties were measured by using laser flash method,

chromel-niobium thermal couples and the Van der Pauw technique.

3.1 Synthesis Methods

In the extensively research of PbTe and Bi2Te3 nanostructures, various synthesis

methods have been used to prepare nanostructural products. Each method has its

advantages and disadvantages, which is showed in table 3. Compared with other

methods, solvothermal method showed its unique merits. Solvothermal method is

facile, commercially available and does not require any complicated technique and

exacting terms. Solvothermal method showed its controllability and stability in

previous works, in which nanostructures of materials were extensively synthesized

with various morphologies. In this project, solvothermal method will be focused.

- 52 -

Figure 3.1 An autoclave used in solvothermal synthesis.

Solvothermal synthesis is similar to the hydrothermal method, where the chemical

reaction takes place in solution rather than in aqueous and sealed in a stainless steel

autoclave (Figure 3.1). By heating the autoclave up to a certain temperature, huge

pressure can be generated to facilitate the reaction allows the growth of nano-sized

crystals. The shape, size, crystallinity of products can be adjust by varying the

experimental parameters including reaction temperature, reaction time, solvent

constituents, surfactants, and reactant types.

In a typical solvothermal route (Figure 3.2), reactants are weighed and dissolved into

specific solvent, which followed the designed stoichiometric and quality. The mixture

is stirred for several minutes or hours to form a well-disperse solution. The solution is

transferred to a Teflon-lined stainless steel autoclave and sealed. Then the

autoclave is put into an oven to be heated at certain temperature for a certain time.

After heating, the autoclave is taken out and cooled to room temperature naturally.

The product is collected and washed by centrifugation and dried for further use.

Table 3 Common synthesis methods of low dimensional materials.

Synthesis

methods

Advantages Disadvantages

Molecular Beam

Epitaxy1, 2

Stain-free

Defect-free

Highly controllable growth

Very expensive

Grow very slow

- 53 -

Form special structures

(superlattice)

Electrochemical

deposition3-5

3-D growth

Controllable process

Toxic electrolytes

High requirement for

substitutes

Sol-gel6 Cheap and low temperature

process

High purity of product

Controllable size and

chemical composition

Expensive

High requirement for

precursors

Time-consuming

Not stable during heating

and drying

Solvothermal7, 8/

Hydrathermal9-12

Cheap and low temperature

process

Controllable size,

crystallization and

morphology

Relatively efficient

Relatively environmental

friendly

Many experimental

variables

- 54 -

Figure 3.2 The process of a typical solvothermal route.

Usually, the temperature of the oven should be higher than the boiling point of the

solvent to generate an extremely high pressure in the autoclave, which drives the

reaction at a relatively low temperature. The advantages of the solvothermal route

are: (1) the sealed autoclave could efficiently avoid the oxidation of the product as

well as providing the high pressure to motivate the crystallization of nanostructures.

(2) The morphology could be adjusted by varying the reaction conditions. On the

other hand, details of the reaction process became out of control when we sealed

the autoclave, which impede us to control the reaction during the process. Because

of this, it is hard to assemble complex structure such as heterostructure through

solvothermal route. In this project, the reaction conditions will be adjusted to an

appropriate combination to achieve the best result.

3.2 Characterization Methods

3.2.1 X-Ray Diffraction (XRD)

X-ray diffraction is a popular method to analyse the composition, atomic or molecular

structure of materials. X-ray is a kind of electromagnetic radiation with wavelengths

between 0.006 and 2 nm generated by striking metal target with high energy electron

beam. X-ray can be diffracted when it penetrate through materials which have

periodic crystal structure. The incident angle (θ), the wavelength (λ) of X-ray and the

- 55 -

interplanar spacing (d) obey the equation of 2dsinθ= n λ (n is an arbitrary integer)

which is known as Bragg equation (Figure 3.3a).

Figure 3.3 (a) The principle of XRD (http://www.pic2fly.com); (b) A XRD

spectrophotometer (http://analyticalinstrumentengineer.com).

Figure 3.3b showed a XRD machine. The X-ray gun scans the sample through a

certain angle range (θ/2θ) while the X-ray detector collects the diffracted X-ray and

generates a diffraction spectrum. As the wavelength of X-ray is fixed for each

machine, therefore, the interplanar spacings can be calculated to identify the

composition of materials.

3.2.2 Scanning Electron Microscopy (SEM)

In SEM, the surface of the sample is scanned by a focused electron beam to

generate signals which contribute to the formation of the image (Figure 3.4a). The

electron beam is usually accelerated by a voltage around 1 to 50 kV, and with a

diameter of several nanometres. When this electron beam hits the surface of the

sample, it penetrates and generates an interaction volume.

http://www.pic2fly.com/

- 56 -

Figure 3.4 (a) Electron beam penetrates into the sample and generate different

signals with different information (http://mee-inc.com); (b) Changes in the interaction

volume with topography (http://www.ammrf.org.au).

The size of the interaction volume is directly proportional to the energy of the

electron beam (in other words, the accelerating voltage), and inversely proportional

to the atomic number of the sample.

Generally, secondary electrons and backscattered electrons are used as the signal

for imaging. Secondary electrons are electrons form sample atomic electron cloud

generated by inelastic scattering, which are with low energy (usually 0~50 kV) and

provide shape contrast of the surface of the sample (Figure 3.4b). Because of the

low energy of the secondary electrons, they are easy to be absorbed by other

sample atoms. The secondary electrons we detected are from a shallow region (5 to

50nm from the surface) so they can only produce the image near the surface of the

specimen. But the very detailed information of the surface and the topography can

be given by secondary electron image.

Basically, the electron beam generated from the electron gun and accelerated by an

anode in SEM. The electron beam is controlled by a condenser lens system to tune

the beam as desired size and direction. The spot size of the beam is adjusted as well

to obtain a suitable probe current. Then the electron beam is focused by objective

lens to scan on the surface of the sample. Detectors collect the secondary electrons

and backscattered electrons for the formation of images. To gain the higher

resolution of images, smaller spot size, small aperture, shorter working distance and

http://www.ammrf.org.au/

- 57 -

higher acceleration voltage should be used. But it should be concerned that short

working distance leads the decrease of depth of field, while higher acceleration

voltage reduce the surface information. For the different samples with different sizes,

surface features and conductivities, the operational condition of SEM should be

varied to obtain the best images with the desired information. By adding extra

detectors for other signals, a SEM (Figure 3.5b) can also be used to do analysis

such as energy dispersive X-ray spectroscopy (EDS).

3.2.3 Transmission Electron Microscopy (TEM)

When the thickness of sample is small (thinner than hundreds of nanometres) and

the acceleration voltage is high (100~1000 kV), the electron beam is able to

penetrate through and transmit the sample. The transmitted or forward-scattered

electrons can be used in image formation in a TEM (Figure 3.5a, b).

Figure 3.5 (a) The schematic outline of a TEM (http://www.hk-phy.org); (b) TECNAI

F20 TEM (http://www.bo.imm.cnr.it).

The electron beam is accelerated from a filament or a field-emission tip and then

condensed by lenses to control the spot size and the brightness of the beam. The

convergence angle and the intensity of the electron beam can be adjusted by the

apertures which fitted with the condenser lenses. After hitting the sample, the

http://www.hk-phy.org/

http://www.bo.imm.cnr.it/

- 58 -

electron beam carries the information of the sample will go through an imaging

system (Figure 3.5a). The signals are selected and magnified by several lenses and

finally form the image on the screen or camera. Because of the transmission of the

electron beam, the formed images are without the topographic information of the

surface of the sample. If the sample is crystalline, the incident electrons could obtain

the crystal information of the sample due to the scattering by the atomic planes or

diffracting by the crystal lattice. These electrons can form the diffraction pattern at

the back focal plane of the objective lens (Figure 3.6a and b). When the back focal

plane of the objective lens is used as the object plane of the intermediate lens, the

diffraction pattern can be obtained on the screen of TEM. When the image plane of

the objective lens is used as the objective plane of the intermediate lens, the image

will be shown instead of the diffraction pattern.

Figure 3.6 The principle of imaging and diffraction in TEM

(http://www.microscopy.ethz.ch).

In this study, the preparation of the sample is facile. The as-prepared sample is

nano-sized and do not require further treatment. The sample can be dispersed in

ethanol and dropped on the copper grill and then can be observed in the TEM.

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Other physical properties of prepared thermoelectric materials should be done for

understanding the relation between structure and thermoelectric performance of

materials.

3.3 Thermoelectric Measurements

Up to now, it is still very hard to get the accurate value of thermoelectric performance

for individual nano-scaled materials. To measure the thermoelectric properties of the

as-prepared nanomaterials, the sample powders were hot pressed to form the

pellets with certain densities. These pellets can be used to measure their thermal

conductivities, Seebeck coefficients and electrical conductivities. Then the overall

figure-of-merit can be calculated.

3.3.1 Thermal Properties

The thermal conductivity of measures samples can be calculated via κ=dCpD,13-15

where d is the density which can be measured using the mass and volume of the

sample pellet, Cp can be measured by the DSC 404 F3 (NETZSCH) and the thermal

diffusivity (D) can be measured by the laser flash method (Netzsch LFA 457).16 The

working principle of the NETZSCH DSC 404 F3 is shown in Figure 3.7: the signals of

the reference and the sample were obtained at a constant heating rate, which can be

used to calculate the Cp using a designed program. The principle of laser flash

method can be seen from Figure 3.8. In adiabatic condition, the sample is mounted

on a carrier system and is heated to reach a predetermined temperature by a pulsed

laser. The relative temperature is then measured by an IR detector as a function of

time. The thermal diffusivity is computed by the software using these time/relative

temperature increase data.

Figure 3.7 Schematic diagram shows the working principle of the Netzsch DSC 404

F3: (a) the furnace and (b) the calculation of Cp using obtained signals.

(https://www.netzsch-thermal-analysis.com)

https://www.netzsch-thermal-analysis.com/

- 60 -

Figure 3.8 Schematic diagram shows the working principle of the Netzsch LFA 457

system (from http://ap.netzschcdn.com).

3.3.2 Seebeck Coefficient

Seebeck coefficient of sample can be measured using chromel-niobium thermal

couples.17 As it is shown in Figure 3.9, the thermal voltage (thermal power) can be

measured by two thermal couples under a designed temperature gradient. Under a

given temperature gradient, several thermal voltages can be measured at different

temperature. According the definition of Seebeck coefficient, the sloped of the

http://ap.netzschcdn.com/

- 61 -

thermal voltage versus the temperature gradient will be the measured Seebeck

coefficient.

Figure 3.9 Schematic diagram shows the working principle of Seebeck coefficient

measurement.17

3.3.3 Electrical properties

The electrical resistivity and the Hall coefficient can be measured via the Van der

Pauw technique (Figure 3.10). Van der Pauw technique was developed to accurately

measure the properties of a sample of any arbitrary shape when the sample is

approximately two-dimensional and the electrodes are placed on its perimeter. From

the measurements, the resistivity of the material, the doping type, the carrier density

of the majority carrier and the mobility of the majority carrier can be calculated. When

a reversible magnetic field was applied, the Hall coefficient can be measured as well

(Figure 3. 11).

- 62 -

Figure 3.10 Schematics of Van der Pauw technique (http://www-

personal.umich.edu).

http://www.google.com.au/url?sa=i&rct=j&q=&esrc=s&frm=1&source=images&cd=&cad=rja&docid=HH6CPa4eNVkEBM&tbnid=cipWccLuR7egAM:&ved=0CAUQjRw&url=http://www-personal.umich.edu/~mandarg/projects.html&ei=hGFfUsKoEsa-kgXCzYGQBA&psig=AFQjCNFaBeUdAv7vNhPiS36Hm6OaoWW3Cg&ust=1382068980397622

- 63 -

Figure 3.11 The measurement of Hall coefficient using Van der Pauw technique

under a reversible magnetic field

(http://en.wikipedia.org/wiki/File:Van_der_Pauw_Method_-_Hall_Effect.png).

- 64 -

Reference

1. Moeck, P.; Kapilashrami, M.; Rao, A.; Aldushin, K.; Lee, J.; Morris, J. E.;

Browning, N. D.; McCann, P. J., Nominal PbSe nano-islands on PbTe: grown by

MBE, analyzed by AFM and TEM. Progress in Compound Semiconductor Materials

IV 2005, 829, 383-388.

2. Dziawa, P.; Sadowski, J.; Dluzewski, P.; Lusakowska, E.; Domukhovski, V.;

Taliashvili, B.; Wojciechowski, T.; Baczewski, L. T.; Bukala, M.; Galicka, M.; Buczko,

R.; Kacman, P.; Story, T., Defect Free PbTe Nanowires Grown by Molecular Beam

Epitaxy on GaAs(111)B Substrates. Crys. Growth Des. 2010, 10, 109-113.

3. Banga, D. O.; Vaidyanathan, R.; Liang, X. H.; Stickney, J. L.; Cox, S.;

Happeck, U., Formation of PbTe nanofilms by electrochemical atomic layer

deposition (ALD). Electrochim. Acta 2008, 53, 6988-6994.

4. Jung, H.; Park, D. Y.; Xiao, F.; Lee, K. H.; Choa, Y. H.; Yoo, B.; Myung, N. V.,

Electrodeposited Single Crystalline PbTe Nanowires and Their Transport Properties.

J. Phys. Chem. C 2011, 115, 2993-2998.

5. Liu, W. F.; Cai, W. L.; Yao, L. Z., Electrochemical deposition of well-ordered

single-crystal PbTe nanowire Arrays. Chem. Lett. 2007, 36, 1362-1363.

6. Bajaj, P.; Woodruff, E.; Moore, J. T., Synthesis of PbSe/SiO2 and PbTe/SiO2

nanocomposites using the sol-gel process. Mater. Chem. Phys. 2010, 123, 581-584.

7. Wang, W. Z.; Poudel, B.; Wang, D. Z.; Ren, Z. F., Synthesis of PbTe

nanoboxes using a solvothermal technique. Adv. Mater. 2005, 17, 2110-2114.

8. Zou, G. F.; Liu, Z. P.; Wang, D. B.; Jiang, C. L.; Qian, Y. T., Selected-control

solvothermal synthesis of nanoscale hollow spheres and single-crystal tubes of

PbTe. Eur. J. Inorg. Chem. 2004, 4521-4524.

9. Chen, X.; Zhu, T. J.; Zhao, X. B., Controllable Synthesis of PbTe Nanosheets

via an Alkaline Hydrothermal Method. Inec: 2010 3rd International Nanoelectronics

Conference, Vols 1 and 2 2010, 1179-1180.

10. Ni, Y. H.; Qiu, B.; Hong, J. M.; Zhang, L.; Wei, X. W., Hydrothermal synthesis,

characterization, and influence factors of PbTe nanocrystals. Mater. Res. Bull. 2008,

43, 2668-2676.

11. Sahoo, A. K.; Srivastava, S. K., Hydrothermal Synthesis of PbTe Nanorods

Using Different Templates. J. Nanoscience and Nanotechnology 2010, 10, 4921-

4928.

- 65 -

12. Tai, G. A.; Guo, W. L.; Zhang, Z. H., Hydrothermal synthesis and

thermoelectric transport properties of uniform single-crystalline pearl-necklace-

shaped PbTe nanowires. Crys. Growth Des. 2008, 8, 3878-3878.

13. Pei, Y.; Heinz, N. A.; LaLonde, A.; Snyder, G. J., Combination of large

nanostructures and complex band structure for high performance thermoelectric lead

telluride. Energy Environ.Sci. 2011, 4, 3640-3645.

14. Pei, Y.; LaLonde, A.; Iwanaga, S.; Snyder, G. J., High thermoelectric figure of

merit in heavy hole dominated PbTe. Energy Environ. Sci. 2011, 4, 2085-2089.

15. Pei, Y.; LaLonde, A. D.; Heinz, N. A.; Shi, X.; Iwanaga, S.; Wang, H.; Chen,

L.; Snyder, G. J., Stabilizing the optimal carrier concentration for high thermoelectric


16. Pei, Y.; LaLonde, A. D.; Heinz, N. A.; Snyder, G. J., High thermoelectric figure

of merit in PbTe alloys demonstrated in PbTe-CdTe. Adv. Energy Mater. 2012, 2,

670-675.

17. Snyder, G. J. Apparatus for measuring Seebeck coefficient of sample e.g. thin

film sample for thermoelectric device, has motorized stage that supports

thermocouple probe for positioning contact point of thermocouple probe at different

locations. US2013044788-A1.

- 66 -

Chapter 4 Controllable Synthesis

of Metal Chalcogenides

Nanostructures and Their

Thermoelectric Performances

4.1 Introduction

In this chapter, Cu2Se nanomaterials with controlled stoichiometry and Te doping

level were synthesized via a facile solvothermal method in order to tuning the

electrical transport properties, therefore, tuning the ZT. The impact of Cu deficiency

and Te doping to the structure and thermoelectric properties of Cu2Se have been

systemically studied. The Cu deficiency and Te dopant were found to trigger the

phase transition of β-Cu2Se to α-phase and significantly affected the thermoelectric

properties; their mechanisms have also been carefully studied.

4.2 Manuscripts for Publication

The results in Chapter 4 have been drafted as two manuscripts, which have been

submitted.

- 67 -

4.2.1 Effects of Cu Deficiency on the Thermoelectric Properties of Cu2-XSe

Nanostructures

Effects of Cu Deficiency on the Thermoelectric

Properties of Cu2-XSe Nanostructures

Lei Yang, Zhi-Gang Chen*, Guang Han, Min Hong, and Jin Zou*

L. Yang, Dr Z.-G. Chen, Dr G. Han, M. Hong, Prof. J. Zou.

Materials Engineering, The University of Queensland, Brisbane, QLD 4072, Australia

E-mail: [email protected], [email protected]

Prof. J Zou

Centre for Microscopy and Microanalysis, The University of Queensland, Brisbane,

QLD 4072, Australia

mailto:[email protected],%[email protected]

- 68 -

Abstract

Non-stoichiometric Cu2-xSe is one of important thermoelectric candidates for

intermediate temperature applications with intrinsically high performance at

800~1000 K. In this study, Cu-deficient Cu2-xSe nanoplates were synthesized by a

facile and controllable solvothermal method and the impact of Cu deficiency on their

corresponding thermoelectric performance was systematically investigated. It has

been found that α-phased Cu2-xSe can be induced by a relatively high level of Cu

deficiency (Cu1.95Se) in the as-synthesized Cu2-xSe nanoplates at room temperature.

The Cu deficiency was also found to reduce the thermoelectric performances, but

had no significant impact to the morphology of as-synthesized products. Overall, with

the existence of full-spectrum phonon scattering mechanism benefited from the

nanostructuring, the stoichiometric Cu2Se nanoplates showed an outstanding ZT of

1.82 at ~850 K due to its significantly reduced thermal conductivity. With increasing

the Cu deficiency, although the Cu2-xSe nanoplates showed a reduced ZT, such as

1.4 at 850 K for Cu1.98Se, it is still much higher than its bulk counterparts under the

same temperature.

Keywords: Copper selenide, Cu deficiency, nanostructures, thermoelectric

materials, induced phase transition.

- 69 -

Introduction

Due to the global energy shortage, the desire to optimize energy utilization rises.1-5

As one of the promising alternative energy sources, thermoelectric materials can

achieve solid-state power generation with applied temperature difference,1, 2, 5-12

which provide a new option to harvest electricity directly from waste or surplus heat.

The key issue for the commercialization of thermoelectric materials is to improve the

converting efficiency, governed by the figure-of-merit (ZT), which can be expressed

as ZT =S2σT/κ,1 where S is the Seebeck coefficient, σ is the electrical conductivity, T

is the absolute temperature, and κ is the thermal conductivity contributed by its

electron (κe) and lattice (κL) components. The major efforts have been made on

increasing σ and S, while reducing κ. However, these parameters are interrelated

and conflict with each other, so that it is a challenge to optimize them to obtain an

overall high ZT.

Recently, copper selenide (Cu2Se) has become one of the most popular

thermoelectric candidates due to its unique properties.3, 13-15 As demonstrated in

Figure 1, α-phased Cu2Se has a monoclinic crystal structure with lattice parameters

of a = 0.7138 nm, b = 1.238 nm, c = 2.739 nm, and β=94.308° at the room

temperature range, in which Cu ions are located in 12 positions.16 It transforms into

the face-center-cubic (FCC) structured β-phase with a lattice parameter of a = 0.58

nm and a space group of Fm ̅m16 when the temperature is higher than 400 K.3, 14, 17,

18 In β-Cu2Se, Cu+ ions partially occupied the 8(c) and 32(f) interstitial sites19-23

exhibiting super-ionic liquid-like behaviour with high mobility within the {111} planes

of the FCC frame formed by Se atoms,14, 15 which has been the key to lead to the

intrinsically low κL and in turn a high ZT.3, 15 So far, Cu2Se-based materials have

been fabricated by various methods, including solid state reaction method,13-15, 24-26

ultrasonic chemical method,27 solvothermal method,3, 23 wet chemistry method28, 29

Schlenk line techniques,30, 31 and self-propagating high-temperature synthesis.32

Among them, bulk Cu2Se demonstrated a ZT of 1.5 at 1000 K,15 while the phase

transition from α-Cu2Se to β-Cu2Se resulted in a very high ZT (>2 at the temperature

around 400 K).14 Interestingly, β-Cu2Se was found as the preferred phase for

nanostructured Cu2Se at room temperature rather than α-Cu2Se because it is

kinetically favoured.3, 23

- 70 -

Figure 1 Schematic diagrams show the structures and phase transition of between

α-Cu2-xSe and β- Cu2-xSe.

For the Cu2Se-based materials, the stoichiometry is crucial for their structures and

thermoelectric performance.15, 25, 33 According to the theoretical calculation,25 the

stoichiometric Cu2Se is a zero-gap material, but its Cu deficiency leads to the non-

stoichiometric and results in intrinsic p-type semiconductors with modified band

structures.15, 22 Practically, the existence of Cu deficiency in Cu2-xSe would not only

change the electrical properties,15, 33 but can also accelerate the cation exchange

reaction30 that provides extra phonon scattering by vacancies.34 Structurally,

increasing Cu deficiency can create extra Cu vacancies, leading to the distortion of

the lattice and instability of the entire structure. When the Cu deficiency reached to a

significant level, the Cu ions can be expected to re-arrange in the distorted Se frame

to form a more stable structure,23 so that it is highly desired to understand the

influence of the Cu deficiency on the thermoelectric performances of Cu2-xSe,

especially for nano-sized Cu2-xSe materials, which have not yet been systematically

investigated.

Our previous study3 revealed a facile solvothermal method to synthesize high-

performance nano-structured Cu2Se with an enhanced ZT of 1.82 at ~ 850 K. Such a

solvothermal method is highly controllable on the compositions of Cu2-xSe. In this

study, stoichiometric and non-stoichiometric Cu2-xSe plate-like nanomaterials were

- 71 -

fabricated with a controllable solvothermal approach.3 After sparking plasma

sintering (SPS) process, their thermoelectric properties and the related impact of the

Cu-deficiency are investigated in detail.

Experimental Section

Analytical pure copper (II) oxide (CuO, 99.995%), selenium dioxide (SeO2,

99.999%), sodium hydroxide (NaOH, 99.99%), ethylene glycol, polyvinylpyrrolidone

(PVP, average molecular weight: 40,000) from Sigma-Aldrich were used as

precursors without any further purification.

In a typical synthesis of Cu2-xSe nanostructures,3 0.4 g of PVP was dissolved in 36

mL of ethylene glycol, and then designed amount of CuO (1.5909 g for Cu2Se,

1.5750 g for Cu1.98Se, and 1.5511 g for Cu1.95Se), 1.1096 g of SeO2 and 4 mL of 5

mol/L NaOH solution were added in and stirred continuously. The solution was put

into a 125 mL Teflon-lined stainless steel autoclave and sealed and then heated at

230 °C for 24 hours. After that, the autoclave was cooled to room temperature. The

synthesized products were collected by centrifuging and washed by deionized water

and absolute ethanol for several times, and then dried at 60 °C for at least 12 hours.

The crystal structures of as-prepared products and corresponding sintered pellets

were characterized by X-ray diffraction (XRD), recorded on an X-ray diffractometer

equipped with graphite monochromatized, in which Cu Kα radiation (λ = 0.15418 nm)

was used. The morphological, structural, and chemical characteristics of as-

synthesized products and sintered pellets were investigated by scanning electron

microscopy (SEM, JEOL 7800, operated at 5 kV) and transmission electron

microscopy (TEM, Philips Tecnai F20, operated at 200 kV). A JEOL JXA-8200

EPMA (Probe) (operated at 20 kV) was used for the electron probe micro analysis

(EPMA). The TEM specimens for in-situ heating were prepared by focused ion beam

(FEI, SCIOS FIB-SEM).

The as-prepared Cu2-xSe powders were sintered into pellets by SPS under 50 MPa

and heated at 800 K for 5 min in vacuum. The Archimedes measurement was

performed to determine the densities (d) of the sintered pellets and their relative

densities (~ 95%).3

Thermal conductivity κ was calculated through κ = DCpd, where D and Cp are the

thermal diffusivity and specific heat capacity, respectively. D was measured by a

laser flash method with a LFA 457 (NETZSCH). A DSC 404 F3 (NETZSCH) was

- 72 -

used to measure Cp of Cu2-xSe pellets. σ and S were measured simultaneously on a

ZEM-3 (ULVAC). The carrier concentrations (n) of Cu-deficient samples were

estimated using a simple valence counting rule35-36 based on the obtained n for the

stoichiometric sample3 that every Cu vacancy contributes one carrier. The

uncertainty of the measurements of S, σ and D was ~ 5%, and the uncertainty for the

measured Cp was ~10%.

Results and Discussion

Figure 2a shows the XRD patterns taken from as-prepared powders with nominal

compositions of Cu2Se, Cu1.98Se, and Cu1.95Se, respectively. Also, diffraction peaks

from standard β-Cu2Se (Standard Identification Card, JCPDS: 06-0680) and a-

Cu2Se (Standard Identification Card, JCPDS: 47-1448) are marked in Figure 2a for

comparison.15, 23 As can be seen, the XRD patterns of Cu2Se and Cu1.98Se can be

exclusively indexed by the β-phase with a FCC structure (JCPDS 06-0680).3 These

results are consistent with our previous study,3 confirming that β-phase is the stable

phase for nanostructured Cu2-xSe. For Cu1.95Se, the major diffraction peaks can be

indexed as β-Cu2-xSe, but a few weak diffraction peaks can be indexed by the α-

phase (JCPDS 47-1448),3, 15 indicating that α-Cu2Se can be obtained when a certain

level of Cu-deficiency is reached, namely, Cu1.95Se in this study. Notably, the 030*

peak of α-phase in Cu1.95Se sample did not show obvious peak shift compared to the

standard value (Figure 2a inset, and the 200* peak is too weak to be observed),

indicating that the lattice of the as-prepared α-Cu2Se nanostructures has no

significant change compared to its equilibrium structure. It should be noted that the

relatively high non-stoichiometry in Cu2-xSe causes the higher Cu deficiency level in

β-Cu2-xSe,23 which promotes the significant lattice disorder and leads to the instability

of β-Cu2-xSe, which finally induces to form α-Cu2Se because the energy state could

be changed with increasing the Cu deficiency.23, 25 These facts could be the reason

why α-phase formed in Cu1.95Se. In β-Cu2-xSe, there are four partially occupied Cu

layers with high mobility superionic behaviour between neighbouring two Se layers

(refer to Figure 1). The Cu deficiency may cause the lattice distortion of β-Cu2-xSe.

When the distortion became significant (with sufficient Cu vacancies), the Cu atoms

will be forced to re-arrange in a different order to form a more stable structure.23

Instead of maintaining the highly distorted β-Cu2-xSe structure with a high Cu

deficiency level, the Cu ions rearranged in the distorted Se frames to form α-Cu2Se23

- 73 -

Therefore, there exists a threshold for the Cu deficiency level in β-Cu2-xSe, beyond

that thermodynamically stable α-phase can be formed. According to our results, the

threshold should be in the range of x = 0.02 ~ 0.05. Referring our XRD results, the

induced α-Cu2Se may prefer to have the Cu site fully occupied and the amount of the

α-Cu2Se should be small based on their weak diffraction peaks. This finding provides

a new phase-controlled synthesis approach of Cu2-xSe without introducing any

impurities. With increasing the Cu deficiency, lattice shrinkage of as-prepared β-

Cu1.98Se and β-Cu1.95Se can be expected. As shown in Figure 2b, clear right-shift

can be observed for the 111* and 400* diffraction peaks of β-Cu1.98Se and β-

Cu1.95Se when compared with β-Cu2Se. As the diffraction peaks of β-Cu1.95Se show

more shift than that of β-Cu1.98Se (Figure 2b), suggesting a larger lattice shrinkage in

β-Cu1.95Se when compared with β-Cu1.98Se due to the expected greater Cu

deficiency level of Cu1.95Se. Accordingly, the lattice shrinkage ratio can be calculated

for Cu1.98Se and Cu1.95Se as ~ 1% and ~ 1.5% along both [111] and [100] directions,

respectively. Moreover, the average grain size of Cu1.98Se and Cu1.95Se

nanostructures can be estimated as ~ 30nm using the Scherrer equation,3, 37 which

is similar as Cu2Se nanostructures.3

- 74 -

Figure 2 (a) XRD of as-prepared Cu-deficient Cu2-xSe nano powders compared with

stoichiometric Cu2Se and the standard cards for α- and β-Cu2Se, inset is the

enlarged 030* peak of α-Cu1.95Se; (b) XRD pattern for 111* and 400* peaks of β-Cu2-

xSe show the peak shift for different Cu deficiency.

To clarify the morphological characteristics of as-prepared products, SEM

investigation was employed. Figure 3 shows typical SEM images of as-prepared

Cu2Se, Cu1.98Se, and Cu1.95Se, respectively, in which stacked plate-like

nanostructures are seen with various lateral sizes (from several tens of nm up to 1

μm) and small thickness.

Figure 3 SEM images show the morphologies of Cu2Se (a), Cu1.98Se (b) and

Cu1.95Se (c) with the insets high magnification SEM images.

TEM was used to understand the structural characterisation of as-prepared

nanostructures. Figure 4a is a TEM image of a typical hexagonal-shaped β-Cu1.95Se

nanostructure, showing a lateral size of ~ 1 μm. To gain high-resolution TEM

(HRTEM) images, the nanoplate was titled to the [110] zone-axis. Figure 4b shows

such a HRTEM image, from which well-crystallized single crystal without observable

defects can be seen. Figure 4c is the corresponding selected area electron

diffraction (SAED) pattern taken from the [110] zone axis. Figure 4d is a TEM image

of another nanoplate found in the Cu1.95Se products, and Figures 4e and 4f are the

corresponding HRTEM image and SAED pattern. By analyzing these figures, the

nanoplate can be identified as α-Cu2Se based on the measured interplanar spacing

of 0.33 nm (corresponding to the {211} lattice spacing) and the index of the SAED

pattern (taken along the [ ̅11] zone axis). From the discussion above, our extensive

- 75 -

electron microscopy analysis are consist with the XRD results, and confirmed the

existence of α-Cu2Se in Cu1.95Se with the major β-Cu2-xSe.

Figure 4 (a) TEM image of a typical β-Cu1.95Se nanostructure; (b) HRTEM image

and (c) the corresponding SAED pattern from [110] zone axis; (d) TEM image of a

typical α-phase in as-synthesized Cu1.95Se product; (e) HRTEM image and (f) the

corresponding SAED pattern from [ ̅11] zone axis.

The thermoelectric properties of obtained Cu2-xSe nanostructures were investigated.

In doing so, SPS process was employed. Figure 5 shows the measured

thermoelectric properties of sintered Cu2-xSe nanostructures. Figure 5a shows the

measured σ versus T with the inset showing estimated n of different samples at

room temperature. As can be seen, the carrier concentrations and σ of Cu1.98Se and

Cu1.95Se obviously increase with increasing the Cu deficiency when compared with

the stoichiometric Cu2Se. This tendency is consistent with reported result.15 From the

nominal chemical ratio, Cu1.95Se sample has the highest Cu deficiency, thus the

highest σ = 1.2 × 105 S m-1 was obtained at room temperature with n = 1.5 × 1021 cm-

3, which decreases to σ = 3.4 × 104 S m-1 at ~ 850 K. In contrast, S reduces with

increasing the Cu deficiency, as shown in Figure 5b. Cu2Se has an outstanding S =

296 µV K-1 at ~ 850 K, while Cu1.98Se reached S = 212 µV K-1 and Cu1.95Se has the

- 76 -

lowest S = 170 µV K-1 at the same temperature.3 Figure 5c and inset plot the

measured D and Cp for determining κ, in which Cu1.98Se and Cu1.95Se have similar

Cp values, which are slightly higher than that of Cu2Se. Compared with bulk

samples,15 the measured Cp in our sintered Cu2-xSe nanostructures show relatively

low values, mainly attribute to the liquid-like super-ionic behaviour of Cu2Se and the

nanostructuring,3, 15 that result in very low κ. Figure 5d shows the determined κ, from

which Cu1.95Se shows the highest κ while Cu2Se has the lowest one, and κ for

Cu1.98Se is significantly lower than its bulk counterparts with similar composition.15

To further understand κ for our sintered pellets with different Cu deficiencies, κe for

Cu2Se, Cu1.98Se and Cu1.95Se were calculated3 using κe = LσT (where L is the

Lorenz number, in this study, L = 2.0 × 10-8 V2 K-2 is used3, 15). Also, based on κL = κ-

κe, Figure 5e plots the calculated κe and κL and shows that Cu1.98Se and Cu1.95Se

have the similar κL between 0.2 W m-1 K-1and 0.3 W m-1 K-1, which is very low and

comparable to the κL of Cu2Se sample.3 Such low κL values should be benefited from

the super-ionic nature of Cu2-xSe, and further reduced by the structural full-spectrum

phonon scattering3 (discuss later). The Cu deficiency did not make any significant

difference for κL but affect κe through modifying the electrical conductivity, as shown

in Figure 5e. Figure 5f showed the determined ZT as a function of temperature, in

which ZT decreases with increasing the Cu deficiency. Overall, for Cu2-xSe samples,

the increased Cu deficiency enhanced σ and κe but harmed S, leading to overall

decreased ZT. The ZT of Cu1.98Se reached the ZT of 1.4 at 850 K, although this

value is lower than that (1.82) of stoichiometric Cu2Se,3 it is still significantly higher

than its bulk15 and nanostructured33 Cu-deficient counterparts at the same

temperature. Cu1.95Se shows the lowest ZT of ~ 1 at 850 K, which is still comparable

to some popular thermoelectric candidates at the same temperature range.1

Additionally, for all the measured thermoelectric properties, fluctuations can be

observed at ~ 400 K, indicating there is a phase transition during the measurement

(will be discussed later). Since the influence of Cu deficiency to the thermoelectric

performance of Cu2-xSe is clear, its structural impact needs to be carefully

investigated.

To further investigate the impact of Cu deficiency to the sintered Cu2-xSe, the XRD

results of sintered samples were studied. Figure 6a shows XRD patterns of sintered

pellets, in which all diffraction peaks can be index as α-phase at the room

temperature, indicating the phase transition from β-phase to α-phase after the SPS

- 77 -

process. Such phase transition was observed and discussed in our previous work,3

the α-phase will transfer to β-phase when the measuring temperature is higher than

400 K (the phase transition is reversible), which is the reason why the fluctuations

can be observed in the measured properties. Although no impurity or secondary

phase were observed, clear diffraction peak shifts can be seen for the Cu1.98Se and

Cu1.95Se pellets (refer to the inset 030* peaks) when compared to the standard card

and α-Cu2Se, revealing that the shrunk lattice of Cu-deficient samples maintained

after the sintering process. Based on the diffraction peaks, the grain-size of the

nanostructures can be estimated using the Scherrer equation,3, 37 from which the

average grain size of all samples can be estimated as 30-40 nm. These XRD results

indicated that the SPS process did not significantly modified the crystal lattice and

grain size of Cu-deficient Cu2-xSe, but caused the phase transition from β-phase to

α-phase. To further identify the structural characteristics of the sintered Cu2Se,

Cu1.98Se and Cu1.95Se, detailed electron microscopy investigations were

comprehensively performed. Figure 6b shows the statistical EPMA results obtained

from multiple samples and confirms the compositions of our Cu-deficient samples as

Cu1.999Se, Cu1.98Se and Cu1.95Se, which are consistent with our nominal

compositions, indicating that the chemical composition of as-synthesized products

can be highly controllable by our facile synthesis method. Figures 6c-e are typical

SEM images of sintered Cu1.999Se, Cu1.98Se and Cu1.95Se pellets, respectively. As

can be seen, sintered Cu2Se, Cu1.98Se and Cu1.95Se samples have the similar grain

size distributions and small grain sizes, which are consistent with the XRD results.

Also, from these SEM images, we found that the grains of all sintered samples are

randomly orientated.

- 78 -

Figure 5 Measured temperature dependence of thermoelectric properties of

stoichiometric Cu2Se, Cu1.98Se and Cu1.95Se: (a) electrical conductivities (n is the

carrier concentration at room temperature); (b) Seebeck coefficient; (c) thermal

diffusivities with the inset specific heat values; (d) thermal conductivities; (e) electron

and lattice contribution to the thermal conductivity of Cu2Se, Cu1.98Se and Cu1.95Se

and (f) ZT values.

- 79 -

Figure 6 (a)XRD results of sintered Cu2Se, Cu1.98Se and Cu1.95Se (as marked) and

the inset shows broadened and shifted 030* peaks; (b) EPMA statistic results of

samples with the nominal compositions of Cu2Se, Cu1.98Se and Cu1.95Se, showing

the real average composition of Cu2Se, Cu1.98Se and Cu1.95Se, respectively; (c)-(e)

SEM images show the grains of sintered Cu2Se, Cu1.98Se and Cu1.95Se samples as

marked, showing the layer-by-layer stacking feature and randomly orientated grains.

- 80 -

TEM analysis was also performed to understand their structural characteristics.

Similar to SEM, their structural characteristics determined by TEM were very similar.

Therefore, results taken from only one pellet are presented. Figure 7a is a typical

TEM image of sintered Cu1.98Se pellet, and show the similar stacking of

nanostructures as shown in SEM investigations (refer to Figures 6c-e) with the

thickness less than 50 nm. Such a stacking manner of nanostructures results in very

high density of grain boundaries in the sintered pellet, which plays a critical role for

enhancing the thermoelectric performance.3 Figure 7b is a HRTEM image taken from

a typical grain boundary region (marked area in Figure 7a) and shows a high density

of lattice defects. The inset is the corresponding fast-Fourier transform (FFT) pattern,

which can be index as the [111] zone axis of α-Cu2Se. Figure 7c is the reversed FFT

image filtered by ±211* reflections, showing interplanar spacing of ~0.33 nm. Several

dislocation cores (marked as “T”) can be identified in this interface region, indicating

that the grain boundary is a small-angle grain boundary,3 which is crucial for

increasing the phonon scattering in order to enhance ZT.

Figure 7. (a) typical TEM image of sintered Cu1.98Se sample; (b) HRTEM image of

sintered sample taken from the marked area in (a) with the inset indexed FFT

pattern, showing a typical small angle grain boundary with can be clearly seen in the

reversed FFT image (c).

To determine the stability of the nano-grains, in-situ TEM heating experiment was

performed. Figure 8a is a typical TEM image of sintered Cu1.98Se sample, in which

multiple grains with clear grain boundaries can be observed. The TEM specimen

was then heated to ~800 K for 3 h and cooled down inside a TEM. Figure 8b is the

TEM image taken after 3 h heating, in which the nano-sized grains were remained,

indicating our sintered nanostructured Cu2-xSe is thermally stable. Furthermore, the

- 81 -

cycling test of thermoelectric performances of sintered Cu1.98Se and Cu1.95Se

samples also proved their stability. Figure 8c plots the cycling tests of Cu2-xSe

samples, showing similar ZT values in 5 cycles’ tests from room temperature to over

850 K for each pellet.

Figure 8. Typical TEM images of sintered Cu1.98Se sample before (a) and after

heating (b) up to 800 K with the inset SAED patterns from [111] zone axis. (c) ZT of

sintered Cu1.999Se, Cu1.98Se and Cu1.95Se samples measured for 5 cycles.

Based on above extensive structural and thermal property analysis, the impact of the

Cu deficiency on the thermoelectric performance of Cu2-xSe can be summarized as

follows. (1) In terms of electrical transport properties, the Cu deficiency in Cu1.98Se

and Cu1.95Se samples leads to increased carrier concentrations, thus increased σ

and decreased S, harming the overall ZT. Accordingly, the stoichiometric Cu2Se has

achieved the highest thermoelectric performance. (2) Morphologically, the nano-

structured Cu1.98Se and Cu1.95Se are similar as the stoichiometric Cu2Se3 in terms of

the grain sizes and their distribution, and all samples showed very fine average grain

sizes of 30-40 nm. Such fine grains were preserved after the SPS process, created a

high density of small-angle grain boundaries accommodated by a high density of

dislocations in the sintered samples, which can strongly scatter the phonons with

intermediate wavelength.3, 38 Combined with the phonon scattering by Cu

vacancies,34 liquid-like Cu ions,15 our Cu2-xSe nanostructures provide a full-spectrum

phonon scatterings,3 that significantly reduce κL, leading to superior thermoelectric

performances compared to their bulk counterparts.15, 33 (3) When the Cu deficiency

reached to a certain level (Cu1.95Se in this study), thermodynamically favoured α-

phase23 can be synthesized within kinetically stable β-phase, indicating that the

phase transition of Cu2-xSe can be triggered by the Cu deficiency without introducing

- 82 -

impurities, which may inspire the relevant studies of phase control on similar material

systems.

Conclusions

Cu-deficient Cu2-xSe (Cu1.999Se, Cu1.98Se and Cu1.95Se) plate-like nanostructures

have been controllably synthesized by a facile solvothermal method. The Cu

deficiency triggered the formation of α-Cu2Se in the major β-phase of Cu1.95Se

nanostructures, and caused lattice shrinkage in β-phase for both Cu1.98Se and

Cu1.95Se. After the SPS process, all pellets have α-phase and are maintained as

nano-sized grains with a high density of small-angle grain boundaries

accommodated by a high density of dislocations, which dramatically reduces κ. Our

Cu2Se nanostructures achieved an outstanding ZT of 1.82 at ~850 K. Although the

carrier concentration of as-prepared samples increases with increasing the Cu

deficiency, resulting in increased σ and reduced S, both Cu1.98Se and Cu1.95Se

samples still achieved relatively high ZT of 1.4 at ~ 850 K for Cu1.98Se, which is very

promising for Cu2-xSe-based thermoelectric materials, while the Cu1.95Se has the ZT

of ~ 1 at the same temperature.

Notes

The authors declare no competing financial interest.

Acknowledgements

This work was financially supported by the Australian Research Council. LY thanks

the China Scholarship Council for providing his PhD stipend. The Australian

Microscopy & Microanalysis Research Facility is acknowledged for providing

characterization facilities.

- 83 -

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- 87 -

4.2.2 Te-induced Phase Transition of Cu2SexTe1-x Nanomaterials and Their

Thermoelectric Properties

Te-induced Phase Transition of Cu2SexTe1-x

Nanomaterials and Their Thermoelectric

Properties

Lei Yang1, Eduardo Cauduro Manriquez1, Zhi-Gang Chen1*, Guang Han1, Min Hon1,

Liqing Haung1 and Jin Zou1,2*

1Materials Engineering, The University of Queensland, Brisbane, QLD 4072,

Australia

2Centre for Microscopy and Microanalysis, The University of Queensland, Brisbane,

QLD 4072, Australia

*E-mail: [email protected], [email protected]


- 88 -

Abstract

Understanding the impact of dopants will direct the design of high-performance

thermoelectric nanomaterials. In this study, we use Te as a dopant to modify the

crystal structure of Cu2Se. It has been found that Te has been uniformly distributed

in the synthesized products to form Cu2Se1-xTex nanoplates with controlled Te

content. Also, a phase transition from the original β-phase Cu2Se nanoplates to α-

phased Cu2Se1-xTex nanoplates was observed with the increasing Te doping level.

From the thermoelectric evaluation of the sintered pellets of Cu2Se1-xTex nanoplates,

Te can effectively modify their thermoelectric properties, especially their electrical

transport properties. Finally, a high ZT value of 1.76 at ~ 850 K has been achieved

for the Cu2Se0.98Te0.02 nanoplates, which were benefited from the good electrical

transport properties and ultra-low thermal conductivity. Especially, a special high

average ZT value ~1.2 with the range from 400 K to 850 K is observed in β-

Cu2Se0.98Te0.02 with an outstanding peak ZT of 1.76 at ~850 K.

Keywords: Tellurium doping; Copper selenide; Nanoplates; Induced phase

transition; Thermoelectric performance.

- 89 -

1. Introduction

As one of important p-type semiconductors, copper selenide (Cu2Se) has been

widely used in photovoltaics,1 thermoelectrics (TEs),2, 3 photocatalysts,4 gas

sensors,5 electrodes,6 and superionic conductors.7 In general, Cu2Se adopts

polymorphic α- and β-phases at different temperature ranges,2, 3 among which α-

Cu2Se is the thermodynamically favoured phase at room temperature while β-Cu2Se

is the kinetically stable phase for bulk Cu2Se.8 β-Cu2Se shows unique superionic

conductivity due to its crystal structure,3 in which Cu ions partially fill the interstitial

sites of the face-centre-cubic (FCC) frame constructed by Se ions, showing liquid-

like travelling behaviour.3 Such a Cu ionic fluidity9 provides strong phonon

scatterings, leading to an ultra-low thermal conductivity (κ) of β-Cu2Se.2, 3

With a band gap of ~1.23 eV3 and intrinsically high Seebeck coefficient (S),2, 3 β-

Cu2Se has drawn much attention as an intermediate temperature (500-900 K) TE

candidate in recent years2, 3, 9-13. The efficiency of a TE material is defined by its

figure-of-merit (ZT), determined as ZT =S2σT/κ,2, 14-16 where σ is the electrical

conductivity and T is the absolute temperature. High ZT values are required to

achieve high TE efficiency,14 therefore, S, σ, and κ need to be optimized to achieve a

high ZT for TE materials. So far, bulk Cu2Se has shown an intrinsic high ZT of 1.5 at

1000 K,3 and an outstanding ZT > 2 during the phase transition (at ~ 400 K).17 To

improve ZT of Cu2Se, several strategies have been developed, such as doping12, 13,

18 and nanostructuring.2, 19 Among which nanostructuring2 has been found as an

effective approach to further reduce the lattice thermal conductivity (κL) and in turn to

enhance ZT, while doping foreign elements into bulk Cu2Se can modify the electrical

transport properties.13 Theoretically, the doping elements can generate additional

point defects to increase the phonon scattering and consequently to reduce κ.20, 21

Therefore, the synergetic combination of nanostructuring and doping in developing

Cu2Se-based nanostructures may provide great opportunities to alter their electrical

transport properties. As a consequence, it is highly needed to fully understand the

impact of doping (dopant kinds and doping levels) on the TE performance of Cu2Se-

based nanomaterials.

In this study, tellurium (Te) was used as the dopant to synthesize Cu2Se1-xTex (x =

0.01, 0.02, 0.05, 0.10) nanoplates. Since, as an anionic dopant, Te2- has the same

valance as Se2-, but larger radii and heavier mass,13 Te doping into Cu2Se may

modify the Cu2Se lattice, enhancing the phonon scattering and also affect the

- 90 -

electrical transport properties via modifying the band structure.13 Thermoelectric

properties of obtained Cu2Se1-xTex and their underlying mechanisms are

investigated. The as-sintered samples show increased σ and reduced S and κL with

the increased Te doping level, and the nanostructured Cu2Se0.98Te0.02 sample shows

outstanding TE performance with ZT > 1.5 when the temperature is higher than 700

K, and reached a peak ZT of 1.76 at 850 K, which is promising compared to the bulk

un-doped Cu2Se3 and Te-doped Cu2Se.13

2. Experimental Section


99.999%), tellurium dioxide (TeO2, 99.999%), sodium hydroxide (NaOH, 99.99%),

ethylene glycol, polyvinylpyrrolidone (PVP, average molecular weight: 40,000) from

Sigma-Aldrich were used as precursors without any further purification.

In a typical synthesis of Cu2Se1-xTex nanostructures, 0.4 g of PVP was dissolved in

36mL of ethylene glycol. Under continuous stirring, 1.5909 g of CuO, varied amount

of SeO2 and TeO2 (for achieving x = 0.01, 0.02, 0.05, and 0.1), and 4 mL 5mol/L

NaOH solution were added in. The mixed solution was put into a 125 mL Teflon-lined

stainless steel autoclave and sealed, and then heated at 230 °C for 24 h in a CSK

thermal oven. After that, the autoclave was cooled to room temperature naturally.

The synthesized products were collected by centrifuging and washed by deionized

water and absolute ethanol for several times, and then dried at 60°C for at least 12h.

To evaluate the TE performance of synthesized products, the as-synthesized

Cu2Se1-xTex powders were sintered by spark plasma sintering (SPS) under 50MPa

and heated at 800K for 5min in vacuum. The Archimedes measurement2 was

performed to determine the density (d) and relative density (~ 95%). In this study, κ

was calculated using κ=DCpd,2 where D and Cp, are the thermal diffusivity and

specific heat capacity, respectively. D was measured by a laser flash method with a

LFA 457 (NETZSCH). A DSC 404 F3 (NETZSCH) was used to measure Cp. σ and S

were measured simultaneously in a ZEM-3 (ULVAC). The uncertainty of the

measurements of S, σ and D was ~ 5%, and the uncertainty for the measured Cp

was ~10%.

The crystal structures of as-synthesized products and sintered pellets were

characterized by x-ray diffraction (XRD), recorded on an X-ray diffractometer

equipped with graphite monochromatized, Cu Kα radiation (λ = 1.5418 ). The

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morphological, structural, and chemical characteristics of as-synthesized products

and sintered pellets were investigated by scanning electron microscopy (SEM, JEOL

7800 - operated at 5 kV) and transmission electron microscopy (TEM, Philips Tecnai

F20 - operated at 200 kV). The TEM specimens of sintered Cu2SexTe1-x were

prepared using focused-ion beam (FEI - SCIOS FIB). A JEOL JXA-8200 (operated at

20 kV) was used for the electron probe micro analysis (EPMA).

3. Results and discussion

Figure 1a is XRD patterns of as-synthesized Cu2Se1-xTex powders, compared with

an un-doped Cu2Se and its standard identification cards (JCPDS 47-1448 for the α-

phase and JCPDS 06-0680 for the β-phase).2, 3 As can be seen, the un-doped

sample (pure Cu2Se) can be indexed as pure β-phase.2, 8 With increasing the Te

doping level (x = 0.01), some diffraction peaks belonging to the α-Cu2Se are

appeared (for example, the 030* peak). Moreover, these diffraction peaks became

stronger with increasing the Te doping level, while the diffraction peaks belong to β-

Cu2Se (Standard Identification Card, JCPDS 06-0680)2 degenerated. When the Te

doping level reaches to 0.1, the XRD pattern can be only indexed as α-phase without

any obvious impurity or secondary phase. Figure 1b is enlarged 111* peaks (for β-

phase) taken from different Cu2Se1-xTex powders, which are clearly shown the phase

transition process with increasing the Te doping level. A slight left-shift of 111* peaks

of β-phase of x = 0.01, 0.02 and 0.05 samples and 410*, 211* peaks (α-phase) of x =

0.1 sample can be observed in Figure 1b for the Te-doped samples, indicating that

there exists a lattice expansion in Te-doped samples, which should be attributed to

the substitution of Se2- by Te2-. The broadened peaks reveal that the crystal grains of

the samples are relatively small, all the Te-doped samples have the similar grain size

of ~50 nm estimated using the Scherrer equation.2, 22

- 92 -

Figure 1 a) XRD patterns of Cu2Se1-xTex samples compared with the un-doped

sample and the standard cards. b) The enlarged peaks show the phase transition

with the increase Te doping level, the broadened peaks and the peak shift with Te

doping.

To further understand the room phase changes for nano-sized Cu2Se1-xTex powders,

the schematic atomic models of β-phase and α-phase are shown in Figure 2. In β-

Cu2Se, the Cu ions partially fill the 8 (c) and 32 (f) interstitial sites of the sub-lattice

formed by Se ions.3, 8, 23-25 For the FCC-structured Se frame (Figure 2), the

substitution of Se2- by Te2- can cause the lattice distortion as Te2- has larger radii

than Se2-. Such a lattice distortion induces the change of energy state of the

nanostructures,8 which in turn results in rearrangement of the lattice to form a more

stable structure (thermal dynamically favoured α-phase).8

- 93 -

Figure 2 Schematics of the phase transition of Cu2Se1-xTex nanostructures from β-

to α-phase, in which Se ions always have ordered structure while Cu is highly

disordered in β-phase but ordered in α-phase. Cu 32 (f) and Cu 8 (c) are the

possible interstitial sites for Cu ions.

To understand the morphologies of as-synthesized Cu2Se1-xTex powders, SEM

investigations were performed, as the results are shown in Figure 3. All the samples

with Te dopant have a wide range of size distribution from tens of nanometres up to

hundreds of nanometres. Although these samples have different Te doping levels,

they did not show significant morphological differences. Additionally, there is no

significant morphological difference between the mixture of β- and α-phase (Figures

3a-c) and the α-phase Cu2Se0.9Te0.1 (Figure 3d).

- 94 -

Figure 3 a)-d) SEM images of Cu2Se1-xTex samples with different doping level of Te

as marked.

Extensive TEM investigations were performed for all samples. As an example,

Figure 4 shows the typical TEM results of the Cu2Se0.99Te0.01 sample for both β- and

α-phase. Figure 4a is a typical TEM image of the β-phase Cu2Se0.99Te0.01 sample

which has hexagonal shape, showing a lateral size ~1.5 μm. From the high

resolution TEM (HRTEM) image (Figure 4b), the nanostructure is well-crystallized

without any obvious lattice defects and a lattice spacing ~0.34 nm can be seen. The

fast Fourier transform (FFT) pattern of the HRTEM image (Figure 4b inset) can be

indexed as a [110] zone axis of a FCC structure. The measured lattice spacing of

0.34 nm corresponds to the 111* lattice spacing of β-phase Cu2Se. In the same

sample, α-phase can also be found. Figure 4c is a typical TEM image taken from the

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α-phase nanostructures, which has a lateral size ~300 nm. A lattice spacing ~0.34

nm can be observed in the HRTEM (Figure 4d), but it corresponds to the 211* lattice

spacing of α-phase according to the FFT pattern (refer to Figure 4d inset). Therefore,

such nanostructures can be confirmed as α-phase. From the XRD and TEM results,

the doping of Te into Cu2Se did not change the crystal structure of β- and α-phase,

but lead to lattice expansions according to the left-shift of XRD peaks (refer to Figure

1b).

Figure 4 a) A typical TEM image of β-phase Cu2Se0.99Te0.01 sample shows

hexagonal plate-like shape. b) The HRTEM image taken long the [110] zone axis

and inset the corresponding FFT pattern confirmed the FCC crystal structure. c) A

typical TEM image of α-phase Cu2Se0.99Te0.01 sample shows hexagonal morphology.

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d) The HRTEM image taken long the [111] zone axis and inset the corresponding

FFT pattern from [111] zone axis reveals the monoclinic crystal structure.

To verify the compositions and how the Te is incorporated in Cu2Se, EDS and its

mapping were applied. Figure 5a is a TEM image of a typical Cu2Se0.98Te0.02

nanostructure and shows the hexagonal-shaped nanostructure with slightly irregular

stack. Figure 5b shows the corresponding EDS profile, in which Cu, Se and Te

peaks can be seen. Figures 5c-e are their EDS maps, respectively, which revealed

the uniform distribution of Te dopant. The brighter parts in those maps are due to the

thickness contrast, which contributed more signals than the thinner part.

Figure 5 a) The EDS patterns of Cu2Se0.98Te0.02 sample as an example. b) A typical

TEM image of Cu2Se0.98Te0.02 sample. c)-e) The EDS maps for (b) of different

elements as marked, showing the Te is uniformly distributed in Cu2Se. f) EPMA

results confirmed the compositions of Te-doped samples close to the nominal

compositions.

To measure their thermoelectric properties, these Cu2Se1-xTex nanoplates were

sintered using SPS. Their actual compositions of the sintered Te-doped Cu2Se

pellets were determined using electron probe micro analysis (EPMA), the results are

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shown in Figure 5f. The evaluated compositions of as-prepared Cu2Se1-xTex pellets

are very close to their nominal compositions as shown in Figure 5f, indicating that

our synthesis method is highly controllable to the Te doping levels. To further

understand the phase of sintered Te-doped Cu2Se pellets, their XRD results are

shown in Figure 6. As can be seen, all of the sintered Cu2Se1-xTex samples can be

indexed as α-phase without any impurities or secondary phases, indicating there

were phase transitions for Cu2Se1-xTex (x=0.01, 0.02 and 0.05) samples from β- to α-

phase after the sintering process. These phase transition after SPS are well matched

with the recently reported results.2 A broadened peak can be observed, which

reveals that Cu2Se1-xTex pellets have maintained the fine crystal size. A clear peak

left-shift can be seen in enlarged 030* peaks regions (the inset of Figure 6), which is

consistent with the XRD results of as-prepared powders before sintering.

Figure 6 The XRD patterns of sintered Cu2Se1-xTex samples with the inset enlarged

030* peaks show the broadened and shifted peaks.

The sintered samples have also been carefully investigated using electronic

microscopy. All of the sintered Te-doped samples have very small average grain

sizes, herein, only the SEM and TEM images of sintered Cu2Se0.98Te0.02 sample are

- 98 -

shown in Figure 7 as examples. Figure 7a is a typical SEM image of the sintered

Cu2Se0.98Te0.02 pellet, which shows a hierarchical sized distribution with very small

average grain size. The inset magnification SEM image of Figure 7a reveals the

layer-by-layer stacking feature of nano-sized grains. These stacking features also

confirmed by the subsequent TEM observations in Figure 7b. As can be seen, the

nano-sized crystals have an average thickness of ~ 50 nm, which aggregate to lead

to a high density of grain boundaries.2 The EDS map of Te element from the marked

area is shown in the inset of Figure 7b, in which Te is uniformly distributed. Figure 7c

is a typical TEM image of a grain boundary area of as-sintered Cu2Se0.98Te0.02, in

which the inset of FFT pattern can be indexed as [111] zone axis. The reversed FFT

image reveals the array of defect core, which indicates such a grain boundary is a

small angle boundary.2. Figure 7d shows another grain boundary region in the as-

sintered pellets, in which two adjacent grains forms a high angle boundary, as shown

in the inset of Figure 7d. According to the analysis results above, the as-sintered

pellets became α-phase (the Cu2Se0.9Te0.1 sample was α-phase before sintering)

without any precipitated phases, and the rapid SPS process preserved the nano-

sized grains of all the samples, creating high density of small angle grain boundaries

accommodated by defects as well as high angle grain boundaries.

- 99 -

Figure 7 a) SEM image of the sintered Cu2Se0.98Te0.02 sample as an example to

show small average grain sizes and the inset is the high magnification SEM image

shows the stacking feature. b) TEM image of sintered Cu2Se0.98Te0.02 sample as an

example to show the small grains with stacking feature, with the inset of Te map

shows the uniform Te distribution after sintering. c) HRTEM shows a typical small

angle grain boundary area, with (1) the inset FFT pattern and (2) reversed FFT

image shows the array of defects. d) TEM image of another grain boundary region,

the inset HRTEM reveals that it is a high angle grain boundary.

The thermoelectric properties of sintered Cu2Se1-xTex pellets are carefully analyzed

in Figure 8. Figure 8a shows the measured σ of sintered Cu2Se1-xTex pellets. The

Cu2Se0.99Te0.01 sample shows the lowest σ compared with other samples (Figure 8a),

revealing that the Cu2Se0.99Te0.01 sample has lower carrier concentration. With the

increase of Te dopant, the σ of Cu2Se0.98Te0.02 sample become higher, and the σ

- 100 -

further increases when the composition reached to Cu2Se0.95Te0.05. However, the σ

of Cu2Se0.9Te0.1 sample is lower than that of Cu2Se0.95Te0.05 when the temperature is

lower than 400 K, and higher than that of other samples in the temperature range of

400-550 K, but rapidly reduces over 475 K until the temperature reaches over 600 K,

then the reduce rate becomes lower (Figure 8a). Such unusual conductive behaviour

of Cu2Se0.9Te0.1 sample may due to the complex doping state of high concentration

of Te in the Cu2Se matrix, which makes the Te became saturated when the

temperature is higher than 475 K.26 The measured S values have been plotted in

Figure 8b. The Cu2Se0.99Te0.01 sample shows an ultrahigh peak S ~312 µVK-1 and

the S is higher than 300 μVK-1 when the temperature is higher than 720 K. With the

increase of Te dopant, the S values decrease as it can be seen in Figure 8b. The S

values of the Cu2Se0.9Te0.1 sample show the similar Te-saturated behaviour as the σ

when the temperature is higher than 475 K. The Cp and D values of Te-doped

samples have been measured and plotted in Figure 8c and 8d to determine the κ.

From Figure 8c, all the Te-doped Cu2Se samples have similar Cp values between

0.38-0.41 Jg-1K-1 for both α- and β-phase, which are higher than the un-doped

sample2 and comparable to the bulk samples.3 Additionally, all the samples have

very low D values which can be seen in Figure 8d. Combining the Cp values and D

values, the κ of Te-doped Cu2Se samples can be calculated as they are shown in

Figure 8e. All the samples have very low κ values between 0.3 Wm-1K-1 and 0.6 Wm-

1K-1, which are comparable to the un-doped sample.2 Such low κ is mainly attributed

by the low κL, which has been shown in the inset of Figure 8e. The principle of low κL

of Cu2Se-based nanostructured materials has been carefully discussed for the un-

doped sample.2 From Figure 8e, with the increase of Te content, the κL becomes

lower because the atomic contrast created by doping can increase the phonon

scattering, leading to a further reduced κL. Finally, the Cu2Se0.98Te0.02 sample shows

the best performance among the Te-doped Cu2Se samples, which has a ZT ~1.76 at

~850 K (Figure 8f). Such a high ZT value is superior in Cu2Se-based TE materials

(Figure 9 a) compared with the bulk Cu2Se3 and Te-doped Cu2-xSe,13 which confirms

the effectiveness of nanostructure engineering to enhance the TE performances of

materials. It should be highlighted that for our Cu2Se0.98Te0.02 sample, the ZT can

reach a high value of >1.5 at only 700 K, and it continually increases in the rest of

testing temperature range, leading to a high average ZT value ~1.2 for the β-phase

Cu2Se0.98Te0.02, which is superior compared to the reported bulk Cu2Se materials as

- 101 -

it can be seen in Figure 9b. With such high average ZT, the application temperature

range of Cu2Se-based thermoelectric materials has been significantly broadened,

making it even more promising as a thermoelectric candidate.

Figure 8 Measured temperature dependence of thermoelectric properties of Te-

doped Cu2Se with Te doping levels (as marked): a) electrical conductivities; b)

Seebeck coefficient; c) specific heat values; d) thermal diffusivities; e) thermal

conductivities with the inset of lattice contribution to the thermal conductivity of Te-

doped Cu2Se samples and f) ZT values.

- 102 -

Figure 9 The comparison of thermoelectric performances of as-sintered

Cu2Se0.98Te0.02 sample and bulk Cu2Se and Cu2-xTe0.08Se0.92 (adapted from Ref. 3

and 13, respectively): a) the ZT values in the whole temperature range and b) the

average ZT of β-phase.

On the basis of above analysis and discussion, Te was confirmed to be uniformly

doped into Cu2Se nanoplates with controlled doping levels. α-phase Cu2SexTe1-x was

found to be induced by Te doping, which gradually increases with the increase of Te

doping level til the product became pure α-phase when the composition reaches

Cu2Se0.9Te0.1. Moreover, Te doping significantly impacted the electrical transport

properties of Cu2Se nanomaterials. With the increase of Te content, the σ of the

products increases while the S decreases. Meanwhile, the Cu2SexTe1-x samples

achieved very low κ due to the intrinsic properties3 and the additional phonon

scattering via nanostructuring2 and Te dopant. Overall, the as-prepared

nanostructured Cu2SexTe1-x materials shows highly controllable TE properties by

tuning the Te doping level, and the Cu2Se0.98Te0.02 sample reached the ZT of 1.76 at

~ 850 K with a high average ZT ~1.2 for the β-phase.

Conclusions

In this work, Te has been successfully doped into Cu2Se nanostructures. The Te

doping did not significantly change the morphologies of the product, but triggered the

phase transition of Cu2Se1-xTex nanoplates gradually from β- to α-phase with the

increased Te dopant, which could be promising for studying the phase control of

Cu2Se-based nanomaterials. Benefited from the good electrical transport properties

and ultralow κ via nanostructure engineering and Te doping, the Cu2Se1-xTex

- 103 -

samples reveal good thermoelectric performances, especially, the β-phase

Cu2Se0.98Te0.02 sample shows an outstanding average ZT ~1.2 and has the high ZT

~1.76 at ~850 K.

Acknowledgements

This work was financially supported by the Australian Research Council. LY thanks


Microscopy & Microanalysis Research Facility is acknowledged for providing


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References:

1. Tyagi, K.; Gahtori, B.; Bathula, S.; Jayasimhadri, M.; Sharma, S.; Singh, N. K.;

Haranath, D.; Srivastava, A. K.; Dhar, A., Crystal structure and mechanical

properties of spark plasma sintered Cu2Se: An efficient photovoltaic and

thermoelectric material. Solid State Commun. 2015, 207, 21-25.



2015, 16, 367-374.



425.

4. Zhao, Y.; Zhu, L.; Jiang, Y.; Xie, H.; Zhang, G.; Ba, N., Microphone shaped

Cu2Se micro/nano architecture: Preparation, formation mechanism and optical

property. Mater. Lett. 2015, 147, 82-84.

5. Xu, J.; Zhang, W.; Yang, Z.; Ding, S.; Zeng, C.; Chen, L.; Wang, Q.; Yang, S.,

Large-Scale Synthesis of Long Crystalline Cu2-xSe Nanowire Bundles by Water-

Evaporation-Induced Self-Assembly and Their Application in Gas Sensing. Adv.

Funct. Mater. 2009, 19, 1759-1766.

6. Yue, J.-L.; Sun, Q.; Fu, Z.-W., Cu2Se with facile synthesis as a cathode

material for rechargeable sodium batteries. Chem. Commun. 2013, 49, 5868-5870.

7. Mahan, G. D., The Seebeck coefficient of superionic conductors. J. Appl.

Phys. 2015, 117. 045101.



20, 15941-15946.


Ultralow thermal conductivity of β-Cu2Se by atomic fluidity and structure distortion.

Acta Mater. 2015, 86, 247-253.

10. Yu, B.; Liu, W.; Chen, S.; Wang, H.; Wang, H.; Chen, G.; Ren, Z.,

Thermoelectric properties of copper selenide with ordered selenium layer and

disordered copper layer. Nano Energy 2012, 1, 472-478.




- 105 -



merit in bulk Cu1.94Al0.02Se. Appl. Phys. Lett. 2014, 105.






thermoelectric materials: current research and future challenge. Prog. Nat. Sci. 2012,

22, 535-549.


design of Bi2Te3 polycrystalline whiskers for thermoelectric applications. ACS Appl.

Mat. Interfaces 2014, 7, 989-995.







Adv. Mater. 2013, 25, 6607-6612.




19. Hu, H.; Ge, X.; Deng, C.; Sun, M.; Xuan, H.; Zhang, K., Copper selenide



20. Abeles, B., Lattice thermal conductivity of disordered semiconductor alloys at

high temperatures. Phys. Rev. 1963, 131, 1906-1911.





Cu2−δSe. J. Alloy Compd. 2003, 361, 57-61.

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23. Yamamoto, K.; Kashida, S., X-ray study of the average structures of Cu2Se

and Cu1.8S in the room temperature and the high temperature phases. J. Solid State

Chem. 1991, 93, 202-211.


Chem. Metals and Alloys 2011, 4, 200-205.

25. Wu, D.; Zhao, L.-D.; Tong, X.; Li, W.; Wu, L.; Tan, Q.; Pei, Y.; Huang, L.; Li,

J.-F.; Zhu, Y.; Kanatzidis, M. G.; He, J., Superior thermoelectric performance in

PbTe-PbS pseudo-binary: extremely low thermal conductivity and modulated carrier

concentration. Energy Environ. Sci. 2015, 8, 2056-2068.

- 107 -

Chapter 5 Enhanced

Thermoelectric Performances of

Metal Chalcogenides via

Nanostructure Engineering

5.1 Introduction

In this chapter, nanostructure engineering was approved as an effective strategy on

multiple materials system, including nanostructured Cu2Se, Bi2Te3 and PbTe. All

these materials were fabricated via highly controllable facile solvothermal method.

The as-synthesized samples showed significantly reduced thermal conductivity with

good electrical transport properties, which mainly benefited from the high density of

grain boundaries and defects introduced by nanostructuring. The underlying

mechanisms were demonstrated in detail.

5.2 Journal Publications and Manuscript

The results in Chapter 5 are included as it appears in Nano Energy 2015, 16, 367-

374.

http://www.sciencedirect.com/science/article/pii/S2211285515003055.

and ACS Applied Materials & Interfaces 2015, 7, 23694-23699.

http://pubs.acs.org/doi/pdfplus/10.1021/acsami.5b07596.

http://pubs.acs.org/doi/pdfplus/10.1021/acsami.5b07596

- 108 -

5.2.1 High-Performance Thermoelectric Cu2Se Nanoplates through

Nanostructure Engineering

High-Performance Thermoelectric Cu2Se

Nanoplates through Nanostructure Engineering

Lei Yang, Zhi-Gang Chen*, Guang Han, Min Hong, Yichao Zou, and Jin Zou*

L. Yang, Dr Z.-G. Chen, Dr G. Han, M. Hong, Y. Zou, Prof. J. Zou.



Prof. J Zou


QLD 4072, Australia

Keywords: Thermoelectric, High ZT, Copper Selenide, Nanostructure engineering


- 109 -

Abstract

As one of promising thermoelectric materials with intrinsic high figure of merit (ZT),

Cu2Se provides opportunities to tackle the global energy crisis via converting waste

heat into electricity. Here, β-phase Cu2Se nanostructures were synthesized using a

facile and large-scale solvothermal method. After sparking plasma sintering, the

resultant Cu2Se pellets show outstanding thermoelectric properties with an ultra-low

lattice thermal conductivity (as low as ~ 0.2 Wm-1K-1) that resulted in a recorded high

ZT of 1.82 at 850K. Through detailed structural investigations, high-densities small-

angle grain boundaries with dislocations have been found in sintered Cu2Se pellets

through nanostructure engineering, which results in additional phonon scattering to

reduce the lattice thermal conductivity. This study provides an important approach to

enhance thermoelectric performance of potential thermoelectric materials.

- 110 -

1. Introduction

Thermoelectric materials, directly harvesting electricity from heat or achieving solid

state cooling without any emissions or vibrational parts,1 offer a promising solution

for tackling the energy crisis.2 So far, extensive investigations have been made to

improve the thermoelectric efficiency, evaluated by the dimensionless figure-of-merit

ZT ( ), where σ is the electrical conductivity, S is the Seebeck

coefficient, T is the absolute temperature, and κ is the total thermal conductivity that

includes the contributions from its electron (κe) and lattice (κL) components.3-6 For an

ideal thermoelectric material, a high power factor (S2σ) and a low κ are required to

obtain a high ZT, so that a high thermoelectric efficiency can be secured. However, it

is always a challenge to optimize the individual parameters of σ, S and κ for

thermoelectric materials due to their interdependent and conflict.7 Up to now,

besides using band engineering through tuning band convergence,4 quantum

confinement,8, 9 and effective mass10 to maximizing S2σ, most successful ZT

enhancement has been achieved via structural and nanostructural engineering11 or

hierarchical architecturing3 to reduce κ.

Copper chalcogenides Cu2-xX (X= S, Se or Te), especially Cu2Se, have drawn much

attention as a group of promising thermoelectric materials due to their unique

properties.12-18 As a low temperature phase, α-phase Cu2Se (represented as α-

Cu2Se thereafter) has a monoclinic crystal structure with relatively low symmetry.19

When the temperature is increased and reaches to ~ 400 K, α-Cu2Se transfers to a

high temperature β-phase with the space group 13, 15, 20, 21. During the phase

transformation, Cu+ ions orderly stack along the <111> directions to form a simple

anti-fluorite structure from a very complex monoclinic structure with 144 atoms per

unit cell.19 Such a phase transformation is reversible through cooling or heating

processes. For the β-Cu2Se structure, Se atoms form a face-center-cubic (FCC)

frame and Cu+ ions are highly mobile and behave as liquid-like with reduced mean

free path for phonons,13 which results in a low κL with a value between 0.4 and 0.6

Wm-1K-1.13, 15 Bulk Cu2Se is an intrinsically p-type semiconductor,13-16, 22, 23 and

recently demonstrated a peak ZT of 1.5 with a S2σ up to 12 µW cm-1 K-1 at 1000 K.13

Furthermore, the phase transition from α-Cu2Se to β-Cu2Se has resulted in a high ZT

(> 2) in I-doped Cu2Se.15 With these potentials, it is necessary to further improve its

thermoelectric performance through novel strategies, such as nanostructuring or

2 /ZT S T

3Fm m

- 111 -

band engineering. Especially, nanostructuring has been theoretically predicted8 and

experimentally demonstrated5, 9 that can efficiently enhance ZT of thermoelectric

materials. Although there have been very success on developing bulk Cu2Se,

significant improvement of thermoelectric performance has not been achieved in

nanostructured Cu2Se.

In this study, we demonstrate a facile solvothermal method to synthesize high-quality

β-Cu2Se nanoplates and plate-like nanostructures. An enhanced ZT up to 1.82 at ~

850 K is observed in Cu2Se pellets after sparking plasma sintering (SPS)

processing. Such an enhanced ZT could be attributed by its very low κ, which

benefits from the strong phonon scattering by high-densities of small-angle grain

boundaries and dislocations within the boundaries via nanostructure engineering.

2. Material and methods

2.1 Materials synthesis


99.999%), sodium hydroxide (NaOH, 99.99%), ethylene glycol, polyvinylpyrrolidone

(PVP, average molecular weight: 40,000) were purchased from Sigma-Aldrich and

used as precursors without any further purification.

In a typical synthesis of Cu2Se nanostructures, 0.4g of PVP was dissolved in 36mL

of ethylene glycol, and then 1.5909 g of CuO, 1.1096 g of SeO2 and 4 mL 5mol/L

NaOH solution were added in with continuous stirring. The solution was put into a

125mL Teflon-lined stainless steel autoclave and sealed and then heated at 230 °C

for 24 hours. After that, the autoclave was cooled to room temperature naturally. The

products were collected by centrifuging and washed by deionized water and absolute

ethanol for several times, and then dried at 60 °C for at least 12 hours.

2.2 Materials Characterizations


characterized by XRD, recorded on an X-ray diffractometer equipped with graphite

monochromatized, Cu Kα radiation (λ = 1.5418 ). The morphological, structural,

and chemical characteristics of as-synthesized products and sintered pellets were

investigated by SEM (JEOL 7800, operated at 5 kV for normal SEM and 15 kV for

back-scattered SEM) and TEM (Philips Tecnai F20, operated at 200 kV). A JEOL

JXA-8200 (operated at 20 kV) was used for the electron probe micro analysis.

- 112 -

2.3 Sintering process

The as-synthesized Cu2Se powders were compressed by SPS under 50 MPa and

heated at 800 K for 5 min in vacuum. The Archimedes measurement24 was

performed to determine the density (d) and relative density (95%).

2.4 Thermoelectric Performance Measurement

Thermal conductivity κ was calculated through ,13 where D and Cp, are the


laser flash method with a LFA 457 (NETZSCH) and obtained D was plotted in Figure

S4a. A DSC 404 F3 (NETZSCH) was used to measure Cp of Cu2Se, and Figure S4b

plots the measured Cp. σ and S were measured simultaneously on a ZEM-3

(ULVAC). The uncertainty of the measurements of S, σ and D was ~ 5%, and the

uncertainty for the measured Cp was ~10%. The standard deviation of the measured

ZT from 5 different samples (Figure S1) was ~3%.

3. Results and Discussion

Nanostructured Cu2Se was synthesized via a facile solvothermal method. Figure 1a

shows the X-ray diffraction (XRD) pattern of as-synthesized products, in which all

diffraction peaks can be exclusively indexed as the FCC structured β-Cu2Se with the

lattice parameter of a = 0.5739 nm (Standard Identification Card, JCPDS 06-0680).13,

21 Detailed analysis of the XRD pattern shows the peak broaden, as shown in Figure

1a inset. Using measured broadened 111* diffraction peak, we can estimate the

average grain size of as-synthesized products by the Scherrer equation25 which

gives the average size of ~30 nm. It is of interest to note that the obtained β-Cu2Se is

stable at room temperature, although it is a high-temperature phase (> 400 K)

according to the binary Cu-Se phase diagram.15, 26 This may be due to the fact that

such a kinetically favored β-phase26, 27 was fabricated by the solution method with

high temperature (230 °C, higher than the phase transformation temperature), high

pressure synthesis conditions (approximately 300 kPa, generated in the sealed

autoclaves).26 According to Ref.,26 low-dimensional β-Cu2Se will not transfer to α-

Cu2Se after cooling process, which is different with bulk Cu2Se. The reason could be

that kinetically favored β-phase is more stable for nano-sized Cu2Se which has very

high surface energy. Figure 1b is a scanning electron microscopy (SEM) image and

shows the typical morphologies of synthesized products, in which hexagonal

pDC d

- 113 -

nanoplates and self-assembled plate-like components can be observed and their

lateral size contribution can be estimated from several hundred nm to 1 µm.

Transmission electron microscopy (TEM) investigations were further employed to

determine their structural characteristics. Figure 1c shows a TEM image of a typical

hexagonal Cu2Se nanoplate with a lateral size of 200 nm. Figure 1d and its inset are

the high-resolution TEM (HRTEM) image and the corresponding selected area

electron diffraction (SAED) pattern of the nanoplate, from which the nature of highly

crystallized signal crystal nanoplate can be seen. Our extensive TEM investigations

on individual nanoplates suggest that they are single crystals with no observable

lattice defects.

- 114 -

Figure 1 (a) XRD patterns of as-prepared Cu2Se nanostructures, compared with the

standard card, inset being the enlarged 111* peak; (b with inset) SEM images for as-

prepared Cu2Se nanostructures; (c) TEM images of an as-prepared Cu2Se

nanoplate; (d) corresponding HRTEM image and SAED pattern (inset) of the Cu2Se

nanoplate.

To measure the thermoelectric properties of our synthesized products, the as-

prepared Cu2Se powders were sintered using the SPS approach to obtain disc-

shaped pellets. Their determined thermoelectric properties are showed in Figure 2.

Figure 2a shows the measured σ values as a function of T. An σ of 6.68×104 S m-1 is

observed at room temperature with a carrier concentration ~ 3×1020, and decrease to

0.95×104 S m-1 with increasing T up to 850 K. This trend is comparable to reported

bulk Cu2Se.13 It should be noted that, a clear fluctuation, around the phase

transformation temperature at ~ 400 K, can be observed here and in the other

measurements, indicating the phase transformation took place during the transport

property measurements in which T is increased. However, understanding the impact

of phase transformation on the thermoelectric properties is not the scope of this

study, so that no further concerns on this observation will be discussed. Figure 2b

shows the relationship between measured S and T, in which the positive S indicates

its p-type nature with the majority of carriers being holes.13 With increasing T, S

increases and reaches to a peak value of 296 µVK-1 at 850 K, which is greater than

the reported S value for bulk Cu2Se.13 Since thermal conductivity κ is determined by

κ = DCpd (where D is the thermal diffusivity and Cp is the specific heat capacity, and

d is the materials density),13 we measured D and Cp, and their relationships with

temperature are shown in Figure S4. The obtained Cp values are between 0.31 and

0.32 at high temperature (>500 K), which is lower than the reported values around

0.36~0.40. Our low measured Cp value could be attributes to the decreased

constant-volume heat capacity (Cv) from 3NkB (where N is the number of particles

and kB is the Boltzmann constant) of a typical solid to (2-2.5)NkB13, 28 of a liquid

behavior β-Cu2Se because the propagation of most transverse vibrational waves can

be disrupted by local atomic jumps and rearrangement in this superionic liquid-like

crystal,13 Furthermore, the existence of high-density small-angle grain boundaries

with high-density dislocations in nano-sized grains may further enhance the

disruption, which finally contributes the low Cp. Figure 2c presents the measured κ

- 115 -

as a function of T, in which a very low κ between 0.4 Wm-1K-1 and 0.6 Wm-1K-1

(except the values measured during the phase transformation) are observed in the

entire temperature range, which is significantly lower than that of bulk Cu2Se (slightly

less than 1 Wm-1K-1).13 To understand the individual contributions of κe and κL, we

calculate κe using κe = LσT (L is the Lorenz number).13 Here, L = 2.0×10-8 V2K-2 is

used to estimate κe,13 and the correspondingly obtained κe is plotted in Figure 2c.

Using κL = κ-κe, the correspondingly κL can also be obtained and plotted in Figure 2c.

The comparison of plots of κe and κL clearly indicates that κe contributes the major

thermal conductivity up to values between 0.17 Wm-1K-1 and 0.43 Wm-1K-1 (T < 750

K) while the κL is only between 0.11 and 0.15 Wm-1K-1 for the α-Cu2Se and around

0.12~0.23 Wm-1K-1 for the β-Cu2Se. As can be seen, our κL of β-Cu2Se is relatively

independent on the temperature ranging from 400 K to 850 K, and is much lower

than that (0.4-0.6 Wm-1K-1) of bulk β-Cu2Se.13 Such a low κL is not only attributed by

the liquid-like behavior from the superionic Cu+ ions, which strongly scatter the

phonons.13, 15, 22-23 but also contributed by the nanostructure engineering (discuss

later). Such an improved S and the very low κ found in our sintered Cu2Se pellets

finally lead a high ZT of 1.82±0.05 at 850 K, as shown in Figure 2d and S1). This

value has achieved over 20% improvement when compared with reported bulk

Cu2Se.13 In fact, this value is comparable to those highest thermoelectric materials

reported so far.2, 4, 29-30

- 116 -

Figure 2 Measured temperature dependence thermoelectric properties of a typical

sintered Cu2Se pellet: (a) electrical conductivity (n is the room temperature carrier

concentration), (b) the Seebeck coefficient, (c) thermal conductivities of Cu2Se and

(d) calculated ZT values.

To fundamentally understand our observed exceptional thermoelectric performance,

detailed structural characterizations were performed on the sintered pellets. Figure

3a is the XRD pattern taken from a sintered pellet at room temperature, in which all

diffraction peaks can be indexed as the monoclinic structured α-Cu2Se with the

lattice parameters of a = 0.7138 nm, b = 1.2382 nm and c = 2.739 nm (identification

card JCPDS 47-1448),19 indicating that the as-prepared β-Cu2Se nanostructures has

transferred to pure α-Cu2Se after the SPS sintering. It should be noted that broaden

of diffraction peaks can also be observed (as shown in Figure 3a inset), indicating

the small grain sizes in sintered Cu2Se pellets. According to the Scherrer equation,25

the average grain size of our sintered Cu2Se pellets can be estimated as ~ 36 nm,

slightly larger than the average grain size (~ 30 nm) measured from as-synthesized

Cu2Se nanostructures. It is of interest to note that, although the as-synthesized

- 117 -

nanostructures are the β-phase while sintered Cu2Se pellets are the α-phase, both

measured at room-temperature; their average sizes do not change significantly. This

suggests that, during the sintering process, although phase transformation took

places that can lead to a significant structural change, the grain sizes do not change

significantly. This may be due to the fact that SPS requires lower sintering

temperature and shorter holding time at the high temperature,31, 32 so that the

nanostructures can be preserved due to the minimization of the Ostwald ripening.33

To clarify this, we investigate the morphology of fractural surfaces of sintered pellets.

Figure 3b shows an example, in which the SEM image shows the morphology of

plate-like nanoparticles stacked together with plate thicknesses < 50 nm (Figure 3b

(1)). Figure 3b inset (2) is the back-scattered SEM image taken from a polished

surface of a sintered Cu2Se pellet, and confirms its high density without obvious

pores and no observed secondary phases. Based on our estimation using the

Archimedes method,24 the relative density of sintered pellets is up to 95%. Electron

probe micro analysis was applied to determine the stoichiometry of sintered Cu2Se

pellets and results are shown in Figure 3c, indicating that the measured

stoichiometry of our sintered pellets is Cu1.999Se with an error bar of 0.1%, which is

almost identical to the nominal stoichiometry.

It should be noted that the above structural characterizations are performed on

sintered α-Cu2Se pellets, so that to understand the exceptional thermoelectric

properties found at high temperature in our Cu2Se pellets, it is necessary to clarify

the structural characteristics of sintered β-Cu2Se. For this reason, we perform the

heating experiment inside a TEM. Figure 3d is a bright-field TEM image taken at 450

K, in which the small grains can be observed. Figure 3d inset is an electron

diffraction pattern taken from the marked nanoparticle. The feature of the [110] zone-

axis diffraction pattern taken from an FCC structure indicates that the heating has

resulted in the formation of β-Cu2Se while the nano-sized grains can be preserved

(Figure S2). Figure 3e is a bright-field TEM image taken from the same area when

the TEM sample is cooled to the room temperature, which confirmed the

preservation of nanograins during heating and cooling process. More nano-sized

features can be observed from other areas (Figure S3) in which Moiré fringes and

strain contrast can be clearly seen. Figure 3f is a [ ̅11] high-resolution TEM image

taken from two adjacent grains, and shows they are well-crystallized with a grain

- 118 -

boundary where Moiré fringes are also shown. The insets are fast Fourier transform

(FFT) patterns of two adjacent grains showing they are slightly misorientated. Figure

3g is the reversed FFT image filtered by ±211* reflections, in which dislocations

(marked) can be clearly seen. Based on the observed distance between dislocation

cores (~5 nm), the dislocation density is very high within the high-density grain

boundaries. Taking both Figure 3f and 3g into consideration, the grain boundaries

found in our sample are small-angle grain boundaries, accommodated by a high-

density of dislocations to release misfit strained caused by the misorientation

between adjacent grains. In fact, according to Figure 3f and Figure S3, the grain

misorientation found in our sintered pellets is caused by the stacking of the plate-like

nanoparticles during the SPS process.

- 119 -

Figure 3 (a) XRD patterns of sintered Cu2Se pellet; (b) SEM image of sintered

Cu2Se pellet with inset being (1) high-magnification SEM image showing the nano-

sized feature and (2) back-scattered SEM image showing no detectable pores and

secondary phases; (c) statistic results of electron probe micro analysis for 5 different

samples; (d) TEM image of sintered Cu2Se taken at 450 K and inset being a SAED

pattern showing a [110] zone axis of FCC structure; (e) corresponding TEM image

taken at room temperature showing no size change during the cooling process and

inset being SAED pattern showing a [111] zone axis of monoclinic structure; (f)

HRTEM image showing a grain boundary with insets being FFT patterns for

individual grains; (g) reversed FFT image showing dislocation cores.

Based on our extensive structural characterizations both at the room temperature

and high temperature, it is clear that the nano-sized and plate-like grains have been

preserved without observable crystal growth during the sintering process (Figure 3,

Figure S2 and S3). This conclusion is also supported by our XRD investigations from

sintered pellets taken before and after the thermoelectric property measurements, in

which the broadening feature in the XRD patterns preserved after several cycles of

heating-cooling experiments (Figure S2). On this basis, our extensive structural

characterizations confirm that the combination of the nanoplate nature of as-

synthesized Cu2Se and the plate-stacking nature found in the sintered pellets leads

to the formation of a high-density of small-angle grain boundaries accommodated by

a high-density of dislocations in our sintered Cu2Se pellets. These structural defect

features can significantly enhance the phonon scattering.34 According to the

frequency-dependent description of κL,35 the contribution of κL is the sum of phonons

with low, high and intermediate frequencies. It is of interest to note that, κL of bulk

materials is primarily (up to 80%)35 contributed by the intermediate frequency

phonons36 with the mean free path of a few hundreds of nanometers. These

intermediate frequency phonons are also believed to contribute most of the lattice

heat transport in our sintered Cu2Se pellets, because the high-mobility and high-

disorder Cu+ ions can efficiently scatter phonons with shorter mean free paths.13

However, in our sintered Cu2Se, nanostructures with such small grains and high

densities of small-angle grain boundaries and dislocations in the grain boundaries,

these intermediate frequency phonons will be strongly blocked, as well as those low-

frequency phonons,36 so that the overall reduced κ can be achieved. Figure 4

- 120 -

illustrates the significantly enhanced phonon scattering achieved in our case, in

which high-density small-angle grain boundaries with high-density of dislocations

within the grain boundaries can efficiently block phonons transition with long 36 and

intermediate mean free paths, and Cu+ ions strongly scatter phonons with short

mean free path. This full-spectrum phonon scattering nanostructure has minor

influence to the electron behaviors because electrons have very short mean free

path which can transport through grains,37 as demonstrated in Figure 4. As a

consequence, κL is remarkably reduced, which is the major contribution to secure a

very high ZT in our sintered Cu2Se nanostructures.

Figure 4 Schematics of the phonon scattering mechanism for β-Cu2Se.

Conclusions

In summary, β-Cu2Se nanoplates and plate-like nanostructures with controlled

stoichiometry have been synthesized via a facile solution method. During the SPS

process, the stack of plate-like nanostructures leads to the formation of high-density

small-angle grain boundaries accommodated by a high-density of dislocations. This

structural feature has efficiently enhanced the thermal scattering, which in turn leads

to an enhanced ZT value of 1.82, measured at 850 K. This study provides a strategy

- 121 -

to further enhance thermal scattering of thermoelectric materials to ultimately

enhance their thermoelectric performances.

Acknowledgements

This work was financially supported by the Australian Research Council, ZGC thanks

QLD government for a smart state future fellowship (2011002414). LY thanks the

China Scholarship Council for providing his PhD stipend. The Australian Microscopy

& Microanalysis Research Facility is acknowledged for providing characterization

facilities.

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- 125 -

Supporting Information

High-Performance Thermoelectric Cu2Se

Nanoplates through Nanostructure

Engineering

Lei Yang, Zhi-Gang Chen*, Guang Han, Min Hong, Yichao Zou, and Jin Zou*

L. Yang, Dr Z.-G. Chen, Dr G. Han, M. Hong, Y. Zou, Prof. J. Zou.



Prof. J Zou


QLD 4072, Australia


- 126 -

Figure S1 (a) ZT values measured for 5 cycles from our sintered pellets; (b) ZT

values measured for 5 different samples.

- 127 -

Figure S2 High magnification TEM images show the nano-sized grain of Cu2Se

sample at (a) 450 K and (b) room temperature.

- 128 -

Figure S3 (a) Bright-field TEM image of sintered α-Cu2Se showing the high-density

of crystal grains and (b) corresponding 211* dark-field TEM image showing nano-

sized grains with different sizes ranging from several nanometers to tens of

nanometers; (c) TEM image showing plate-like stacks with plates thicknesses < 50

nm with Moiré fringes seen in some regions.

- 129 -

Figure S4 Measured (a) thermal diffusivity and (b) specific heat capacity as a

function of temperature.

- 130 -

5.2.2 Enhanced Thermoelectric Performance of Nanostructured Bi2Te3 through

Significant Phonon Scattering

Enhanced Thermoelectric Performance of

Nanostructured Bi2Te3 through Significant Phonon

Scattering

Lei Yang,a Zhi-Gang Chen,a* Min Hong,a Guang Han,a and Jin Zou a ,b*

a. Materials Engineering, The University of Queensland, Brisbane, QLD 4072, Australia.

b. Centre for Microscopy and Microanalysis, The University of Queensland, Brisbane, QLD 4072,

Australia.

KEYWORDS: Nanostructure engineering, Bi2Te3, low thermal conductivity,

enhanced thermoelectric properties

- 131 -

Abstract

N-type Bi2Te3 nanostructures were synthesized using a solvothermal method and in

turn sintered using sparking plasma sintering. The sintered n-type Bi2Te3 pellets

reserved nano-sized grains and showed an ultra-low lattice thermal conductivity (~

0.2 Wm-1K-1), which benefits from high-density small-angle grain boundaries

accommodated by dislocations. Such a high phonon scattering leads an enhanced

ZT of 0.88 at 400 K. This study provides an efficient method to enhance

thermoelectric performance of thermoelectric nanomaterials through nanostructure

engineering, making the as-prepared n-type nanostructured Bi2Te3 as a promising

candidate for room temperature thermoelectric power generation and Peltier cooling.

- 132 -

Introduction

Solid-state thermoelectric cooling and power generation can directly convert

between heat and electricity without any emissions or vibrational parts,1-5 offering the

opportunity to overcome the upcoming energy crisis. To achieve high-efficiency

energy conversion, extensive progress has been made to improve the thermoelectric

performance, which governed by the dimensionless figure-of-merit ZT, defined as ZT

= S2σT/ κ = S2σT/ (κe + κl), where σ is the electrical conductivity, S is the Seebeck

coefficient, T is the absolute temperature, and κ is the total thermal conductivity

including the contributions from electron (κe) and lattice (κl).2, 6-8 Intrinsically, an

overall high ZT needs a large power factor (S2σ) and/or a low κ. However, these

transport properties (σ, S and κ) of thermoelectric materials are highly

interdependent and conflicted with each other, which make it a challenge to optimise

them to obtain an enhanced ZT.9-11 Up to now, band engineering, including band

convergence,2 quantum confinement,12, 13 tuning effective mass,14 and distorting the

density of states,15 have been extensively employed to improve S2σ to achieve a

high ZT, while another strategies, such as nanostructure engineering16, 17 or

hierarchical architecturing,6 have been adopted to reduce κ.18

As one of the best thermoelectrics at room temperature range,5, 17, 19-25 Bi2Te3 is a

narrow band gap (~ 0.15 eV) semiconductor26 with high valley degeneracy and

anisotropic effective mass,26, 27 resulting in an intrinsically high σ and S. Bulk Bi2Te3-

based materials have been reported with recorded high ZT via introducing doping

elements or ternary phase19, 22, 28-31 to further increase σ and S, from which the

highest S2σ with 4.7× 10-3 W m-1K-1 has been obtained by doping or alloying Bi2Te3

with Se31 or Sb.22 However, the relatively high κ of Bi2Te3-based bulk materials has

become the drawback to achieve higher ZT in bulk materials.28 Recently, low-

dimensional Bi2Te3 nanostructures20, 24, 25, 32 have been developed to target even

high ZT according to the theoretical predictions on quantum confinements12, 13, 33 to

enhance S2σ and nanostructuring to reduce κl.34, 35 Additionally, bulk Bi2Te3-based

materials are anisotropic thermoelectrics,28, 30 while the nanostructured Bi2Te3

materials tend to have isotropic properties20, 25 because of the random stacking of

nano-sized grains. A low κl ~ 0.3 Wm-1K-1 has been achieved in nanostructured

Bi2Te3 materials,17, 32 but it is still crucial to clarify the relationship between Bi2Te3

- 133 -

microstructure and the increased phonon scattering to fully understand the

mechanism of κl reduction in nanostructured Bi2Te3.

In this study, nanostructure engineering was employed to enhance the thermoelectric

performance of nanostructured pure Bi2Te3. Plate-like Bi2Te3 nanostructures were

synthesized via a solvothermal method, and then sintered by sparking plasma

sintering (SPS) with a short period of time to avoid grain growth. During the SPS

process, a high density of small-angle grain boundaries accommodated by a high-

density of dislocations is formed, which strongly scatter the phonons and in turn

significantly reducing κl. As a consequence, an enhanced ZT with a peak value of

0.88 at ~ 400 K is obtained from sintered sample. Such a value represents one of the

highest reported ZT value for n-type nanostructured pure Bi2Te3.

Experimental

Analytical grade bismuth oxide (Bi2O3, 99.9%), tellurium dioxide (TeO2, 99.999%),

sodium hydroxide (NaOH, 99.99%), ethylene glycol, polyvinylpyrrolidone (PVP,

average molecular weight: 40,000) were purchased from Sigma-Aldrich and used as

precursors without any further purification.

The detailed synthesis procedure is outlined as follows. Firstly, 0.2 g PVP was

dissolved in 18 mL ethylene glycol to form a clear solution, followed by the additions

of 0.2330 g Bi2O3 powders, and 0.2396 g TeO2 powders. The prepared solution was

then mixed with 2 mL NaOH solution (5mol/L), the resulting suspension was stirred

vigorously for 30 min, and subsequently sealed in a 125 mL Teflon-lined steel

autoclave. The autoclave was heated to 210 °C for 24h and then naturally cooled to

room temperature in air. The synthesized products were collected by a high-speed

centrifugation and washed by the distilled water and absolute ethanol, and finally

dried at 50 °C for at least 12 hours.36, 37


characterized by X-ray diffraction (XRD), recorded on an X-ray diffractometer (Bruker

D8 Advance), equipped with graphite monochromatized, Cu Kα radiation (λ = 1.5418

) was used). The morphological, structural, and chemical characteristics of as-

synthesized products and sintered pellets were investigated by scanning electron

microscopy (SEM, JEOL 7800, operated at 5 kV for normal SEM and 15 kV for back-

- 134 -

scattered SEM) and transmission electron microscopy (TEM, Philips Tecnai F20,

operated at 200 kV).

The as-synthesized Bi2Te3 powders were compressed by SPS under 50 MPa and

heated at 550 K for 5 min in vacuum. The Archimedes measured were performed to

determine the density (d) and relative density (95%) of sintered pellets.

The thermoelectric properties of sintered pellets were studied in both parallel (∥) and

perpendicular (⊥) to the press direction. κ was calculated through κ= DCpd, where D

and Cp, are the thermal diffusivity and specific heat capacity, respectively. D was

measured by a laser flash method with a LFA 457 (NETZSCH). A DSC 404 F3

(NETZSCH) was used to measure Cp. σ and S were measured simultaneously on a

ZEM-3 (ULVAC). The uncertainty of the all measurements (S, σ and D) is estimated

as ~5%. The combined uncertainty for the experimental determination of ZT is up to

20%, and the standard deviation of the measured ZT from several different samples

is ~3%.


Figure 1a shows a typical XRD pattern of as-synthesized products, which can be

indexed exclusively as a rhombohedra structured Bi2Te3 phase with lattice

parameters of a = 4.386 Å and c = 30.478 Å and a space group of R ̅m (JCPDS

#15-0863).25, 36, 37 There is no other diffraction peaks can be observed, indicating the

high purity of the as-synthesized Bi2Te3. Figure 1b is a typical SEM image taken from

as-synthesized Bi2Te3, in which hexagonal plate-like nanostructures can be

observed. The lateral size distributions of these nanostructures are varied from 100

to several hundreds of nanometres. Their typical thickness can be observed in the

high magnification SEM (Figure 1c), which is around 20 nm. The crystal structure of

Bi2Te3 nanoplates are further examined by TEM (Figure 1 d-f). Figure 1d is a TEM

image of a typical hexagonal-shaped Bi2Te3 nanostructure. From the high resolution

TEM image (Figure 1e), the measured periodic fringe spacing of 0.22 nm

corresponds to the lattice spacing between the (112̅0) planes, which can be further

confirmed by the selected area electron diffraction (SAED) pattern (Figure 1f). The

clear lattice fringes in Figure 1e indicate that the nanostructure is well-crystallized.

- 135 -

Figure 1 Characterization of as-synthesized Bi2Te3 nanostructures: (a) XRD; (b) Low

magnification SEM image; (c) High magnification SEM showing the thickness of a

typical plate-like nanostructure; (d) TEM image; (e) High resolution TEM image; (f)

[0001] zone-axis SAED pattern.

To measure their thermoelectric properties, the as-synthesized Bi2Te3 nanostructures

were sintered using SPS to obtain disc-like pellets. Their thermoelectric properties

were measured in the temperature range from 300 K to 550 K. As mentioned above,

the bulk Bi2Te3 thermoelectrics showed anisotropy properties. To clarify the nature of

isotropy of our nanostructures for their thermoelectric properties, the samples were

measured from both parallel (∥) and perpendicular (⊥) to the press direction. Figure 2

- 136 -

shows various measurement results. As can be seen in Figure 2a, the sintered

Bi2Te3 shows similar σ⊥ and σ∥. The highest σ of 7.2×104 S m-1 at 300K is comparable

to reported results for pure Bi2Te3,38, 39 and then keeps decreasing with increasing

the temperature. The measured S (Figure 2b) shows negative values, indicating an

n-type pellet. The S values determined from different directions are also isotropic,

reached the peak value of -143 μV K-1 at ~ 400 K, which is comparable to the

reported results.38 To obtain κ, D (Figure 2c) and Cp (Figure 2d) were measured. The

sintered pellets have similar D∥ and D⊥, and so for the Cp, leading to similar κ values

(Figure 2e). The κ values of sintered pellets are between 0.58 Wm-1K-1 and 0.86 Wm-

1K-1, which is significantly lower than those of pure Bi2Te3,40-41 but is comparable to

the best reported Bi2Te3-based materials.20, 32 To investigate κe and κl, κe was

calculated using κe = LσT, where L is the Lorenz number.42 Here, L = 1.5×10-8 V2K-2

is used for estimating κe,38, 41 and the obtained κe is plotted in Figure 2f. Using κl = κ -

κe, the correspondingly κl can be obtained, plotted in Figure 2e. From which, an ultra-

low κl between 0.2 Wm-1K-1 and 0.37 Wm-1K-1 can be obtained, indicating that the

phonons have been strongly scattered. Benefiting from such a low κ, the ZT⊥

reached the peak value of 0.88 at 400 K, while the ZT∥ reached almost the same

value (Figure 2g). The ZT⊥ value is very stable after several cycles of measurement

(as shown in Figure 2h) and similar ZT values can be obtained for 5 different

samples (Figure 2i), suggesting that the sintered samples are highly stable and

durable.

- 137 -

Figure 2 Plots of temperature dependent thermoelectric properties of sintered

Bi2Te3: (a) Electrical conductivity; (b) Seebeck coefficient; (c) Thermal diffusivity; (d)

Specific heat values; (e) Thermal conductivity; (f) κ include the contribution of κe and

lattice κl; (g) Calculated ZT values; (h) ZT measured for 5 cycles, and (i) ZT

measured for 5 samples.

To understand the fundamental reason for such a low κl, detailed structural

characterizations were performed on the sintered pellets. Figure 3a is a XRD

pattern, which can be again indexed as rhombohedral structured Bi2Te3 without any

impurities. Figure 3b is a typal SEM image and shows that the nano-sized features

were preserved in the sintered Bi2Te3, and the Bi2Te3 nanostructures showed a

random stacking with each other. From the back-scattered SEM image of the

- 138 -

polished sample (Figure 3b inset), no secondary phase and pores can be observed,

which confirms that the sintered pellets are high purity and dense. Figure 3c and d

are typical TEM images of sintered Bi2Te3. Figure 3c shows that Bi2Te3

nanostructures can stack to each other with the average thickness of the plate-like

grains being approximately 20 nm, resulting in a high density of stacked grains. It

should be noted that such a thickness is close to the original Bi2Te3 nanoplates,

indicating that no significant grain growth occurring during the SPS process. Figure

3d shows nanosized grains with clear gain boundaries, suggesting the random

stacking of the Bi2Te3 nanostructures in our pellets.

To better understand the structural characteristics at grain boundaries, high-

resolution TEM (HRTEM) investigation was employed. Figure 4a is a TEM image

showing several grains stacked together. Figure 4b and c are HRTEM images taken

from inside a grain and a grain boundary, respectively. Figure 4b shows the grain is

well-crystalized with measurable periodicities of lattice spacings of 1 nm and 0.37

nm, which respectively correspond to the lattice spacing between the (0003) planes

and (10 ̅1) planes. Figure 4c show a clear grain boundary taken from two adjacent

grains with two insets showing the fast Fourier transform (FFT) patterns of the two

grains. As can be seen, the upper grain shows clearly lattice image, precisely viewed

along the [112̅0] direction and confirmed by the inset FFT pattern. In contrast, the

lattice image of the bottom grain is relatively faint, suggesting there exists

misorientation between the two grains, which can be further confirmed by the

difference of two inset FFT patterns. Figure 4d is another example, and the inset is

the reversed FFT images filtered by ±0001* reflections, in which two dislocations can

be clearly seen. A measured misorientation between these grains is about 4.5

degree, which can be believed to belong to small-angle grain boundaries. Their

structural model is illustrated by Figure 4e.43 The observed small grain misorientation

should be caused by the stacking of the plate-like nanostructures under high

pressure (50 MPa) during the SPS process.

- 139 -

Figure 3 (a) XRD pattern for sintered Bi2Te3; (b) SEM image of sintered Bi2Te3 with

inset of back-scattered SEM image of polished sample; (c) TEM image of sintered

Bi2Te3 showing the stacking of nano-sized grains; (d) TEM image of sintered Bi2Te3

showing nanosized grains with clear grain boundaries.

- 140 -

Figure 4 (a) TEM image of sintered Bi2Te3 showing nanosized grains and grain

boundaries; (b) HRTEM image showing clear crystal lattice within the grain; (c)

HRTEM image showing the grain boundary with the inset FFT patterns showing

slightly misorientation between two grains; (d) HRTEM image and reversed FFT

image showing dislocation cores; and (e) Schematic showing the formation of small

angle grain boundary with high density of dislocations.

On the basis of above extensive structural characterizations and analysis, we

propose a following mechanism for such a low κ, as illustrated in Figure 5. Firstly, as

demonstrated in Figure 3, the nano-sized Bi2Te3 grains have been well preserved

- 141 -

without significant grain growth during the sintering process. The stacking of Bi2Te3

nanostructures under such a high pressure during the SPS process leads to the

formation of a high-density of dislocations accommodated in the small-angle grain

boundaries of the sintered Bi2Te3 pellets. Such a high density of structural defect

features can significantly enhance the phonon scattering in materials.43, 44 In general,

the transport of phonons with low, intermediate and high frequencies contribute to

the κl according to the frequency-dependent description of κl.34 Especially, the

intermediate frequency phonons35 are believed to contribute the most of κl.34 In our

sintered Bi2Te3 pellets, the existence of fine-grain nanostructures and a high density

of dislocations accommodated in a high density of small-angle grain boundaries can

strongly block the intermediate frequency and low-frequency phonons with the mean

free path of a few hundreds of nanometers or larger,34, 35 to remarkably reduce κl and

to achieve an overall low κ. It should be noted that point defects often play an

important role for the high frequency phonon scattering,44 and TeBi anti-site defects

are often found in the as-prepared Bi2Te3 that may contribute to the intrinsic n-type

conductivity.45 Therefore, their contributions should not be ignored. Interestingly, our

sintered pellets have shown comparable electrical transport properties with reported

Bi2Te3 thermoelectrics, indicating that our sample may do not introduce significant

change in point defects compared with the reported Bi2Te3 thermoelectrics. However,

high density small-angle grain boundaries accommodated by a high-density of

dislocations were observed in our as-prepared Bi2Te3 pellets. As a consequence, we

considered that the significantly reduced κl found in our pellets should be attributed

to these high-density of small-angle grain boundaries accommodated by a high-

density of dislocations. As illustrated in Figure 5, the scattering of phonons transition

with long and intermediate mean free paths was significantly enhanced by a high-

density of small-angle grain boundaries with a high-density of dislocations, providing

a full-spectrum phonon scattering nanostructure. Consequently, κl is remarkably

reduced, which is the major contribution to secure a high ZT in our sintered Bi2Te3

nanostructures.

- 142 -

Figure 5 Schematics of the phonon scattering mechanism for sintered Bi2Te3.

Conclusion

In summary, hexagonal plate-like Bi2Te3 nanostructures with uniform morphology

have been synthesized by using a facile solvothermal method. After the sintering, a

high-density of small-angle grain boundaries accommodated by a high density of

dislocations is formed due to the stack of plate-like Bi2Te3 nanostructures during the

SPS process. These structural features reduce the overall κ, and in turn lead to an

enhanced ZT of 0.88 at 400 K. This study suggests a strategy to further enhance the

phonon scattering of thermoelectric materials to ultimately enhance their

thermoelectric performances.

AUTHOR INFORMATION

Corresponding Author

*Email: [email protected] (J.Z)

*Email: [email protected] (Z.G.C)

mailto:[email protected]

mailto:[email protected]

- 143 -

NOTES

The authors declare no competing financial interests.

ACKNOWLEDGEMENTS

This work was financially supported by the Australian Research Council, ZGC

thanks QLD government for a smart state future fellowship (2011002414). LY thanks


Microscopy & Microanalysis Research Facility and the Queensland node of the

Australian National Fabrication Facility are acknowledged for providing


- 144 -

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- 149 -

5.2.3 Manuscript

Bi-doped PbTe nanocubes with enhanced

thermoelectric properties

Lei Yang1, Zhi-Gang Chen1*, Guang Han1, Lihua Wang2, Deli Kong2, Liqing Huang1,

Yichao Zou1, Min Hong1and Jin Zou1, 3*

1Materials Engineering, The University of Queensland, Brisbane, QLD 4072,

Australia

2Institute of Microstructure and Properties of Advanced Materials, Beijing Key Lab of

Microstructure and Property of Advanced Material, Beijing University of Technology,

Beijing, China

3Centre for Microscopy and Microanalysis, The University of Queensland, Brisbane,

QLD 4072, Australia

*E-mail: [email protected], [email protected]

- 150 -

Abstract

Bi-doped PbTe nanocubes with controllable doping levels were synthesized using a

facile solvothermal method, the impacts of Bi doping into the PbTe have been

studied in details. Morphologically, high Bi doping concentration (x=0.05) for the Pb1-

xBixTe was found to induce a <100> dominant growth mechanism to form a miss-

cornered cubic nanostructure instead of the original <111> dominant growth for un-

doped and Bi-doped PbTe with lower doping level. In terms of thermoelectric

properties, Bi doping effectively suppressed the bipolar effect of un-doped PbTe,

significantly improved the electrical transport properties of PbTe. The as-sintered Bi-

doped PbTe materials also show very low lattice thermal conductivity due to the high

density of grain boundaries and strained defects via nanostructure engineering,

leading an overall ZT ~1.35 at 675 K of Pb1-xBixTe (x=0.01).

- 151 -

Introduction

Thermoelectric materials1-5 are promising to tackle the global energy crisis by

converting waste heat directly into electricity, achieving the solid-state and emission

free power generation and refrigeration without any moving parts. The thermoelectric

efficiency is governed by figure of merit of thermoelectric materials which is defined

as ZT=S2σT/κ, where S is the Seebeck coefficient, σ is the electrical conductivity, κ

is the thermal conductivity and T is the absolute temperature. Since the S, σ and κ

are highly interdependent and conflict,6 they have to be optimized to obtain high ZT.

Lead telluride2, 3 (PbTe) is one of the best thermoelectric candidates in intermediate

temperature (400-800 K) with an intrinsically high ZT of approximately 0.8.6, 7

Extensive efforts have been made to improve the thermoelectric performance of

PbTe including doping PbTe with various element to modify its band structure,3, 8, 9

carrier concentration10 and defects or interfaces,11, 12 thus the S, σ and κ of PbTe can

be tuned to improve the thermoelectric efficiency.

Bismuth (Bi) is an important dopant for PbTe as a donor impurity.13-15 In PbTe matrix,

Bi ions can replace the Pb2+ or take the Te position to form the anti-site defect16, 17 as

well as being interstitials,16 which is expected to tune the PbTe as a n-type

semiconductor. Experimentally, Bi had been doped into bulk PbTe15, 16, 18, 19 and

enhanced their thermoelectric properties.16, 18, 19 However, complex Pb-Bi-Te ternary

compounds14 were found in the doped bulk PbTe, in other cases, Bi secondary

phase and/or Bi-rich precipitates16, 18 were observed in PbTe matrix instead of

uniformly doping in PbTe, therefore, it is necessary to carefully investigate the effect

of Bi-doping to improve the thermoelectric performance of PbTe. In recent decade,

nano-sized PbTe-based materials have been extensively fabricated via various

methods with enhanced thermoelectric performance.20-23 According to the theoretical

and experimental studies, the introducing of substantial numbers of grain interface of

PbTe nanostructures can significantly scatter the phonons11, 24, 25 and can also

improve the energy filtering effect for electrons,26, 27 which could achieve a harvest of

a very low thermal conductivity approaching the amorphous limit of PbTe. By

combining such advantages, a significant enhancement of thermoelectric

performance of PbTe can be expected due to the improved electrical transport

properties via Bi doping and low thermal conductivity via nanostructure engineering.

- 152 -

In the present study, Bi was uniformly doped into PbTe nanocubes via a facile

solvothermal method. The products with the nominal compositions of Pb1-xBixTe

(x=0, 0.005, 0.01, 0.02, 0.05) were carefully characterized. The cube-shaped

products have average sizes ~120 nm, Bi was confirmed to be doped into PbTe by

various analysis methods and started to affect the morphology when the

concentration reached x=0.05, which lead to a miss-cornered cube-like morphology

due to <100> dominant growth mechanism. After SPS process, the as-synthesized

samples were densified and their thermoelectric properties were measured. The

rapid SPS process preserved the nano-sized grains of the samples, created high

density of grain boundaries and strained defects, leading to reduced κ due to the

significantly enhanced phonon scattering. The Bi doping effectively improved the σ of

samples, also suppress the bipolar conduction to stabilize the S, resulting a high ZT

~1.35 at 675 K for the Pb1-xBixTe x=0.01sample.

Experimental

Analytical pure sodium telluride (Na2TeO3, 99.999%), lead oxalate (PbC2O4,

99.999%), bismuth chloride (BiCl3, 99.999%), ethylene glycol, polyvinylpyrrolidone

(PVP, average molecular weight: 40,000) and sodium hydroxide (NaOH, 99.99%)

were purchased from Sigma-Aldrich and used as precursors without any further

purification.

In a typical synthesis of PbTe nanostructures, 0.2g of PVP was dissolved in 36mL of

ethylene glycol, and then 0.1108g of Na2TeO3, 0.1476g of PbC2O4 and 4 mL 5mol/L

NaOH solution were added in with continuous stirring (proportional BiCl3 was added

for Bi-doped PbTe samples with the reduced amount of PbC2O4). The solution was

put into a 125mL Teflon-lined stainless steel autoclave and sealed and then heated

at 230 °C for 4 hours. After that, the autoclave was cooled to room temperature

naturally. The products were collected by centrifuging and washed by deionized

water and absolute ethanol for several times, and then dried at 60 °C for at least 12

hours.


characterized by XRD, recorded on an X-ray diffractometer equipped with graphite

monochromatized, Cu Kα radiation (λ = 1.5418 ). The morphological, structural,

and chemical characteristics of as-synthesized products and sintered pellets were

- 153 -

investigated by SEM (JEOL 7800, operated at 5 kV) and TEM (Philips Tecnai F20,

operated at 200 kV, Philips Tecnai F30, operated at 300 kV). A JEOL JXA-8200

(operated at 20 kV) was used for the electron probe micro analysis (EPMA).

The as-synthesized Bi-doped PbTe powders were compressed by SPS under 60

MPa and heated at 673 K for 5 min in vacuum. The Archimedes measured were

performed to determine the density (d) and relative density (90%).

Thermal conductivity κ was calculated through , where D and Cp, are the


laser flash method with a LFA 457 (NETZSCH). A DSC 404 F3 (NETZSCH) was

used to measure Cp of Cu2Se. σ and S were measured simultaneously on a ZEM-3

(ULVAC). The uncertainty of the all measurements (S, σ and D) was ~5%, and the

uncertainty for the measured Cp was ~10%. The standard deviation of the measured

ZT was ~5%.


The XRD patterns of Pb1-xBixTe (x=0, 0.005, 0.01, 0.02, 0.05) samples are shown in

Figure 1a. As it can be seen in Figure 1a, all the samples can be indexed as face-

centre-cubic PbTe (Standard Identification Card, JCPDS 78-1905, the orange lines

in Figure 1a)22, 28 without any impurities or secondary phases. Right shifts of peaks

can be observed with the Bi doping as it is shown in Figure 1a inset, indicating there

were lattice shrinkages when Bi was doped in to PbTe. The lattice shrinkage

gradually increased (Figure 1a inset) with the increased Bi doping level, which

reveals the substitution of Pb2+ by the Bi3+ as Bi3+ has smaller ionic radius compare

to Pb2+.13

pDC d

- 154 -

Figure 1 XRD patterns of Bi-doped PbTe (Pb1-xBixTe) with different doping levels as

marked, the inset shows the peak shift of (200) peaks.

The morphologies of the as-prepared samples have been verified by SEM and

shown in Figure 2. From Figure 2, all the Bi-doped PbTe samples show cubic

morphologies and have very uniform size distribution with average sizes around 120

nm. Interestingly, all the as-prepared Bi-doped PbTe nanostructures are intact cubes

(Figure 2a-d and insets) except the x=0.05 sample (Figure 2 e and f) which shows

cubes without eight corners. Such different morphology may due to different growth

mechanism, which will be carefully investigated later.

- 155 -

Figure 2 SEM images of Pb1-xBixTe samples with different Bi doping level: (a) x=0;

(b) x=0.005; (c) x=0.01; (d) x=0.02; (e) and (f): x=0.05, the insets are high

magnification SEM images for each sample.

TEM analysis was applied to study the crystal structure of as-prepared samples. A

typical TEM image of un-doped PbTe was shown in Figure 3a as an example of

intact nanocubes. In Figure 3a, the nanocubes show the size contribution from ~100

nm to ~150 nm, which is consistent with the SEM results. The contrast in the TEM

image is thickness contrast due to tilted nanocubes. The high resolution TEM

(HRTEM) image of un-doped PbTe is shown in Figure 3b, in which a lattice spacing

~0.23 nm can be measured, corresponding to the 220* lattice spacing of FFC PbTe.

The FCC crystal structure can be confirmed by the corresponding SAED pattern in

Figure 3b inset which was taken from the [111] zone axis. The x=0.05 sample

(Figure 3c) shows rounder projection shape compared with un-doped sample due to

the missed corners and they have similar size as which is consistent with the SEM

- 156 -

results. Figure 3d shows the HRTEM of x=0.05 sample, in which a lattice spacing of

~0.32 nm can be obtained, corresponding to the 200* lattice spacing of PbTe. The

SAED pattern from [001] zone axis confirmed the cubic crystal structure, revealing

that there is no structural changed after Bi was doped into PbTe although the

morphology has been slightly modified.

Figure 3 (a) Typical TEM image of un-doped PbTe nanocubes; (b) HRTEM image of

un-doped PbTe and the corresponding SAED pattern; (c) typical TEM image of Pb1-

xBixTe x=0.05 nanostructures and its HRTEM image (d) with the inset of

corresponding SAED pattern.

To study the growth mechanism, the Pb1-xBixTe samples have been synthesized in

different reaction time and their SEM images are shown in Figure 4. The un-doped

PbTe and x=0.05 samples were chosen to demonstrate the typical growth process.

According to the experimental results, the chemical reactions start to take place after

- 157 -

1 hour of heating process, so that the samples were collected from 1.5 hours.

Octahedral crystal seeds were found in 1.5 h un-doped sample as it is can be seen

in Figure 4a. The octahedrons further developed and form tetrakaidecahedrons

which can also be observed in Figure 4a and was modelled in Figure 4a inset. The

nanostructures kept growing and form a further developed tetrakaidecahedrons at 2

hours as it is shown in Figure 4b, and finally grew to the fully developed nanocubes

after 3 hours. The x= 0.05 sample shows the similar morphology at the early stage of

the reaction (1.5 h, Figure 4c), but it grew to a significantly different shape at 2 hours

as it can be observed in Figure 4e, and the miss-cornered nanostructures can be

obtained after 3 hours. Our reaction time for synthesizing Bi-doped PbTe

nanostructures was set to 4 hours to allow the reaction fully competed.

Figure 4 Un-doped PbTe samples synthesized in (a) 1.5 h; (b) 2 h; (c) 3h; and Pb1-

xBixTe x=0.05 samples synthesized in (d) 1.5 h; (e) 2 h; (f) 3h with the insets of

models.

- 158 -

According to the SEM results, the growth mechanism of PbTe nanocubes and

x=0.05 miss-cornered nanostructures can be verified. The shape of nanomaterials

with FCC crystal structure determined by the growth rate ratio in the <111> and

<100> crystalline directions.29 At the early stage of the reaction, both samples form

octahedral crystal seeds. All eight facets of such octahedrons should be 111* facets

considering the FCC crystal structure.30 For the nanocube samples, the <111>

growth dominated the crystal growth (has higher growth rate along <111>

directions), these 111* facets continue developing and the 100* facets would be

seen in the following stage (Figure 4a, Figure 5a), and they will fully developed to

form intact nanocubes. For the miss-cornered nanostructures, the growth rate along

<100> directions became higher29 due to the sufficient Bi doping, 100* facets

dominated the growth instead of 111* facets. The different processes can be

demonstrated by Figure 5a. Figure 5b is the SEM image of Pb deficient sample with

a nominal composition of Pb0.95Te, in which the nanostructures shows un-uniform

morphology but no miss-cornered structures. From the discussion above, such

different growth mechanism should be caused by the high doping level of Bi into

PbTe rather than the Pb deficiency. Actually, when the concentration of the

precursors was increased, the <100> dominated growth mechanism became more

obvious. In a control experiment, all the precursors’ concentration was increased by

10 times to synthesize the Pb1-xBixTe x=0.05 sample, and the SEM image of the

product is shown in Figure 5c. In Figure 5c, the morphology of the product is varied,

but all of them showed the <100> dominant growth shapes,29 in which some

structures with six symmetric components can be observed (as marked in Figure 5c).

Each one of these six components should be developed from six 100* facets.

- 159 -

Figure 5 (a) schematic of different growth mechanisms for nanocubes and x=0.05

miss-cornered nanostructures; (b) SEM image of Pb deficient Pb0.95Te sample does

not show the miss-cornered morphology; (c) SEM image of the x=0.05 sample

synthesized with the 10 times concentration shows some typical <100> dominant

growth structures.

The as-synthesized samples have been sintered using SPS process. The as-

sintered samples have been characterized using SEM and TEM. Only the typical

SEM and TEM images of x=0.01 sample were shown in Figure 6 as all the as-

sintered samples show similar fine grain structures. The EPMA results have been

shown in Figure 6a, which reveal good agreement with the nominal compositions.

From the SEM image in Figure 6b, grains with varied sizes from tens of nanometres

to ~200 nm can be observed, indicating there was minor crystal growth (compared to

the original sized of Bi-doped PbTe ~120 nm) during the sintering even with a

relatively low sintering temperature (673 K). However, the rapid SPS process

prevented the grains from further growth,31 so that these samples still have very

- 160 -

small grain size and high density of grain boundaries. Some pores can be observed

from the SEM image, which may due to the randomly stacked cubic nanostructures

and lead a relatively low density ~90% for the as-sintered samples. The fine grains

with varied size of as-sintered sample can also be observed from the TEM image

(Figure 6c), in which high density of grain boundaries are obvious. Figure 6d is

another TEM image, which shows the dense-sintered nanograins, also reveals the

defects within the grains (in the red cycle and enlarged in Figure 6e) with higher

magnification. The selected area electron diffraction (SAED) pattern taken from the

marked area can be indexed as the [111] zone axis of the FCC nature

- 161 -

Figure 6 (a) Statistic results of EPMA (b) SEM and (c) TEM images of as-sintered

Bi-doped PbTe samples (x=0.01) as examples to show the nano-sized grains after

sintering process;(d) a TEM image with higher magnification shows dense sintered

nano-sized grains and defects within grains, the inset SAED pattern can be indexed

as the [111] zone axis; (e) the enlarged TEM image shows the defects in the grain.

The defects within the grains were further investigated using TEM. Figure 7a is a

high magnification TEM image of a grain taken along the [111] zone axis. Such grain

contents multiple defects in the marked regions, which can be seen in their inversed

Fast Fourier Transform (IFFT) images in Figure 7b and c, respectively. In Figure 7b

and c, the cores of edge defects can be clearly observed. The inset of Figure 7b

illustrated the formation of such edge defects, in which an extra half-plane of atoms

insets through the crystal, causing regional distortions. From both Figures 7b and c,

the density of defects is very high, which could cause significant strain in the grain.

Figure 7d is the strain map of Figure 7c showing the distribution of strain, from which

a large strain up to 25% can be observed on the defect cores, leading to significant

lattice distortion nearby the defect cores. The distorted regions have an average

lateral size ~ 2nm, which can provide strong scattering to the phonons with similar

mean free path,32-34 so that significantly reduced κ can be expected for the as-

sintered samples.35

From Figure 6 and Figure 7, the rapid SPS process effectively preserved the nano-

sized grains after sintering, create not only high density of grain boundaries in the

as-sintered samples, but also strained defects accommodate within the grains

(Figure 6c and d), which is essential for reducing the κ and improve the

thermoelectric performance6, 26, 31-33, 35, 36 through increasing the phonon scattering.

- 162 -

Figure 7 (a) A typical high magnification TEM image of the defect region taken along

the [111] zone axis as can be seen in the inset FFT pattern; (b) and (c) the IFFT

images from the marked area in (a) clearly show the defect cores as marked, the

inset of (b) is a schematic of edge defect view from [111] zone axis; (d) the strain

map of (c) shows strain distribution around dislocation, the colour bar indicates 25 to

–25% strain; (d) schematic of as-sintered samples shows the high density of grain

boundaries of nano-sized grains and defects located in the grain.

The thermoelectric properties of as-sintered samples were measured from 300 K to

800 K. The obtained electrical transport properties (σ and S) are plotted in Figure 8.

Figure 8a is the temperature dependent σ values for all the as-sintered samples with

- 163 -

compositions as marked (n is the room temperature carrier concentration). It can be

seen that the un-doped PbTe has very poor σ, which was significantly improved by

Bi doping via increasing the n. With the increase of Bi doping level, the σ values

increase, while all the Bi-doped samples show a decrease with the increased

temperature. Notably, the S (Figure 8b) of un-doped PbTe sample shows dramatic

change versus the temperature: it gradually increases from 300 K to ~425 K with

positive values (p-type conductivity), then decreases and turn into a n-type

semiconductor at ~600 K, and then continually increases till 800 K to reach -210

μVK-1. The original p-type conductivity of un-doped PbTe at low temperature may

due to the slight off stoichiometry of the as-synthesized sample. The electrical

transport behaviour of un-doped PbTe is due to the bipolar conductivity and a two

valence band conduction mechanism3, 37 of PbTe, in which both holes and electrons

contribute to the σ. At room temperature, holes mainly contribute to the conductivity

for the p-type un-doped PbTe and the Fermi level is closer to the valence band due

to the p-type nature. With the increased temperature, the light valence band

degenerates,3 so the band gap will increase, leading to a decreased σ and increased

S.37 When the temperature further increase, more electrons started to be thermally

activated through the band gap, the bipolar effect become significant, leading a

decreased S and finally turn to a n-type semiconductor. The doping of Bi suppressed

such bipolar conductivity and stabilized the S, making S increases with the increased

temperature until ~725 K (Figure 8b).

Figure 8 The temperature dependent electrical transport properties of as-sintered Bi-

doped PbTe samples with various compositions: (a) σ and (b) S values.

- 164 -

The Cp and D have been measured to determine the κ. As it can be seen in Figure

9a, the measured Cp values are between 0.157 and 0.159 Jg-1K-1 from 300 K to 800

K. The measure D values are plotted in Figure 9b, and the κ can be calculated and

seen in Figure 9c. From Figure 9c, the κ of as-sintered un-doped and Bi-doped PbTe

is lower than most of bulk PbTe-based materials38 and comparable to some fine-

grained nanostructured materials.26 The κ of un-doped PbTe is lower than that of Bi-

doped PbTe, which may due to the low electronic contribution (κe) of un-doped

sample to the total κ. To exclude the κe from the κ, the lattice contributions (κL) for

the κ have been calculated using κL= κ- κe according to the Wiedemann-Franz law,39

where κe= LσT, L is Lorenz number. The L value could be varied depending on the

materials, but it can be estimated using the measured S values as L=1.5+exp[-

|S|/116],39 notably L is in 10−8 WΩK−2 while S is in μV/K in this equation. The L has

been calculated and plotted in Figure 9d as a function of temperature. As is can be

seen in Figure 9d, all the L values for un-doped and Bi-doped PbTe are between the

1.5 × 10−8 WΩK−2 for acoustic phonon scattering and the 2.44 × 10−8 WΩK−2 for

degenerate limit,39 showing a typical semiconductor behaviour.39 Then the κL can be

obtained as it is shown in Figure 9e. From Figure 9e, all the samples show very low

κL between 0.6 and 1 Wm-1K-1 when the temperature is higher than 500 K, but there

is no significant difference between each sample. Such low κL may benefit from the

introducing of mass contrast and point defects via doping with Bi,18 and the

increased grain boundaries and interfaces via nanostructure engineering.31, 36

Finally, the improved electrical transport properties of Bi-doped PbTe result in

significantly enhanced ZT compare to the un-doped sample (Figure 9f). Also

benefited from the low κ via nanostructure engineering, the as-sintered Bi-doped

PbTe samples show higher ZT values than some other n-type PbTe-based bulk

materials,38 reached a peak ZT of ~1.35 at 675 K for the x=0.01 sample (Figure 9f).

Compared with the un-doped PbTe sample (Figure 9f), the x=0.01 sample achieved

a 125% increase on the peak ZT.

- 165 -

Figure 9 Temperature dependent properties of as-sintered un-doped and Bi-doped

PbTe samples: (a) Cp; (b) D; (c) κ values; (d) the calculated Lorenz numbers; (e) κL

and (f) ZT values.

The as-sintered x=0.01sample was tested by 5 cycles and the ZT values are shown

in Figure 10, revealing very good thermal stability of the sample during the test.

There was neither any significant degeneration of the electrical transport properties

nor obvious change of the thermal conductivity.

- 166 -

Figure 10 Cycling test ZT results for as-sintered x=0.01 sample show good thermal

stability.

Conclusions

In this work, Un-doped and Bi-doped PbTe nanocubes were synthesized via a facile

solvothermal method with controllable Bi doping levels. The heavy doped (x=0.05)

sample shows a miss-cornered cubic nanostructure due to the <100> dominant

growth mechanism. The doping of Bi significantly improved the electrical transport

properties of PbTe, suppressed the bipolar conduction, leading to an increased n-

type electrical conductivity and slightly reduced S. Additionally, the SPS process

effectively preserved the nano-sized grain, which provided high density of grain

boundaries for extra phonon scattering which significantly reduced the thermal

conductivity, resulting an high ZT of 1.35 at 675 K for the x=0.01 sample.

- 167 -

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- 171 -

Chapter 6 Conclusions and

Recommendations

In this thesis, literature study has been done to establish a comprehensive view of

the state-of-art development of thermoelectric materials; metal chalcogenides have

been focused due to their intrinsically good thermoelectric performance and huge

potential which has been shown in previous theoretical and experimental studies.

Nanostructure engineering has been applied as the major strategy to enhance the

thermoelectric performance of metal chalcogenides, which was supplemented by

compositional controlling, includes doping and creating vacancies. For the

experimental work, Cu2Se-, Bi2Te3- and PbTe-based thermoelectric nanomaterials

were synthesized via facile and controllable solvothermal methods, the products

were sintered and their thermoelectric properties have been investigated in detail.

The conclusions of the thesis can be summarized as follows:

Cu2Se-based nanomaterials have been synthesized with controllable Cu

deficiency (Cu2-xSe). The Cu deficiency was found to impact not only the

structure but also the electrical transport properties: α-phase was

observed in the major β-phase Cu2-xSe when the Cu deficiency reached

Cu1.95Se; the stoichiometric Cu2Se showed the best thermoelectric

performance while the electrical transport properties can be harmed by the

Cu deficiency, leading to reduced ZT. Te was doped into Cu2Se

nanomaterials in order to tune the electrical transport properties. The as-

prepared samples showed increased σ and decreased S with the

increased Te doping level. Overall, the Cu2Se0.98Te0.02 sample reached the

peak ZT of 1.76 at ~850 K. In these studies, the electrical transport properties

of Cu2Se-based nanomaterials can be effectively controlled by both

compositional adjustment and elemental doping to achieve the optimized

results, which is very important for achieving high ZT.

As the highlighted strategy in this thesis, nanostructure engineering was

carefully investigated on different metal chalcogenides. Nano-sized Cu2Se,

Bi2Te3 and PbTe were synthesized, and sintered using SPS. The as-sintered

- 172 -

samples maintained the nano-sized crystal grains with very high density of

grain boundaries. For Cu2Se and Bi2Te3, the sintered samples have well-

crystallized grains without obvious defects, while high density of edge defect

arrays were found accommodated within the grain boundaries, forming small

angle grain boundaries. Such nano-sized grains combined with high density of

small angle grain boundaries provided the full spectrum phonon scattering

mechanism significantly reduced the κ without decrease the electrical

transport properties, resulting in enhanced ZT compared to their bulk

counterparts. For PbTe samples, Bi was doped into PbTe to improve the

power factor. High density of strained defects was found within the nano-sized

PbTe grains, which can also strongly scatter phonons and reduce the κ,

leading to a significantly increased ZT. In these studies, nanostructure

engineering was approved as an effective strategy for enhance the

thermoelectric performance via reducing the κ without decrease the power

factor.

According to the investigations which have been done in this thesis, some

suggestions and recommendations can be stated:

The effectiveness of nanostructure engineering on reducing the κ via

increasing the phonon scattering was approved in metal chalcogenides

nanomaterials, but the universality is still need to be verified. The

nanostructure engineering should be applied on more materials

systems and the mechanisms should be studied.

As the high density of grain boundaries are crucial to enhance the

phonon scattering, the effect of the density of grain boundaries and

defects to the phonon scattering should be systemically investigated.

For example, it should be clarified that what is the optimum density of

grain boundaries and defects for certain materials, or how small the

grains should achieve to obtain the lowest κ with relatively high power

factor.

More studies should be done on combining the nanostructure

engineering with band engineering and optimizing the carrier

concentration as they are all efficient strategies to obtain high ZT with

great potential.

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