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ADVANCES IN SCIENCE – Volume 1

Managing Editor: Omid Zadakbar Publisher: Science Network – Online Open Access Publisher ISBN: 978-0-9869554-0-2

Cover Designer: Mona Atari First published in March, 2013 Printed in Canada

A free online edition of this book is available at www.sciencenetwork.ca Additional hard copies can be obtained from reprint@sciencenetwork.ca

Copyright © 2013 Science Network All Books published by Science Network are licensed under a Creative Commons AttributionNonCommercial-NoDerivs 3.0 Unported License.


ADVANCES IN

SCIENCE

 Science Network Online Open Access Publisher


TABLE OF CONTENTS CHAPTER 1 CONDENSED MATTER PHYSICS ON CURVED GEOMETRY INTRODUCTION

1

1. QUANTUM TRANSPORT IN CURVED NANOMATERIALS 1-1- CURVATURE-INDUCED POTENTIAL FIELD 1-2- RESISTIVITY JUMP IN CORRUGATED FILMS 1-3- COLLECTIVE EXCITATION IN POLYMERS 1-4- ELECTRON-PHONON INTERACTION IN POLYMERS

3 3 6 7 10

2- ORIENTATIONAL ORDER ON CURVED SURFACES 2-1- CURVATURE-INDUCED FRUSTRATION 2-2- LIQUID CRYSTAL FILM ON SPHERICAL SURFACES 2-3- LIQUID CRYSTAL FILM ON A GAUSSIAN BUMP

12 12 13 15

3- AQUOS FOAM ON CURVED SURFACES 3-1- COARSENING DYNAMICS ON THE FLAT PLANE 3-2- COARSENING DYNAMICS ON CURVES SURFACES 3-3- EXPERIMENTAL TEST OF GENERALIZED VON-NEUMANN FORMULA 3-4- STABILITY IN THE EQUILIBRIUM FOAM STRUCTURE

17 17 18 20 21

SUMMARY AND PERSPECTIVE ACKNOWLEDGEMENT REFERENCES

22 23 23

CHAPTER 2 GRAPHENE-METAL INTERACTION INVESTIGATED BY SOFT X-RAY SPECTROSCOPIES 1- INTRODUCTION

32

2- SOFT X-RAY SPECTROSCOPY STUDIES ON GRAPHENE/METAL & GRAPHENE/SIO2 SYSTEMS 2-1- GRAPHENE/CU 2-2- GRAPHENE/PT 2-4- COMPARISON OF GRAPHENE ON DIFFERENT METAL SURFACES 2-5- GRAPHENE/SIO2

38 38 46 51 53

3- SOFT X-RAY SPECTROSCOPY STUDIES ON GRAPHENE OXIDE (GO) & RELATED MATERIALS 3-1- CHEMICALLY MODIFIED GO 3-2- GO-S NANOCOMPOSITE FOR HIGH PERFORMANCE LI/S CELLS

58 58 61

CONCLUSIONS ACKNOWLEDGEMENTS REFERENCES

65 66 66

I


CHAPTER 3 BIOCOMPATIBLE CONJUGATION OF HYDROGELS FOR REGENERATIVE MEDICINE INTRODUCTION

77

1- MATERIALS FOR HYDROGELS 1-1- NATURAL MATERIALS 1-2- SYNTHETIC MATERIALS

80 80 81

2- HYDROGEL SYSTEMS 2-1- PHYSICAL CROSSLINKING 2-1-1- AFFINITY INTERACTION 2-1-2- NUCLEOBASE PAIRING 2-2- CHEMICAL CROSSLINKING 2-2-1- SCHIFF-BASE REACTION 2-2-2- DIELS-ALDER ADDITION

82 82 82 86 89 89 93

CONCLUSIONS REFERENCES

98 99

CHAPTER 4 FOUR-ELEMENT PRINCIPLE OF ORGANIC/POLYMER PI-SEMICONDUCTORS 1- INTRODUCTION 1-1- ELECTRONIC STRUCTURE ANALYSIS AND DESIGN 1-2- STERIC HINDRANCE ANALYSIS AND DESIGN 1-3- CONFORMATION/TOPOLOGY ANALYSIS AND DESIGN 1-4- SUPRAMOLECULAR ANALYSIS AND DESIGN 2- SUMMARIES AND OUTLOOK ACKNOWLEDGEMENTS REFERENCES

104 109 114 117 120 38 126 127

CHAPTER 5 INTEGRATED SILICON PHOTONICS APPLIED FOR OPTICAL INTERCONNECTS INTRODUCTION

133

1- HISTORY OF SILICON PHOTONICS

136

2- PASSIVE DEVICES 2-1- WAVEGUIDES 2-2- COUPLERS

141 141 144

II


2-3- RING RESONATORS

148

3- MODULATORS

152

4- DETECTORS

158

5- LIGHT SOURCES

162

6- COMMERCIAL PROGRESS

169

CONCLUSIONS REFERENCES

172 173

III


Advances in Science

Chapter

1 Condensed Matter Physics on Curved Geometry Hiroyuki Shima Associate Professor; Department of Environmental Sciences and Interdisciplinary Graduate School of Medicine and Engineering, University of Yamanashi, 4-4-37, Takeda, Kofu, Yamanashi 400-8510, Japan E-mail :

hshima@yamanashi.ac.jp

S

urface curvature of a low-dimensional host material with non-flat geometry often leads to drastic alteration in the nature of physical systems living upon it. This chapter provides a bird’s-eye review on the latest findings of curvatureproperty relation, i.e., intriguing correlation between geometric surface curvature and physical property of the material considered. Special attention will be paid to the following three issues: quantum transport in curved nanocarbons, liquid crystal membranes with curved shape, and aqueous foam confined to curved surfaces. Theoretical approaches based on differential geometry, which facilitates quantitative description of the curvature effects, are also explained.

Introduction The phrases of “geometry" and “curvature", originally used in the mathematics community alone, are becoming commonplace in diverse fields of condensed matter physics. Nowadays, physical consequences of geometric curvature can be observed not only in spacetime distorted by a gravitational field [1,2], but also in various low-dimensional materials such as

graphene nanoribbons [3, 4, 5],

corrugated nanocarabon cylinders [6, 7], liquid crystal films, [8, 9, 10], and cell

1


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membranes [11, 12]. In all the condensed matter abovementioned, the underlying geometric curvature can affect dominantly the physical properties of the systems. From a theoretical perspective, the body of research has relied on differential geometry, the mathematical discipline that gives a quantitative description of curvature-property relations in physics. In fact, differential geometry allows us to formulate explicitly, for instance, the effective Hamiltonian of quantum excitations in curved nanomaterials [13, 14, 15]. It also enables us to appreciate beautiful interplay between surface curvature and topological defect configuration in orientationally ordered systems in two dimensions [16, 17]. In particular, curvature-property relations become salient in soft matters that are mechanically deformable; liquid crystal membranes and monolayered aqueous foam are the cases in point. This chapter provides an overview of the up-to-date findings on the curvatureproperty relations in low-dimensional materials endowed with nonzero surface curvature. Special focus is placed on the following three issues:

(i) Anomalous quantum transport in curved nanomaterials: Following an introductory guide to quantum mechanics on curved surfaces, we discuss the way how such the “curvature effect" can be manifested in real systems. In particular, the first experimental verification of such the curvature effect observed in a specific class of curved nanocarbons (i.e., peanut-shaped polymers) will be underlined.

(ii) Configurational order in liquid crystal membranes with curved shape: Non-trivial couplings between the surface curvature and topological defect configurations, which are theoretically predicted to be observed in liquid crystal membranes with curved shape, will be explained in detail. We will learn that the energetically preferred configuration is determined by the spatial distribution of the curvature.

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(iii) Time-evolution of aqueous foam confined to curves surfaces. The last topic is the coarsening dynamics of two-dimensional dry foam sandwiched by two parallel transparent sheets. The surface curvature of the paired sheets is known to dominate the fate of polygonal gas bubbles; the theoretical background as well as the first experimental test for the predicted dynamics will be demonstrated.

Throughout the discussions, we will see that geometric curvature causes a drastic alteration in the nature of the systems, which imply the development of curvedshaped functional materials based on curvature-property relations.

1- Quantum transport in curved nanomaterials 1-1- Curvature-induced potential field Rapid advances in nanotechnology have made it possible to manufacture quasi one- and two-dimensional nanostructures with non-flat geometries. One of the most intriguing realizations is a microscopic “knot" or “link" made from organic materials. The presence of knotted molecules was first identified in DNA [18], and subsequently found in naturally occurring proteins [19]. After the field of chemical topology blossomed, diverse molecular knots have become a reality [20, 21], in which some topological effects may be observed in their quantum mechanical states. Besides soft materials, a variety of hard nanomaterials with curved geometry has been synthesized, such as wrinkled

nanomembranes

[22, 23] and mechanically buckled carbon nanotubes [24]. These experimental achievements arouse renewed interest in the effect of structural geometry on the quantum-mechanical properties. Surprisingly, the quantum mechanics of a particle whose motion is constrained to curved surfaces has been an old but new problem of theoretical physics for more than 50 years. The difficulty arises from operator-ordering ambiguities [13], which permit multiple consistent quantizations for a curved system. 3


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Figure 1: Curvilinear coordinate systems based on the surface of parametric equation

One of the conventional methods used to resolve the ambiguities is the confiningpotential approach [14, 15]. In this approach, the motion of a particle on a curved surface is regarded as being confined by a strong force acting normal to the surface. Because of the confinement, quantum excitation energies in the normal direction are raised far beyond those in the tangential direction. Hence, we can safely ignore the particle motion normal to the surface, which leads to an effective Hamiltonian for propagation along the curved surface with no ambiguity. Given below is an outline of the derivation of the effective Hamiltonian in accord with the confining-potential approach. Suppose non-interacting spinless electrons confined to a general two-dimensional curved surface S embedded in a threedimensional Euclidean space. A point p on S is represented by (1) where

is a curvilinear coordinate spanning the surface and

are the

Cartesian coordinates in the embedding space. Figure 1 illustrates spatial interrelation between the two coordinate systems. Using the notation (2) 4


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Figure 2: Definition of the two principal curvatures at a point of a curved surface

We introduced the following quantities: |

(3)

|

where n is the unit vector normal to the surface. From the confining-potential approach [14, 15], it follows that the wavefunction

of a noninteracting electron

on a curved surface satisfies the Schrödinger equation defined by [

Where

(

]

√ )

[25], and

(4)

is the effective mass of electrons. The

quantities (5) are the so-called Gaussian curvature and mean curvature, respectively, both of which are functions of

. The term proportional to

in Eq. (4) is

called the effective scalar potential induced by surface curvature. Geometric meanings of the two classes of curvature are accounted for by the following discussion. Suppose at each point p on a curved surface, one choose a unit normal vector (see Fig. 2). A normal plane at p is one that contains the normal, and thus contains a unique direction tangent to the surface and cut the surface in a plane curve indicated by C in Fig. 2. This curve usually has different curvatures for different normal planes at p. Among infinite choices, the set of maximum and minimum values of this curvature are called the principal curvatures at p.

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Now we are ready; the Gaussian curvature

of the surface at p is the product of

the two principal curvatures at p. Similarly, the mean curvature

at the same

point is defined by the arithmetic mean of the principal curvatures. As seen from Eq. (4), the effective Hamiltonian involves an effective scalar potential whose magnitude depends on the local surface curvature. This means that quantum particles confined to a thin curved layer behave differently from those on a flat plane, even in the absence of any external field. Such curvature effects have gained renewed attention in the last decade, and thus many intriguing phenomena of non-interacting electron systems with curved geometry have been considered as fascinating subjects.

1-2- Resistivity jump in corrugated films In nano-scaled materials, low-dimensionality often enhances the influence of electron-electron interactions on the quantum nature of the system. An important consequence can be found in the low-temperature resistivity of two-dimensional corrugated semiconductor films [26, 27]. It was shown that the electrical resistivity of nano-corrugated semiconductor films (Fig. 3) exhibits a stepwise increase with the corrugation amplitude.

Figure 3: (Left) Schematic view of thin corrugated surface represented by z=acos(γx), with a being the corrugation amplitude and 2π/γ the period of corrugation. The curvilinear coordinates (ξ,η) and the period of a curvature-induced potential U(ξ), denoted by Λ , are indicated. (Right) Diagram of the first-order Umklapp ξ scattering process between two-electron states

6

and

. After Ref. [26].


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The stepwise increase in the resistivity is attributed to contributions of electronelectron Umklapp scattering processes. When the amplitude of corrugation takes specific values determined by the Fermi energy, then Umklapp processes due to the curvature-induced periodic potential cause a change in the total electron momentum that results in a sudden jump in the resistivity (Fig. 4).

Figure 4: Stepwise increase in Inset: contribution to

with corrugation amplitude a.

shows sudden jumps at

and 2.6.

from the normal process (m=0) and Umklapp process of the first (m=1) and second (m=2) orders. After Ref. [26].

Corrugation amplitudes that lead to resistivity jumps are determined by the relation (6) where

is the Fermi wave vector and

is the corrugation-induced reciprocal-

lattice vector associated with the curvature-driven periodic potential. The magnitude of the enhanced resistivity is within the realm of existing experiments, which confirms the relevance of the theoretical predictions to realistic applications based on curvature-property relations.

1-3- Collective excitation in

polymers

Another important consequence of geometric curvature shows up in the case where curved nanomaterials are quasi one-dimensional systems. It is well known that in one-dimension, the conventional Fermi-liquid theory breaks down so that 7


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the system is in a Tomonaga-Luttinger liquid (TLL) state [28]. In a TLL state, many physical quantities exhibit a power-law dependence stemming from the absence of single-particle excitations near the Fermi energy; this situation naturally raises the question as to how geometric perturbation affects the TLL behaviors of quasi-one-dimensional curved systems. Peanut-shaped

polymers [29, 30] are exemplary materials to be considered for

studying TLL behaviors. They are synthesized by electron-beam irradiation of pristine two-dimensional

films, having coalesced structures via the general

Stone-Wales rearrangement [31] between adjacent

molecules. The

polymer belongs to a class of π-electron conjugated system, thus exhibiting metallic properties. In addition, they are thin, long, and hollow tubules whose radius is periodically modulated along the tube axis; this simply that the polymer has both positive and negative Gaussian curvatures (see Fig. 5), which thus belongs to a novel class of nanocarbon materials distinct from other well-known π-electron systems such as carbon nanotubes [32] and graphene ribbons [33]. Hence, the periodic surface curvature intrinsic to the systems will produce sizable effects on their TLL properties. Recently, the curvature-induced alteration in the TLL nature has been explored by a collaborative project in which the author was involved. From a theoretical perspective, the author and collaborators developed a quantum theory based on differential geometry [34]. They predicted that spatial variation in the surface curvature of

polymers causes a significant increase in the power-law exponent

of the single-particle density of states; the increase in the exponent is thought to originate from a curvature-induced potential that attracts low-energy electrons to region that has large curvature [35]. Eventually, the prediction was verified in a photo-emission experiment [6].

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Figure 5: Hypothesized atomic configuration of a peanut-shaped

polymer with an uneven peanut-shaped

structure. The area colored in light blue represents a quasi-two-dimensional curved space to which the motion of π-electrons is constraint. Alternation in the sign of Gaussian curvature (designated by k in the plot) along the tubular axis is highlighted by a color map. After Ref. [6].

Figure 6 shows the photo-emission spectra polymers in the vicinity of the Fermi level

of the peanut-shaped . It exhibits the power-law

dependences with respect to both the binding energy ( ) and temperature ( ). The TLL exponent, evaluated from the data within the energy range of

meV,

was

as found in Fig. 6(a). On the other hand, the power-law

dependence of

on

result of

in the temperature range of

. Hence, it is reasonable to conclude that the TLL

exponent Îą for the 1D uneven peanut-shaped significantly larger than

led us to the

polymer is

, which is

for metallic single-walled carbon nanotubes [36].

Emphasis should be put on that the curvature-induced shift in the TLL exponent, reported in Ref. [6], is the first experimental realization of the curvature-property relation in low-dimensional quantum systems. Furthermore, the remarkable consistency between theory and experiment tells us that the confining potential approach is effective to explore the curvature-property relations in and possibly other curved nanomaterials.

9

polymers


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Figure 6: Right: Photo-emission spectra of a peanut-shaped

polymer in the vicinity of the Fermi level.

The power-law dependences on the binding energy (left) and temperature (right) are observed. After Ref. [6].

1-4- Electron-phonon interaction in

polymers

In addition to the shift in the TLL exponent,

polymers exhibit distinct features

far from conventional nano-carbons. In the case of phononic excitations, infrared spectral measurements evidenced a rapid growth of specific eigenmode peaks with the increase in the electron-beam irradiation time [37]; this peak growth is attributed to the anomaly in the phonon density of states peculiar to quasi-1D systems [38]. A more important observation is a coupling-induced metal-semiconductor transition (i.e., the Peierls transition) in the system. Pump-probe spectroscopy measurements have revealed a spontaneous energy gap formation in the conduction band of the

polymers at a critical temperature

of several tens

[39]. This energy gap formation suggested the Peierls transition driven by electron-phonon couplings, implying that the conducting nature of the polymers vanishes below

. Hence, it is fundamentally important to unveil the

transition mechanism; however, the problem has been left to be considered until recently. A breakthrough has been made by the theoretical work [40] that was intended to artificial manipulation of

. It was suggested that carrier doping to the

polymers causes a sizable shift in the Fermi level from which 10

can be tuned artificially.

to

, by


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Figure 7: Transition temperature Without doping, we have

as a function of the Fermi level shift

because of carrier doping.

as indicated by a vertical line. Inset: Schematic of the commensurate

effect, illustrating the electron band structure and relevant electron-phonon couplings via the phonon wave number q. After Ref. [40].

Figure 7 shows the μ-dependence of

deduced theoretically in Ref. [40]. In the

analysis,

was defined in accord with the assumption that each carbon atom

within a

molecule gives a π-electron and each band accommodates two

electrons per a unit cell due to spin degeneracy. The plot indicates that decreases monotonically as μ increases; this decrease results from a reduction in the commensurate effect [41] that maximizes

at the half-filling state (i.e.,

meV in the present condition). The rapid decrease in

with the increase in μ is significant in light of controlling

the low-temperature electronic conductivity of the increment value of

eV causes a drastic reduction in

from 130 K to 1 K; this

coincides with the upward shift in the Fermi level that is observed

when one additional alkali atom per cavity of the

polymers. For example, an

molecule is inserted into the hollow

polymer. As a result, the metallic nature of the

polymers can

survive even at a few , unlike the case of the undoped system. This phenomenon provides a new avenue for controlling the low-temperature conductivity of polymers by doping, which is highly advantageous in actual device applications.

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Figure 8: Snapshots of the vortex dynamics observed in a superconducting nanoshell. The color scale indicates the Cooper pair density from blue (high) to red (low). The supercurrents are indicated by the arrow field. After Ref. [43].

2- Orientational order on curved surfaces 2-1- Curvature-induced frustration Another important consequence of surface curvature in low-dimensional materials is the complex interaction of geometry with orientational order and topological defects. Vortices in curved superfluid films [42], superconducting films [43], and magnetic thin films [17] are typical examples of such topological defects. When interacting constituents on a curved surface have orientational degrees of freedom, they can no longer show perfect orientational order. The loss of perfect orientational order is due to the non-commutative property of parallel transport of vectors [25]. On a curved surface, parallel transport of a vector along a closed loop does not maintain its direction but yields a rotation after the round trip (Fig. 9). This makes it impossible for all vectors to orient the same direction, yielding multiple frustrated states at low temperatures even if the system 12


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possesses no disorder. The imposition of curvature, therefore, causes disruption of the long-range order, leading to geometrically induced stresses [44]; to relax these stresses, defects and/or disclinations emerge spontaneously even in the ground state and often localize to regions of high Gaussian curvature [45].

Figure 9: Parallel translation of a vector along a closed loop on a curved surface. The arrow migrates from the position 1 to 9 (or 7) in turn with keeping its direction during the round trip.

2-2- Liquid crystal film on spherical surfaces Among many ordered systems, liquid crystal membranes with curved shape have been under vigorous investigations and regarded as the most fascinating for those in search of curvature-property relations in soft matters. Liquid crystals consist of aggregates of anisotropic molecules with orientational ordering [46], i.e., described through a unit vector n (called the director). The existence of a locally preferred direction and the resulting anisotropic optical properties make liquid crystals very suitable for electro-optic devices. When nematics are confined on curved surfaces, the molecular field is influenced by both geometrical and topological constraints. Theoretically, upon covering a microsize sphere with a thin layer of nematic liquid crystal, four half-strength topological defects are expected to emerge. This defect configuration is consistent with the PoincarĂŠ constraint that the total surface defect charge on a closed surface should be equal to just two [47]. Further interesting is the fact that the four 13


Advances in Science

defects, which reside on the vertices of a regular tetrahedron, represent high-free energy spots, potentially suitable for a chemical attack. Hence, the defectinvolving nematic sphere may enable to build complex colloidal architectures [48, 49] such as those for photonic applications. In this perspective, it is very important to control the coordination number and the valence angles of the colloidal atoms by changing the number of surface defects and/or their position on the spherical particle. The possibility of defect manipulation aforementioned, based on the electric field application to the nematic sphere, has been suggested in Ref. [51] using largescale Monte Carlo simulations. Figure 10 shows numerical results of director field and defect positions of nematic liquid crystals on a spherical shell. In the absence of an external electric field, a director configuration with four defect lines penetrating the shell as expected. Once turning on a homogeneous external electric field along the z axis, the gradual evolution of the defect line position on the sphere was observed as presented from Figs. 10(b)-(e). On increasing the field strength, first the defect line tetrahedron deforms and aligns as required by the field, and finally the four defect lines penetrating the shell partially coalesce to form two pairs of point defects at the inner or outer sphere poles [Figs. 10(d) and 10(e)]. Here the bipolar axis is parallel to the field direction. The field-driven structural transition suggested by Ref. [51], from a four halfstrength defect line structure to a bipolar structure with two pairs of defects on the poles, provides a promising way to design supramolecular atoms with controllable valence[47], and thus far different defect structures have been observed in colloids coated with thin nematic films [10].

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Figure 10: Director field and defect positions of nematic liquid crystals on a spherical shell. (a) The field-free conditions, and (b-e) The cases under the z-oriented homogeneous electric field with small (b,c) and large (d,e) magnitudes. The defects are indicated by red color, the director field by streamlines. The dark sphere represents the inner shell surface; the outer shell is omitted. After Ref. [51]

2-3- Liquid crystal film on a Gaussian bump Aside from closed surfaces, nematic membranes on an “open" surface with nonzero Gaussian curvature show non-trivial coupling between geometry and defects [52]. The coupling effect can be found in defect configuration of nematic films constrained on a Gaussian bump (see Fig. 11). The bump is one of the exemplary curved surfaces because it has a continuous rotation symmetry around a vertical axis, showing a positive (negative) Gaussian curvature close to (far from) the top of the bump. The coexistence of a positively curved region and a negatively one on the same surface allows a simultaneous observation of the curvature effects at the two distinct regions with oppositely-signed Gaussian curvature.

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Figure 11: A positive-charge defect of a liquid crystal film on a Gaussian bump.

In a continuum approximation, preferred configuration of the director field on a Gaussian bump is extracted by the minimization of the free energy given by âˆŤ

*

where h is the film thickness, and

+

and

(7)

are elastic constants of the director field,

indicates the covariant derivative [25]. Monte Carlo simulations have

revealed that point disclinations are attracted to the inflection point of the bump regardless of their charges. The result appears controversial with the prior work [44], which suggested that a point disclination with +1 charge is trapped to the top of the bump due to the curvature potential. The apparent controversy between the two studies was caused by the difference in the assumed bump widths. In Ref. [44], the bump width is so small that all the molecules over the whole bump consist of one single disclination; this situation is expressed by saying that the coherence length, , of the nematic order fluctuation exceeds the scale of the bump width. In contrast, the work performed in Ref. [52] was based on the assumption that the bump width is larger enough than

and the system can

involve many small disclinations. These arguments imply that the ratio of the 16


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bump width to

plays a decisive role in determining the stable position of

disclinations on a Gaussian bump. Interestingly, the two contrasting results imply the existence of a threshold ratio across which the stable position of disclinations alters. The conjecture was proved via the finding [53] of a sudden shift in the stable position of a +1 disclination with increasing the bump width. It was numerically confirmed that once the bump width exceeds a threshold value, the stable position is suddenly switched from the top to the inflection point of the bump. Further remarkably, the positional shift is described by a power law function of the bump width, as is analogous to the second-order phase transitions of various physical systems whose properties are characterized by power-law behaviors. The result thus indicates a novel kind of phase transition that is peculiar to curved nematic films.

3- Aquos Foam on Curved Surfaces 3-1- Coarsening dynamics on the flat plane The last topic is two-dimensional aqueous foam that also exhibits a good interplay between geometry and physics. With time, foam consisting of polyhedral bubbles evolves into the equilibrium structure, during which internal gas diffuses from a gas bubble to others through thin curved liquid interfaces [54, 55]. Diffusion is driven by the pressure difference between bubbles; assuming a constant diffusion coefficient, the pressure difference in two adjacent bubbles is proportional to the geometric curvature of their common boundary interface. Each boundary moves toward its concave sides due to the inter-bubble gas transfer, where the velocity of the boundary motion is again proportional to curvature [56]. As a result, some bubbles dilate while others shrink and eventually disappear, which results in a progressive increase in the average bubble size, i.e., the coarsening of foam.

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Figure 12: (Left) Foam assembly composed of many polygonal bubbles. (Right) Triangular-shaped bubble cell surrounded by three curved edges.

Foam that we encounter in our daily life, such as shaving cream and beer head, consists of a three-dimensional agglomerate. Its coarsening dynamics as well as equilibrium cellular structures have required formidable efforts for clarification because of geometrical and topological complexity. This is partly the reason why a large degree of attention has been paid to two-dimensional counterparts, i.e., a foam monolayer confined between two membranes. It is remarkable that in twodimensional foam confined between flat planes, the time evolution of a polygonal cell of bubble depends only on the number of its sides, regardless of its shape or area. The growth-rate of the area S of an n-sided cell is given by [57] (8) This formula states that a cell is stationary if constant rate if

but it grows (shrinks) at a

is larger (smaller) than 6. The disappearance of shrinking cells

causes a topological change in the network of liquid interfaces, whose effects on the stability of foam has also been largely investigated [58, 59, 60].

3-2- Coarsening dynamics on curves surfaces The coarsening behavior on the flat plane alters drastically when the foam is constrained to a curved surface [61, 62]. In the latter case, the evolution of cells is characterized by the Gaussian curvature

of the underlying surface. When the

surface has a positive (negative) curvature, n-sided cells with ( stationary, yielding 18

) can be

, only if S equals to a specific value that depends on


Advances in Science

and

(see Eq. (13) below). Furthermore, the stability of those stationary cells is

sensitive to the sign of

in addition to the value of . For instance, no cell on a

positively curved surface is stable; once a cell grows (shrinks) slightly under perturbation, then it keeps growing (shrinking). In contrast, all stationary cells on a negatively curved surface are stable; therefore, the equilibrium configuration consists of various n-sided cells each having a specific area determined by and . Such two-dimensional foam that spreads over a curved surface could be realized on an elastic confining plate or on a phase boundary with another fluid medium that repels the foam. The law of foam coarsening also applies to normal grain growth in metal, and an experimental set-up to observe surface curvature effects has been proposed [63]. The usefulness of statistical lattice models defined on curved surfaces [64, 65, 66, 67] to the foam coarsening dynamics [62] is also interesting from academic viewpoints. Let us assume that the monolayer foam is confined between two rigid membranes with spatially uniform Gaussian curvature . The gap between the membranes is smaller than the typical length of boundary curves. The growth-rate of the area of an n-sided cell with internal pressure p is described by ∑

Here,

(9) is the pressure difference between the cell and its jth neighbor,

ℓj is the length of the jth boundary curve separating the two cells, and γ>0 is a

diffusion constant. Equation (9) captures the simple idea that if the cell has a higher pressure than the jth neighbor (i.e., neighbor, and vice versa. With local equilibrium, tension

), then gas escapes to the is balanced by the line

along the interfaces, satisfying the generalized Laplace-Young law [61] (10)

where

is the geodesic curvature of the jth interface. From viewpoints of

differential geometry, any n-sided polygon on a surface with curvature

satisfies

Gauss-Bonnet’s theorem expressed by ∑ 19

(11)


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is the internal angle of the ith vertex and must be equal to 2Ď€/3 for all i

where

according to Plateau’s lemma [68]. Consequently, we obtain the result *

provided

+

(12)

is spatially uniform. Equation (12) is called a generalized von-

Neumann formula describing the foam coarsening on a rigid membrane [61].

3-3- Experimental test of generalized von-Neumann formula The first experiment to verify the theoretical predition, Eq. (12), has been recently performed [12] by realizing a curved two-dimensional foam confined into gap between a pair of hemispherical membranes. The set-up consists of two hemispherical polycarbonate domes; the smaller dome has an outer diameter of 12.5 cm, and the larger dome has an inner diameter of 13.3 cm, creating a 4-mm gap. Optical image processing was used to quantify individual bubble areas and their time evolution. Left panel in Fig. 13 displays the growth of particular three six-sided bubbles with different initial areas. There the initial area of each bubble was subtracted off so that the traces are easily comparable. The lines are a linear fit to the data, giving a constant growth rate

that is positive. The key feature in the plot is that the

six-sided bubbles grow, and that the larger ones grow faster. This agrees with the theoretical predition, Eq. (12), and contrasts strongly with the case of a planar (flat) system in which six-sided bubbles neither grow nor shrink as mentioned earlier. Right panel in Fig. 13 shows the growth rate for all bubbles; the vertical axis is the coarsening rate, and the horizontal axis is the expected proportionality corresponding to the hemisphere membranes (with constant curvature) used in the experiment. The inset is an enlarged view of all the six-sided bubble data, in which the three bubbles featured in the Left panel and highlighted. It is clearly shown that the growth rates for six-sided bubbles are all positive, furthermore, increases with bubble size, which is in a good agreement with the behavior deduced from Eq. (12). 20


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Figure 13: (Left) Time evolution of six-sided bubble areas A with initial area A0. Data extracted from three different bubbles are plotted. (Right) Growth rate of the bubble area versus the expected factor from the generalized von- Neumann formula. A is the bubble area and R is the radius of the dome. Inset: Blowup of the data for the n = 6-sided bubbles. Growth rates for the three representative bubbles in the left panel are circled. After Ref. [12].

3-4- Stability in the equilibrium foam structure Formula (12) accounts for the stability properties of n-sided cells on rigid curved surfaces. For

, such cells that satisfy | (

and

)|

(13)

can be stationary, although all stationary cells are unstable. For instance, if becomes slightly larger than

due to perturbation, then the quantity in the square

brackets in Eq. (12) becomes positive. Therefore, we obtain

after the

perturbation, which signifies a persistent growth in the cell. On the contrary, all stationary cells are stable for

, since S being larger (smaller) than

makes

negative (positive). In this context, the case of a flat plane is marginal, in which the stationary cell of n=6 is neither stable nor unstable against perturbation. Furthermore, the stability of those stationary cells is sensitive to the sign of

in

addition to the value of n. For instance, no cell on a positively curved surface is stable; once a cell grows (shrinks) slightly under perturbation, then it keeps growing (shrinking). In contrast, all stationary cells on a negatively curved surface are stable; therefore, the equilibrium configuration consists of various n-sided cells each having a specific area determined by n(>6) and

. Such two-

dimensional foam that spreads over a curved surface could be realized on an

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elastic confining plate or on a phase boundary with another fluid medium that repels the foam. It is also noteworthy that if the two confining membranes are mechanically deformable, then the time-varying deformation of the membranes gives rise to a significant alteration in the formula (12); the alteration results from the correlation between the rate of inter-cell gas transfer and temporal fluctuation in surface curvature within a cell domain [69]. Membrane deformation may cause a flow of liquid through the films, as analogous to foam drainage in response to gravity and capillarity [70, 71]. In fact, the deformation of the confining membranes leads to rearrangement of the liquid film network as well as cell configuration, and thus, inducing pressure gradient in the liquid. Theoretically, it was shown that fluid flow on a surface with time-varying surface curvature exhibits three kinds of dynamical responses depending on geometric and material constants [72]. Hence, deformation-induced fluid flow in the current system may behave differently from the case of a rigid membrane, which gives rise to three-cornered coupling of surface deformation with gas transfer and associated fluid flow.

Summary and Perspective We have had an overview of the progress achieved in the field of “curvatureproperty relations" in condensed matters. Three individual topics were discussed: quantum mechanics, liquid crystals, and foam dynamics. We saw that the Hamiltonian of a quantum particle confined to a curved surface involves a curvature-induced effective potential, which strongly affect the elementary excitations of particles and their correlated behaviors. Liquid crystal membranes with curved shape were found to show anomalous couplings between the surface curvature and topological defect configurations; the energetically preferred configuration is determined by the spatial distribution of the curvature. Lastly, we have learned that aqueous foam sandwiched by two parallel transparent sheets shows coarsening dynamics, in which the surface curvature sways the fate of polygonal gas bubbles. 22


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There are a variety of other topics relevant to curvature-property relations in physics, although they were omitted due to the page number limitation. In particular, quantum mechanics on curved geometry come to a flurry of theoretical activity: only a few to mention are surface plasmon polaritons on curved surfaces [73], electron transport [74] and light propagation [75] through helical wires, and relativistic effects on spin-1/2 particle under electromagnetic field [76]. It is hoped that the discussion of the curvature-property relations presented in this chapter may stimulate theoretical physicists to seek other unveiled curvature effects on condensed matter physics, which will trigger experimental efforts to pave the way for the design of measurements to detect the theoretical predictions.

Acknowledgement This work was supported by JSPS KAKENHI (Grant Number 22760058). Financial support from the Asahi Grass Foundation is greatly acknowledged.

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[40] S. Ono and H. Shima (2011) Flexible control of the Peierls transition in metallic C60 polymers, EPL (Europhys. Lett.) 96, 27011. [41] G. Grüner (1994) Density Waves in Solids, (Addison-Wesley, Reading, MA) [42] H. Kuratsuji (2012) Stochastic theory of quantum vortex on a sphere, Phys. Rev. E 85, 031150. [43] J. Tempere, V. N. Gladilin, I. F. Silvera, J. T. Devreese, and V. V. Moshchalkov (2009) Coexistence of the Meissner and vortex states on a nanoscale superconducting spherical shell, Phys. Rev. B 79, 134516. [44] C. D. Santangelo, V. Vitelli, R. D. Kamien and D. R. Nelson (2007) Geometric Theory of Columnar Phases on Curved Substrates, Phys. Rev. Lett. 99, 017801. [45] V. Vitelli and A. M. Turner (2004) Anomalous Coupling Between Topological Defects and Curvature, Phys. Rev. Lett. 93, 215301. [46] P. G. de Gennes and J. Prost (1995) The Physics of Liquid Crystals, (Oxford University, New York). [47] D. R. Nelson (2002) Toward a Tetravalent Chemistry of Colloids, Nano Lett. 2, pp.1125-1129. [48] S. C. Glotzer and M. J. Solomon (2007) Anisotropy of building blocks and their assembly into complex structures, Nat. Mater. 6, pp.557-562. [49] I. Muševič, M. Škarabot, U. Tkalec, M. Ravnik and S. Ţumer (2006) TwoDimensional Nematic Colloidal Crystals Self-Assembled by Topological Defects, Science 313, pp.954-958. [50] H. Shin, M. J. Bowick and X. Xing (2008) Topological Defects in Spherical Nematics, Phys. Rev. Lett. 101, 037802. [51] G. Skačej and C. Zannoni (2008) Controlling Surface Defect Valence in Colloids, Phys. Rev. Lett. 100, 197802. [52] I. Hasegasa and H. Shima (2010) Point-Defect Haloing in Curved Nematic Films, J. Phys. Soc. Jpn. 79, 074607.

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[53] I. Hasegasa and H. Shima (2011) Continuous transition of defect configuration in a deformed liquid crystal film, Mod. Phys. Lett. B 25, pp.581588. [54] D. Weaire and N. Rivier (1984) Soap, cells and statistics: random patterns in two dimensions, Contemp. Phys. 25, pp.59-99. [55] M. F. Vaz (2008) Liquid foams: an introduction, Phil. Mag. Lett. 88, 627636. [56] D. J. Durian, D. A. Weitz and D. J. Pine (1991) Scaling behavior in shaving cream, Phys. Rev. A 44, R7902-R7905. [57] J. von Neumann (1952), in Metal Interfaces, p. 108., edited by C. Herring (American Society for Metals, Cleveland,) [58] M. Durand and H. A. Stone (2006) Relaxation Time of the Topological T1 Process in a Two-Dimensional Foam, Phys. Rev. Lett. 97, 226101. [59] S. J. Cox, G. Graner, and M. F. Vaz (2008) Screening in dry two-dimensional foams, Soft Matter 4, pp.1871-1878. [60] J. -P. Raven and P. Marmottant (2009) Microfluidic Crystals: Dynamic Interplay between Rearrangement Waves and Flow, Phys. Rev. Lett, 102, 084501. [61] J. E. Avron and D. Levine (1992) Geometry and Foams: 2D Dynamics and 3D Statics, Phys. Rev. Lett. 69, pp.208-211. [62] P. Peczak, G. S. Grest and D. Levine (1993) Monte Carlo studies of grain growth on curved surfaces, Phys. Rev. E 48, pp.4470-4482. [63] D. Levine, J. E. Avron and A. Brokman (1992) Grain Growth on Curved Surfaces, Mat. Sci. Forum 94-96, pp.281-284. [64] H. Shima and Y. Sakaniwa (2006) Geometric effects on critical behaviours of the Ising model, J. Phys. A: Math. Gen. 39, pp.4921-4933. [65] H. Shima and Y. Sakaniwa (2006) The dynamic exponent of the Ising model on negatively curved surfaces, J. Stat. Mech.: Theor. Exp., P08017. [66] Y. Sakaniwa and H. Shima (2009) Survival of short-range order in the Ising model on negatively curved surfaces, Phys. Rev. E 80, 021103.

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[67] S. K. Baek, H. Shima and B. J. Kim (2009) Curvature-induced frustration in the XY model on hyperbolic surfaces, Phys. Rev. E 79, 060106(R). [68] J. Plateau (1873) Statique Expŕimentale et Thórique des Liquides Soumis aux Seules Forces Molćulaires, (Gauthier-Villars, Paris). [69] H. Shima (2010) Growth of aqueous foam on flexible membranes, J. Phys. Soc. Jpn. 79, 074601. [70] A. Saint-Jalmes (2006) Physical chemistry in foam drainage and coarsening, Soft Matter 2, pp.836-849. [71] K. Feitosa and D. J. Durian (2008) Gas and liquid transport in steady-state aqueous foam, Eur. Phys. J. E 26, pp.309-316. [72] M. Arroyo and A. DeSimone (2009) Relaxation dynamics of fluid membranes, Phys. Rev. E 79, 031915. [73] G Della Valle and S Longhi (2010) Geometric potential for plasmon polaritons on curved surfaces, J. Phys. B: At. Mol. Opt. Phys. 43, 051002. [74] G. Cuoghi, A. Bertoni and A. Sacchetti (2011) Effect of quasibound states on coherent electron transport in twisted nanowires, Phys. Rev. B 83, 245439. [75] H. Taira (2011) Spin transfer of light waves in twisted optical waveguides, J. Phys. B: At. Mol. Opt. Phys. 44, 195401. [76] T. Kosugi (2011) Pauli Equation on a Curved Surface and Rashba Splitting on a Corrugated Surface, J. Phys. Soc. Jpn. 80, 073602.

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Chapter

2 Graphene-metal interaction investigated by soft X-ray spectroscopies Liang Zhang and Junfa Zhu* National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230029, China E-mail :

T

hshima@yamanashi.ac.jp

he synthesis of graphene on metal surfaces (such as Ni, Ru, Pt, Cu, Ir and Fe) by chemical vapor deposition (CVD) is one of the most promising methods to prepare single-layer and large-area graphene, which is a prerequisite for the fabrication of graphene-based electronic devices. Therefore, understanding the fundamental properties of graphene/metal interfaces is of great significance because of their important roles in graphene synthesis by CVD method as well as developing graphene-based nanoelectronics. Recently, great efforts have been devoted to understanding the graphene-metal interfacial interaction both experimentally and theoretically. In this chapter, we will review the latest progress in the investigations of graphene/metal interfacial properties by the means of X-ray absorption spectroscopy (XAS), X-ray emission spectroscopy (XES) and resonant inelastic X-ray scattering (RIXS). It has been found that the RIXS spectra of graphene on different metal surfaces change dramatically, which can be directly related to the different strengths of hybridization between graphene and metal substrates. These significant spectra changes have been proved to be an effective measure for the bonding strength of graphene/metal interface: strong band dispersion can be observed when the interaction between graphene and metal substrate is weak, while the band dispersion is seriously disturbed when a strong hybridization between graphene and metal surface exists. In addition, potential applications of XAS, XES and RIXS in the electronic structure characterization of graphene-based functional nano-materials are also reviewed on the basis of recent experimental studies.

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Introduction Graphene is a new material-star with a one-atom-thick planar sheet of sp2-bonded carbon atoms packed in a hexagonal lattice. Since its first discovery by “scotch tape� in 2004, graphene has attracted extensive attention due to its unique electronic structure and extraordinary physical properties [1]. For example, graphene has a high thermal conductivity (~5000 W/mK) [2], high carrier mobility (~250, 000 cm2 V-1 S-1) [3], strong Young’s modulus (~1 TPa) [4] and large surface area (~2600 m2/g). Due to these outstanding properties, graphene is expected to have wide applications in nano-electronics [5], sensing [6], catalysis [7], energy storage [8-10], as well as bio-electronic device [11] in the future. On the other hand, graphene is the mother element of other carbon allotropes: it can be wrapped into zero-dimensional fullerenes, rolled into one-dimensional nanotubes, and stacked into three-dimensional graphite (Figure 1) [5]. As a consequence, graphene-related study is of great significance for a basic understanding of the relationship between the electronic structure and performance of the nanostructured carbon materials. There are mainly four methods to prepare graphene layers: micromechanical cleavage of graphite crystal [1], epitaxial growth on a SiC substrate at high temperature [12], chemical vapor deposition (CVD) growth on metal surfaces (such as Ni, Ru, Cu, Co and Pt) [13-16] and chemical synthesis from graphite [17, 18]. Among them, CVD is the most widely used method to prepare defect-free, large-area and single-layer graphene [14].

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Figure 1. Basis of all graphitic forms.

Graphene is a 2D building unit for other carbon allotropes [3]. Graphene formed on metal substrates by CVD method has been known for more than 40 years and it was first observed on Pt and Ru single crystals [19, 20]. It has been found that only monolayer or few-layer graphene can be formed on metal surfaces [21]. In addition, the graphene moirĂŠ patterns arising from the lattice mismatch or rotation of the graphene lattice with respect to that of metal substrates have also been utilized as the templates for self-assembly of metal clusters or organic molecules [22, 23]. The current boom in the research of free-standing graphene has led to the renewed interests in graphene/metal systems. It has been demonstrated that the graphene-metal interaction is different for graphene on different metal surfaces, and, moreover, charge transfer or band gap can be induced by direct contact of graphene with metal surfaces [24-26]. Therefore, the knowledge of the interfacial interactions and electronic properties of graphene/metal systems is useful for many aspects, ranging from interest in understanding the basic properties of material to applied areas of research.

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Figure 2. Schematic illustration of C K-edge XAS and XES process.

The challenge in probing the electronic band structure of graphene on metal substrates is to separate the contributions of graphene from those of the underlying substrates. X-ray spectroscopies, such as X-ray absorption spectroscopy (XAS), X-ray emission spectroscopy (XES) and resonant inelastic X-ray scattering (RIXS), provide an atom-specific probe of the electronic structure which allows for such separation [27-31]. For the C K-edge XAS process, the C 1s electrons are excited to the empty electronic states in the conduction band, and the dipole selection rule provides a powerful measurement to study the local C 2p characters of these unoccupied states (Figure 2). In principle, the resonance in electronic transitions from core states to unoccupied Ď€* or Ďƒ* states is strongly enhanced given the electric filed vector of the incident synchrotron light is parallel to the

or Ďƒ* orbital, which makes XAS a sensitive

method to probe the alignment and orientation of thin films [32-38]. In the C K-edge XES process, the core holes left by the X-ray excited C 1s electrons are filled by valence electrons (Figure 2), which can give the direct chemical bonding information of the valence band [27, 28]. In addition to probing the local valence density of states (DOS), resonant XES, namely RIXS, allows the crystal-momentum-resolved (k-resolved) electronic structure to be studied due to the conservation of crystal momentum during the scattering process. The spectral 34


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properties of the RIXS process can be described by the following KramersHeisenberg equation [25, 26]:

[

]

Here, ⟨

∑ ∑ |

| and |

|

| ⟩⟨ |

|

|

(1)

⟩ are the conduction and valence band states at k point with

energies Eck and Evk, respectively; | ⟩ is the core hole state with energy Es; while pa and pβ denote the momentum operators for the incident and outgoing beam polarization, with energies 1 and 2, respectively. The combination of XAS, XES and RIXS has been frequently utilized to characterize carbon-based materials, providing deep insight into the electronic strucutres of both occupied and unoccupied DOS [28, 39-41]. In this chapter, we will give a brief overview of the structural and electronic properties of graphene/metal systems investigated by XAS, XES and RIXS. For comparison, the graphite and graphene/SiO2 systems are also taken into consideration. In addition, applications of XAS, XES and RIXS in the electronic structure characterization of graphene-based functional nano-materials are also reviewed on the basis of recent experimental studies.

1- Soft X-ray spectroscopy studies on graphite Graphite, with a layered structure and large interlayer space, has been the subject of a great deal of experimental and theoretical efforts [42-46]. Figure 3 shows the C K-edge XAS spectra of graphite as a function of incidence angle . The intensity of peak A increases significantly as can be ascribed to C 1s to peak A when 35

transition

is increased. Therefore, this feature

unequivocally. The residual intensity of the

is 0° may be caused by the sample misalignment or incomplete


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polarization of the synchrotron light [42]. For electric dipole transitions excited by the polarized light, the intensity should vary as a function of graphite, the intensity of C 1s to while the intensity of C 1s to

. Therefore, for

transition should be proportional to sin2Îą, transition should be proportional to

. As

shown in the Figure 4, the intensity of peak A increases linearly as a function of corroborating the assignment of the first peak.

Figure 3. Angle-dependent XAS spectra of graphite [42].

Figure 4. Relative intensity of peak A in Figure 3 as a function of

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Figure 5. Resonant and non-resonant XES spectra of HOPG at different incident energies. The upper panel is the graphite band structure with the energy axis matched to the photon energy axis of the XES data [45].

Figure 5 shows the resonant and non-resonant XES spectra of highly oriented pyrolytic graphite (HOPG) at different excitation energies [45]. The uppermost spectrum is the non-resonant case where the excitation energy is far above the absorption threshold (hνin = 400.0 eV). It shows a wide feature centered at ~276.0 eV and a high energy peak centered at ~281.0 eV, which result from the

and

states, respectively [43, 45]. The resonant spectra are shown below the nonresonant spectrum. The resonant spectral shape changes significantly at different excitation energies and most of the spectra have a drastically different spectral shape compared to the non-resonant one. The dispersive peaks in the resonant spectra are highlighted by the dashed lines, as numbered 1−7 in the figure. The emission spectrum changes most obviously when increasing the photon energy from 285.0 eV to 285.5 eV. Such a large change in the RIXS spectrum cannot solely be attributed to simple intensity modulation, which was explained in terms of sweeping through critical points in the Brillouin zone [47]. Clearly, there is a correlation between the dispersive emission features and band structure of HOPG 37


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due to the conservation of crystal momentum during the resonant inelastic X-ray scattering process. In other words, by tuning the incident photon energy to excite the core electrons into unoccupied states with a well-defined momentum, we are able to get the band structure of HOPG. Because the band structures of graphite and graphene are very similar (Figure 6) and the only difference is just a marginal splitting of the bands in graphite due to the interlayer coupling. Therefore, similar band dispersions should also be observed in the RIXS spectra of graphene.

Figure 6. Calculated band structure of graphite (a) and graphene (b). The marked areas indicate the band structure difference between graphite and graphene.

2- Soft X-ray spectroscopy studies on graphene/metal and graphene/SiO2 systems 2-1- graphene/Cu Another Fabrication of graphene on Cu foils by CVD is the most promising method to prepare large-area and single-layer graphene [14, 48] due in part to the low solubility of C in Cu [49]. It is essential to understand the interfacial interaction between graphene and Cu surface to better control the growth parameters of CVD process and improve the quality of graphene films [38].

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Figure 7. (a) Schematic illustration of the XAS and XES experiments. (b) C K-edge XAS spectra of graphene/Cu, graphene/SiO2 and HOPG. (c) C K-edge XAS spectra of graphene/Cu collected at different incident angle .

Figure 7 displays the C K-edge XAS spectra of graphene/Cu [38]. The data were obtained by rotating the sample in the theta direction of the manipulator (the experimental geometry is shown in Figure 7a). Figure 7b demonstrates a comparison of the XAS spectra for HOPG, graphene/Cu and graphene/SiO2. For HOPG, three resolvable features can be observed, which are ascribed to the state (peak A), excitonic state (peak E) and

state (peak F), respectively [35, 42,

43]. In contrast, three more features (features B, C and D) can be observed in the XAS spectrum of graphene/Cu, indicating that these new electronic states only exist in the conduction band of graphene/Cu. The feature B is caused by the edge states of graphene with C-H bonds due to the strong reducing conditions of CVD [33]. For feature C, it can be mainly assigned to the interlayer interaction between graphene and Cu [33]. Moreover, the absence of feature C in the XAS spectrum of graphene/SiO2 confirms our assignment considering the negligible interaction between graphene and SiO2 [50-53]. The peak D is attributed to the defect states in graphene formed during the CVD process [54]. Figure 7c shows the angledependent XAS spectra of graphene/Cu [38]. Similar as the results of graphite, the resonance of graphene shows strong dichroism, indicating a strong anisotropic 39


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effect for graphene which results from the alignment ordering between graphene and Cu. However, there is still remnant intensity of

state at normal incidence,

indicating the warping or crumpling of graphene on Cu. Figure 8 demonstrates the resonant and non-resonant XES spectra of graphene/Cu along with the corresponding spectra of HOPG [38]. The XES data were acquired at an incident angle of 60째 and emission angle of 30째. The non-resonant XES spectra were collected with an incident energy of 320.0 eV. The non-resonant XES spectra of graphene/Cu and HOPG are very similar except for some intensity modulation for the states at higher emission energies, which may be caused by the charge transfer from Cu to graphene or the difference in the band structures of graphene and HOPG.

Figure 8. RIXS spectra of graphene/Cu and HOPG recorded at different excitation energies [38].

The resonant XES spectra were collected at different excitation energies corresponding to the absorption features in XAS spectrum. The spectral shape changes significantly with the variation of excitation energies, especially when it is close to the absorption threshold, indicating that the corresponding transitions occur from states with a well-defined crystal momentum [27, 28, 38, 45, 55]. Because the band structures of HOPG and graphene are very similar, the resonant 40


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XES spectra of graphene/Cu should be treated with the same RIXS formalism as that of HOPG [56]. From Figure 8 we can see that the dispersive features in the resonant XES spectra of graphene/Cu due to the conservation of crystal momentum are clearly observed, indicating that the graphene on Cu preserves the intrinsic momentum as that of HOPG due to the weak interaction between graphene and Cu [57-59]. However, some obvious differences can still be found when comparing the RIXS spectra of graphene/Cu with the corresponding spectra of HOPG. Firstly, the intensities of the elastic peaks in the resonant XES spectra of graphene/Cu are always stronger than those of HOPG at different excitation energies. Generally speaking, the intensity of elastic peak is strongly dependent on the extent of the localization of the excited photoelectrons at the core-level site [41, 60]. The stronger elastic peak for graphene indicates that graphene-derived emission states are more excitonic in nature. Secondly, broad contributions at higher emission energies can always be observed for graphene/Cu, which may be caused by the stronger electron-phonon scattering in graphene. Finally, subtle shifts of the main emission features for graphene/Cu to higher emission energies are found compared with those of HOPG, which can be a result of the different DOS for HOPG and graphene/Cu. In addition, the chemical shift due to the charge transfer between graphene and Cu may also induce the corresponding peak shift.

2-2- graphene/Pt Graphene/Pt is another excellent system for probing the effect of weak graphenemetal substrate interaction [16, 37, 62-64]. The C 1s XPS spectrum of graphene/Pt shown in Figure 9a exhibits a single sp2 peak, indicating that the graphene samples are very clean without any impurities. The low-energy electron diffraction (LEED) pattern displayed in Figure 9b confirms the periodicity of the graphene overlayer with respect to the Pt(111) surface [37, 62, 65, 66]. Figure 9c 41


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shows the well-developed surface and bulk components of Pt 4f7/2 for pure Pt and graphene/Pt, which are located at 70.55 and 70.92 eV, respectively. After the presence of graphene layer on Pt, the intensity of the surface component is found to attenuate only slightly, indicating the weak interaction between graphene and Pt. Moreover, an additional peak located at 71.40 eV appears for graphene/Pt, which can be assigned to the increased graphene-substrate interaction at specific Pt-C sites.

Figure 9. (a) C 1s XPS of graphene/Pt fitted with one sp 2 component. (b) LEED pattern of graphene/Pt. (c) XPS spectra of Pt 4f7/2 for pure Pt (red dash line) and graphene/Pt (black line). From low to high binding energy, the three deconvoluted components indicate the surface, the bulk state of Pt, and the graphene-Pt state [61].

Figure 10 shows the σ- and π-resolved C K-edge XAS spectra of graphene/Pt(111). In the π-resolved XAS spectrum, the

feature at 285.5 eV

exhibits a weak shoulder at 284.4 eV. The pre-edge feature could arise from weak orbital mixing between Pt(111) and graphene. This assignment is supported by the evidence that the pre-edge features are more pronounced for graphene on strongly interacting transition metals [67]. The σ-resolved XAS spectrum displays an excitonic 42

state at 291.0 eV. In addition, the spectrum also shows a weak,


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symmetry-forbidden feature related to the

resonance. This weak feature can be

caused by the small rippling of the graphene overlayer which may lead to minor tilting of the individual carbon atoms.

Figure 10. Out of plane (black line) and in-plane polarized (red line) XAS spectra of graphene/Pt. The arrows indicate the excitation energies where RIXS spectra were recorded [61].

The influence of substrate effect on the valence-band structure of graphene was monitored by RIXS. The C K-edge RIXS spectra in grazing- and normal-emission angle for graphene/Pt are shown in Figure 11. The spectra are displayed with respect to the Fermi level referred to C 1s binding energy at 284.0 eV. The nonresonant XES with the excitation energy far above the absorption threshold (320.0 eV in the current case) can be regarded as a measure of the valence DOS of graphene. The RIXS process for graphene can be interpreted within a one-electron picture, although the core-hole effects could alter the relative intensity contributions from coherent and incoherent processes [47]. In a one-electron RIXS picture of graphene, due to the conservation of crystal momentum during the scattering process, excitation energies of 284.5, 286.0 and 292.0 eV should give rise to resonant emission from states close to K, M and Γ points, respectively, which is in good accordance with the experimental results (Figure 11). The corresponding emission features are therefore assigned to specific high-symmetry points in the k space, as demonstrated in Figure 11. 43


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By comparing the RIXS spectra between graphene/Pt and graphite, one could get valuable information on the influence of substrate on the band structure of graphene. Figure 12 shows such a comparison at incident photon energies of 284.5 and 320.0 eV, respectively. The spectra excited at 320.0 eV for graphene/Pt and HOPG are very similar, indicating the similar valence DOS between them. However, the carbon-projected graphene DOS cuts the Fermi level directly, whereas graphite has a low DOS at this energy range due to its semimetallic character [68]. In contrast, significant changes in the RIXS spectra of graphene/Pt at excitation energy of 284.5 eV can be observed when compared with the corresponding spectrum of graphite: (1) strong emission features near E F where emission could be expected to occur, and (2) additional emission intensities from 6 to 9 eV and 2.5 to 6 eV for the σ and π states, respectively. The first observation could arise from the substrate hybridization effects. Moreover, because Pt has a large DOS near the Fermi level, the weak dispersion of the emission feature is consistent with the contribution of Pt to the associated graphene valence states. Two possible origins could be responsible for the second observation. One possibility is the break of the momentum coherence induced by the weakly dispersing C-Pt state at the Fermi edges. Another possibility is the surface umklapp scattering arising from the superperiodic potential induced by Moiré formation.

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Figure 11. (a) Calculated band structure of graphene. (b) RIXS spectra of graphene in grazing (black lines) and normal (red lines) emission geometry at different excitation energies [61].

Figure 12. RIXS spectra of graphene and HOPG at excitation energies of 284.5 eV and 320.0 eV. The dotted rectangle indicates the emission region due to surface umklapp process [61].

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2-3- graphene/Ni Different from graphene/Cu and graphene/Pt, the interaction between graphene and Ni for graphene/Ni is very strong due to the different configurations of metal 3d states among Ni, Pt and Cu [25, 69, 70]. Therefore, the influence of Ni on the electronic properties of graphene is expected to be different from those of Cu and Pt. Figure 13a displays the scanning tunnel microscopy (STM) image of graphene/Ni(111) surface acquired at room temperature. All investigated terraces show the same atomic structure with well-resolved honeycomb lattice of graphene (Figure 13b), indicating the high quality of graphene films. The LEED pattern in the insert of Figure 13a reveals a well-ordered 1×1 structure as expected due to the small lattice mismatch between graphene and Ni [71]. Figure 13c shows the angle-dependent photoemission spectra of graphene/Ni taken at excitation energy of 40.8 eV along the Γ-M direction. Compared with the photoemission spectra of graphite, the binding energy of the π states for graphene/Ni is 2.3 eV higher than that of graphite, in good agreement with the theoretical prediction of 2.35 eV [72]. This shift indicates the existence of hybridization of graphene π bands with the Ni 3d bands as we have stated above.

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Figure 13. (a) STM image of graphene/Ni. The insert shows a LEED image obtained at 63 eV. (b) Atomic structure of graphene monolayer. (c) Angle-dependent photoemission spectra of graphene/Ni measured with photon energy of 40.8 eV [73].

Figure 14a displays the C K-edge XAS spectra of graphite and graphene/Ni. Two pronounced regions of 283-289 eV and 289-315 eV can be ascribed to C 1s to and C 1s to

transitions, respectively. Comparing the spectra of graphite and

graphene/Ni, it can be clearly seen that the 1s to

and

transitions are changed

significantly, indicating the strong chemisorption for the latter case. Same phenomena were also observed for graphene/Rh and graphene/Ru, indicating the strong interaction between graphene and these metal surfaces [67]. The broadened

and

resonances can also be caused by the strong orbital

hybridization and electron sharing at the graphene/Ni interface. Figure 14b shows the C K-edge XAS spectra of graphene/Ni as a function of incident angle Îą. Similar to the angle-dependent XAS spectra of graphite [42] and graphene/Cu [38], the spectra of graphene/Ni also show strong dichroism, indicating the high alignment of graphene films on Ni substrate. Compared with the calculated C K47


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edge electron loss spectra of graphene/Ni [72], the first peak at 285.5 eV can be attributed to the transition of the electrons from 1s core level to the interface state above the Fermi level (around K-point in the BZ) originating from the C pz-Ni 3d hybridization states. These states correspond to the antibonding between carbon atoms C-top and interface Ni atoms. The second peak at 287.1 eV can be assigned to the transition of the 1s core electrons to the interface states above the Fermi level (around M-point in the BZ), which originate from the C Pz-Ni px, py, 3d hybridization. In contrast, these states correspond to a bonding between the two carbon atoms, C-top and C-fcc, which involves the nickel interface atoms.

Figure 14. (a) C K-edge XAS spectra of graphite and graphene/Ni. (b) Angle-dependent XAS spectra of graphene/Ni [74].

Figure 15 shows the X-ray magnetic circular dichroism (XMCD) spectra of the graphene/Ni at (a) Ni L2,3-edges and (b) C K-edges in total-electron-yield (TEY) and partial-electron-yield (PEY) modes, respectively. The bulk values of spin- and orbital-magnetic moments

= 0.69 and

= 0.07 for Ni based on the sum-rules

from the Ni L2,3 XMCD spectrum are in good accordance with the results published before [75]. However, a relative large XMCD contrast at the C K-edge can also be observed. The results indicate that the major magnetic response stems from the transitions of the 1s electrons to the 48

states. In contrast, no such


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magnetic response can be observed in the transitions of the 1s electrons to the Ďƒ* states, indicating that only C 2pz orbitals are polarized which hybridize with the Ni 3d band. The main reason is that the feature at the 1s to

absorption edge is

originated from hybridized C pz-Ni 3d and C pz-Ni px, py, 3d states. Because the XMCD at the K edge provides only the information on the orbital moment, it can be concluded that the orbital moment of the carbon atoms in the graphene layer is parallel to the spin and orbital moments of the Ni substrates. This observation is similar to that for Fe/C multilayers where the magnetism is related to the hybridization of the Fe 3d orbitals and the C pz orbitals [76].

Figure 15. XMCD spectra of the graphene/Ni system: (a) Ni L2,3 and (b) C K-edge [74].

Figure 16 displays the RIXS spectra of graphene/Ni excited at different excitation energies. The spectrum excited at 320.0 eV corresponds to the non-resonant one. This spectrum is similar to that of graphite, indicating the analogous electronic configurations between them. In addition, the resonant emission spectra also show considerable dispersive features at different excitation energies, indicative of the conservation of crystal momentum for graphene/Ni to some extent. However, the RIXS spectra of graphene/Ni are quite different from the corresponding spectra of graphene/Cu (Figure 8) when the excitation energies are close to the absorption threshold, suggesting that the electronic structures of graphene have changed 49


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because of the stronger interfacial interaction between graphene and Ni surfaces. Since the electronic structures of the 4s and 4p states of Cu and Ni are similar, while their 3d states are quite different, we mainly ascribe the variations in the spectral shape to the different bonding effect of 3d states: strong for graphene/Ni and weak for graphene/Cu. For Cu, it is difficult to hybridize with the p band of graphene because the d orbital of Cu is fully occupied, resulting in a weak bonding between graphene and Cu. Therefore, the emission spectra of graphene/Cu and HOPG are similar (Figure 8). DFT calculations also indicate that the bonding between graphene and Cu is so weak that the unique electronic structure of graphene is preserved [25]. While for Ni, the d orbital is only partially occupied. As a consequence, strong covalent bonding arising from the hybridization between d orbitals of Ni and pz orbitals of graphene can be formed, which results in considerable charge redistribution and symmetry breaking of graphene. This observation is similar to the result for h-BN on Ni and Cu: h-BN was strongly chemisorbed on Ni, while on Cu only weakly chemisorbed h-BN monolayer was observed [77].

Figure 16. RIXS spectra of graphene/Ni excited at different excitation energies.

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2-4- Comparison of graphene on different metal surfaces To investigate the influence of metal substrate on the electronic properties of graphene, Preobrajenki et al. studied the electronic structure of graphene on different transitional metal surfaces by XAS and photoelectron spectra (PES) [67]. Figure 17 shows the comparison of C K-edge XAS spectra of graphene on Pt, Ir, Rh and Ru, respectively, along with their corresponding LEED patterns. As has been mentioned above, the peaks A, B, and C in the spectrum of HOPG result from the excitation of core electrons to

state, excitonic state and

state. The

overall spectral shapes between HOPG and graphene/Pt are very similar, indicating the weak orbital hybridization between graphene and Pt [62]. This results in a very weak chemical bonding of graphene to Pt, as can be seen from the incommensurate structure in the LEED pattern (Figure 17b). However, a small shoulder A' at about 284 eV appears for graphene/Pt, which can originate from the weak orbital mixing between graphene and Pt.

Figure 17. (a) C K-edge XAS spectra of graphene on different metal surfaces and HOPG. (b)-(e): Corresponding LEED patterns for graphene/metal [67].

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More visible changes can be observed in the C XAS spectrum when going from graphene/Pt to graphene/Ir (Figure 17a). A new feature named A" appears in the energy range between the

and

resonances, which indicates the increase of

orbital mixing compared with graphene/Pt. The increase is also confirmed by the coincidence lattice structure with just one single domain observed in LEED (Figure 17c). This can be mainly ascribed to the reduction in the 5d occupancy in going from Pt (d9) to Ir (d7), which leads to the growing covalent bonding. However, in general, the bonding between graphene and 5d metal surfaces is weak. In contrast, the C 1s XAS spectrum of adsorbed graphene changes significantly when going from 5d (Ir, Pt) to 4d (Rh, Ru) metals. In addition to the increase of the intensities of chemisorption-induced separation between

and

bands A' and A", the energy

bands is also decreased. The reduction of the π-σ

energy separation is a signature of the π bond softening for graphene on Rh and Ru substrates due to the electron sharing with the substrates. The broadening of the excitonic state is another evidence for a strong orbital hybridization and electron sharing at the graphene/Rh and graphene/Ru interfaces due to the strong delocalization of the corresponding core-excited states. Although the covalent bonds at the graphene/Rh and graphene/Ru are much stronger than that of graphene/Ir, all of them show similar moiré patterns in LEED due to similar lattice mismatch. However, the C 1s PES can help to reveal the variation of graphene morphology. Figure 18 shows the C 1s PES spectra of HOPG and all four interfaces under study and their respective peak fittings. For HOPG, the C 1s spectrum contains a single peak with the Lorentzian full width at half maximum (LFWHM) of 150 meV. For graphene/Pt and graphene/Ir, both the C 1s spectra have a single feature with the LFWHMs of 130 and 150 meV, respectively. The C 1s lines split into two distinct components with energy separations of 0.53 and 0.60 eV for graphene 52


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on Rh and Ru, respectively. The double-peak shape can be associated with the corrugation of graphene, as illustrated in Figure 18. The feature with lower binding energy (C1) is due to the nonbonding part of the graphene layer, while the feature with higher binding energy (C2) is ascribed to the strongly bonded parts. It is known that the substrate-adsorbate distance for graphene on metal surfaces is strongly dependent on the chemical bonding [21]. For example, the distance between graphene and the topmost Ni layer is only 2.1 Ă… due to the strong chemical bonding between them [78], which is much smaller than that for the weakly bonded graphene/Pt interface (3.7 Ă…) [79]. Based on the PES results, we can see that the corrugation induced by the chemical bonding between graphene and metal substrates can strongly modify the electronic properties of graphene.

Figure 18. C 1s photoelectron spectra (PES) taken at 400 eV for graphene on different metal surfaces. For graphene/Ru, the Ru 3d signal is subtracted [67].

2-5- Graphene/SiO2 Graphene/SiO2 is another important system to be considered due to its interesting electrical transport properties [50, 51, 80-84]. Therefore, understanding the 53


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electronic structure of graphene/SiO2 is a prerequisite to understand the transport properties of graphene and improve the performance of graphene-related devices [85]. On the other hand, the uniqueness of single-layer graphene can offer the opportunity to test the influence of core-hole effect on the electronic band structure of graphene [85]. Figure 19a shows the C K-edge XAS spectrum of graphene/SiO2, which is similar to the results from graphene/Cu and graphene/Pt, indicating the weak interaction between graphene and SiO2 [38, 61]. Figure 19b shows the RIXS spectra of graphene/SiO2 at different excitation energies. For comparison, the uppermost XES spectrum with the excitation energy of 320.0 eV is also shown. The RIXS spectra show strong dependence on the excitation energy. However, due to the presence of defects in graphene [38], strong electron-phonon scattering exists in the scattering process, which can be regarded as the incoherent part in the RIXS spectra. Therefore, the k-conserving resonant contribution (coherent part) is only part of the spectrum, and it contributes less for the spectra at higher excitation energies [86, 87]. The incoherent part, which can be viewed as the non-resonant XES contribution, is maximally subtracted from the RIXS spectra and leads to the indicated fractions of the coherent part of spectrum shown in Figure 19c.

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Figure 19. (a) C K-edge XAS spectra for graphene/SiO2. (b) XES and RIXS spectra recorded at different excitation energies as indicated in the figure. (c) Coherent fraction of the RIXS spectra in (b) [85].

Figure 20 shows a comparison of the XAS and XES spectra with different theoretical models. According to the final state rules, the XAS spectrum should represent the PDOS in the presence core hole. The presence of core hole significantly changes the PDOS and pulls a bound state out of the conduction band as illustrated in Figure 20. In addition, it can be found that the core hole can influence not only the atom on which the core hole is located but also its nearest neighbors. A shift of 1.7 eV can be found between the core-hole bound state and the π* peak in the pure graphene. While for XES, it should be represented by the PDOS without the core hole based on the final state rule. In order to align the PDOS with XES spectrum, the calculated spectrum has to be shifted down by 1.7 eV (Figure 20). This is because the core hole, which pulled down the PDOS by about 1.7 eV, is not present in the final state of XES. As a consequence, the experimental spectrum should be shifted up by 1.7 eV to undo the core-hole shift. Because we know precisely the relative position of Dirac point in the band structure to the 55

peak of the PDOS without core-hole, which is very close to the


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position of the

bound state with core hole, the Dirac point can be identified as

occurring at 284.7 eV in the XAS spectrum.

Figure 20. XAS and XES spectra compared with calculated results [85].

To understand the k-dependent RIXS spectra, we have simulated the RIXS spectra of graphene based on the Kramers-Heisenberg equation [88]. According the discussion above, the RIXS spectrum at excitation energy of 284.7 eV should correspond to the excitation from C 1s to Dirac point. In that case, the calculated spectrum should shift up by 283.0 eV as we did to the non-resonant XES spectrum because the core hole does not influence the position of the RIXS spectra. The calculated spectra accompanied with the experimental spectra are shown in Figure 21. The experimental spectra shown here are the RIXS spectra after removing the incoherent fractions based on the procedure mentioned earlier. The main features are labeled with letters for the ease of the following discussion.

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Figure 21. RIXS spectra compared to calculated spectra at different excitation energies [85].

At low excitation energies, the emission features are mainly from σ band. The low excitation energies are close to the Dirac point in the BZ. However, there is very little DOS at this point. Therefore, the spectral weight is very low. Here, it is worth to point out that the calculated spectra in Figure 21 are scaled by peak height, but the absolute intensity is weak for the low excitation energies. As increasing the excitation energy, the lower band (feature B) moves down and the upper one (feature A) moves up because the excitation energy is moved away from K towards Γ and M. At the excitation energy of 285.5 eV, a new feature C appears which disperses and grows in intensity at higher excitation energies. This is because we approach the M point in the BZ and there is a large DOS due to a saddle point in the band structure. The feature D in the theory corresponding to the

emission at M point is not visible in the experimental results. This may

be because that the emission by the π bands is symmetry forbidden since the and 57

states of the two carbon atoms in graphene unit cell have an opposite phase


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relation [45, 61]. On the other hand, it may also be due to the fact that it occurs in an energy range close to the strong elastic peak, and therefore the subtractions of the incoherent part and elastic peak may have removed the feature all together. At higher excitation energies, the relation between the spectral shape and k-points becomes less clear because the excitation energies intersect the conduction bands at various points including near Γ. In all, the dispersion of the features A, B and C is similar to the results of graphite and shows the similarity of the overall band structure of graphite and graphene.

3- Soft X-ray spectroscopy studies on graphene oxide (GO) and related materials 3-1- Chemically modified GO The GO, which is composed of a single-layer of graphite oxide and is usually produced by the oxidation of graphite by chemical method [89-93], is another specific branch of graphene research. The graphene prepared by reduction of GO also has various potential applications including sensing, catalysis and energy storage [94, 95]. To date, the atomic structure of GO is still unclear because of its nonstoichiometry, and several models have been proposed to study the functional groups and their arrangements across the carbon plane [96-99]. Therefore, studies of the fundamental electronic and structural properties of GO as well as the reduced GO are important for understanding its possible applications. Two general reduction approaches, chemical and thermal reduction, are proposed to achieve highly reduced GO materials [18, 92, 93, 100]. In this section, we will introduce the soft X-ray spectroscopy studies on chemically modified GO.

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Figure 22. Digital photographs of graphene oxide films (A) and KOH/hydrazine-reduced graphene films (B). SEM image of a KOH/hydrazine-reduced graphene film (C) and graphene oxide film (D) [101].

Figure 22 shows the digital photographs, optical microscopy images and scanning electron micrographs (SEM) of exfoliated GO and GO reduced by KOH/hydrazine on ITO substrates. The color of the GO films changes from light to dark for the thickness varied from 9.6 to 146 nm. After reduction with KOH and hydrazine, the obtained reduced GO suspensions are black in color and produce darker films upon electrophoretic deposition. The films are smoother upon deposition in constant-current mode as compared to deposition with applied voltage constant. The optical and SEM images indicate the uniform films over large area without any significant cracking. For the thicker electrophoretically deposited films, some corrugations at the boundaries between the graphene sheets are observed as shown in Figure 22. Figure 23a displays the C K-edge XAS spectra of the samples mentioned above. The lowest energy peak at 285.5 eV labeled as to the

can be attributed to transitions

symmetry around the M and L points in BZ. The sharp feature at 292.5

eV is ascribed to the transitions to dispersionless unoccupied states possessing symmetry at the Γ point of graphene BZ. Here, it should be noted that the intensity of 59

feature for the KOH/hydrazine-reduced graphene increases


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compared to the GO films before reduction, indicating the restoration of sp2 carbon network by reduction. The KOH/hydrazine-reduced graphene sample shows two additional distinctive features at ~297.5 and 300.2 eV arising from the transitions to K 2p3/2 and 2p1/2 states, respectively [101]. The peaks for K L-edge are shifted by ~3 eV compared with the literature values, indicating the formation of strongly ionic potassium-functional-group interaction for the KOH/hydrazinereduced graphene samples [102]. The GO films also show two spectral features at ~287.9 and 289.1 eV between the

and

resonances. The narrow structure at

289.1 eV can be attributed to the transitions to

resonance of C=O states. While

the relatively broad peak at ~287.9 eV may represent the transition to the resonance of C-O states derived from epoxide moieties. For the KOH/hydrazinereduced graphene films, the feature at 287.9 eV is far less prominent, indicating the reduction of GO films. In contrast, the transitions to

resonance of C=O

states at 288.9 eV originating from the carboxylate moieties at edge states are retained in the samples, indicating the inertness of COOH groups to hydrazine.

Figure 23. C K-edge (a) and O K-edge (b) for graphene oxide film and reduced graphene oxide films [101].

Further evidence for the reduction of GO by KOH/hydrazine is provided by O Kedge XAS spectra shown in Figure 23b. The spectra were only normalized by the pre-edge to show the difference in the edge jump heights of GO films before and after hydrazine reduction. The results clearly demonstrate the loss of oxygenated 60


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functional groups upon reduction. Two prominent features are observed at the O K-edge spectra: a weak peak at ~531 eV and a strong convoluted feature around 537 eV, which can be ascribed to transitions to C=O O-H and C-O

states and transitions to

states, respectively. Notably, after hydrazine reduction, it can be

clearly seen that intensity of

features decreases seriously, while the C=O

spectral feature is largely unaffected. This observation corroborates the role of hydrazine mainly reacting with the epoxide and phenolic moieties on graphene [101].

4-2- GO-S nanocomposite for high performance Li/S cells

Figure 24. SEM images of GO (a) and GO-S nanocomposite (b) [103].

Sulfur (S) is an attractive cathode material in rechargeable lithium batteries because it can yield almost the highest theoretical specific capacity of ~1675 mAh/g [8]. In addition, S is abundant in nature and environmental benignity. Therefore, Li/S batteries have the great potential for high-performance rechargeable lithium batteries [8, 103-107]. However, because S is insulated and its reaction products have high solubility in the electrolyte, S cannot be used as the cathode material for rechargeable lithium batteries directly [8]. We used a simple chemical method to prepare the graphene-oxide (GO-S) nanocomposite as the cathode material of Li/S cells [103]. Figure 24 shows the SEM images of the GO and GO-S nanocomposite after heating in Ar at 155 ยบC for 61


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12 hours. The layered nanostructures with porous structure of GO-S nanocomposite can be clearly seen in Figure 24b. The unique structure of GO-S can not only provides a good conductive matrix between GO and S for electron conduction but also accommodates large volume expansion/shrinkage during the discharge/charge process. This can be illustrated by the excellent electrochemical cyclability shown in Figure 25. The as-prepared nanocomposite shows a specific capacity of 1550 mAh/g in the initial cycle, and it remains above 900 mAh/g after more than 50 cycles, demonstrating its excellent electrochemical performance as a cathode material in Li/S cells.

Figure 25. (a) Galvanostatic discharge/charge profiles of GO-S nanocomposite at 0.02C rate; (b) cycling performance of GO-S nanocomposite at a constant current rate of 0.1C after initial activation processes at 0.02C (1C = 1675 mAh/g) for two cycles [100].

To better understand the properties of the GO-S nanocomposite, we investigated the electronic structure and chemical bonding between GO and S using XAS and XES. Figure 26 contrasts the C K-edge (a) and O K-edge (b) XAS spectra accompanied with their peak fittings for GO and GO-S. As we have discussed above, for C K-edge, the lowest energy peak can be attributed to transitions from the C 1s level to the

state, while the feature at ~292.8 eV is assigned to the

dispersionless Ďƒ* states at the Γ point. The peak position of the

for GO-S is 0.2

eV higher than that of GO, which is caused by the loss of oxygenated functional groups in GO [103]. In addition, the features between 62

and

states attributed


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to the carbon species in different carbon-oxygen functional groups are almost undetectable in the spectrum of GO-S. Instead, a broad feature centered at 287.5 eV, originating from the convolution of peaks of C-H and C-S bonds [108, 109], is observed, confirming the reduction of GO and formation of C-S bond in GO-S nanocomposite. Further evidence for the reduction of GO in GO-S can be found in the O K-edge XAS spectra of GO and GO-S as displayed in Figure 26b. The intensity of each feature for GO-S is weaker than that of the corresponding feature for GO. In addition, a new peak at ~537.0 eV attributed to the O-S bond appears in the spectrum of GO-S [110], indicating the formation of chemical bond between S and O for GO-S nanocomposite.

Figure 26. C K-edge (a) and O K-edge (b) of GO and GO-S nanocomposite [100].

Figure 27a shows the non-resonant and resonant C K-edge emission spectra for GO and GO-S nanocomposite, along with the spectra of HOPG for comparison. The non-resonant spectra for the three samples are very similar, indicating the analogous DOS among them. However, no strong band dispersive features can be observed in the RIXS spectra of GO and GO-S as that of HOPG, indicating that the oxygen functional groups in GO have pronounced influence on the valence band structure of graphene which will disturb the intrinsic crystal momentum during the scattering process [103].

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Figure 27. Resonant and non-resonant XES spectra of C K-edge (a) and O K-edge (b) of GO and GO-S nanocomposite [100].

It is also important to note that the RIXS spectra for GO and GO-S are similar, indicative of similar band structures for both materials. The O K-edge emission spectra of GO and GO-S are shown in Figure 27b. Two main features at 527.0 and 522.8 eV can be observed, which are mainly attributed to the O 2p-C 2p state and O 2p-C 2s state, respectively [100]. The shape of the low energy feature changes obviously at different excitation energies, which can be caused by the different chemical environments of oxygen in different functional groups. For GO-S, when the excitation energy is 538.0 eV, the state from the O-S bond can also contribute to the emission spectrum, resulting in the increased intensity at lower emission energies when comparing with the spectrum of GO at the same excitation energy. Same phenomenon can also be observed in the non-resonant XES spectra. Moreover, the main peak in the non-resonant spectrum of GO-S is found to shift about 0.5 eV towards the lower emission energy, which may be caused by the lower binding energy of O-S bond compared with those O species from other functional groups. Based on the results, we can see that the high performance of Li-S cells using the GO-S nanocomposite as the cathode material may be related 64


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to the following factors: (1) the S can partially reduce the GO and then improve the conductivity of the GO and (2) the mild interaction between GO and S can stabilize S by bonding with the GO sheet.

Conclusions Carbon allotropes such as graphene and carbon nanotubes have been developed from basic research to prospective applications. However, several challenges for their large-area and high-quality preparation are yet to be addressed. For the case of graphene, CVD and chemical approach are the two most promising ways towards the controlled synthesis of graphene with designed morphology and structure. For the CVD method, graphene layers are epitaxially grown on the transitional metal surfaces. Therefore, understanding the interfacial interaction between graphene and metal surfaces is of great importance for both understanding the growth mechanism of graphene and the development of high-performance graphene-based devices. It has been shown that soft X-ray spectroscopy, including XAS, XES, and RIXS, can serve as a powerful technique for investigating the electronic properties of carbon allotropes. From the results of graphene on different metal surfaces investigated by soft X-ray spectroscopy, we can see that the electronic structure of graphene on different metal surfaces is totally different, which can be directly related to the different strength of hybridization between graphene and metal substrates. These significant spectra changes have been proved to be an effective measure for the bonding strength of graphene/metal interface: strong band dispersion can be observed when the interaction between graphene and metal substrate is weak, while the band dispersion is seriously disturbed when a strong hybridization between graphene and metal surface exists. GO has also attracted intense attention because of its wide use as a precursor material for the mass production of graphene-based composite materials, which 65


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hold great potential in various practical applications. Insight into the electronic structure of GO, reduced GO and GO-based composite material are essential for better understanding the distribution of functional groups, defects and structure evolution during the reduction process. Soft X-ray spectroscopy studies can supply the electronic information on both conduction band and valence band of GO and GO based material, which can help us shed more light on the reduction mechanism of GO and understanding the structural and electronic properties of GO-based composite.

Acknowledgements J. F. Zhu thanks the financial supports from the Natural Science Foundation of China (Grant No. 21173200), National Basic Research Program of China (2010CB923302) and the Specialized Research Fund for the Doctoral Program of Higher Education of Ministry of Education (Grant No. 20113402110029). L. Zhang thanks the financial support from the Scholarship Award for Excellent Doctoral Student Granted by Ministry of Education of China.

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Chapter

3 Biocompatible Conjugation of Hydrogels for Regenerative Medicine Huaping Tan School of Materials Science and Engineering, Nanjing University of Science and Technology, Nanjing 210094, China E-mail :

hptan@njust.edu.cn

B

iodegradable and biocompatible hydrogels have many different applications in the field of regenerative medicine. The hydrogels could be utilized as delivery systems, cell carriers, and scaffolds for drug delivery and tissue engineering. The hydrogels are appealing scaffolds because they are structurally similar to the extracellular matrix of many tissues, can often be processed under relatively mild conditions, and may be delivered in a minimally invasive manner. This chapter will discuss recent advances in the field of biodegradable and biocompatible hydrogels, including both physical and covalent crosslinking, which can be potentially used in regenerative medicine applications.

Introduction Regenerative medicine, which combines of tissue engineering and drug delivery, takes the multidisciplinary principles of materials science, medicine, and life science to generate tissues and organs of better biological structures and functions (Liu et al., 2011). Regenerative medicine is to implant scaffolding materials for regenerating the tissue based on the recruitment of native cells into the scaffold, and subsequent deposition of extracellular matrix (ECM). Cell scaffolds play a 77


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crucial role because they act as an artificial ECM to provide a temporary environment for cell supporting to infiltrate, adhere, proliferate and differentiate (Langer et al., 1993). The scaffolds provide the initial structural support and retain cells in the defective area for cell growth, thus playing an important role during the development of engineered tissues (Tan et al., 2009). A major issue in the use of scaffolds and tissue reconstruction procedures in general, is that the establishment of a biomimic environment into the engineered tissue is necessary for engineered tissue survival (Tan et al., 2007). Hydrogels are promising substrates for regenerative medicine applications due to high tissue-like water content, ability to homogeneously encapsulate cells, efficient mass transfer, easily manipulated physical properties and minimally invasive delivery (Drury et al., 2003). Highly hydrated hydrogels can better mimic the chemical and physical environments of ECM and therefore are ideally cellular microenvironment for cell growth. Therefore, many biodegradable hydrogels have served as highly functional scaffolds for cells and growth factors (GFs) delivery which increase the effectiveness of the GFs in regulation of cell proliferation and migration (Tememoff et al., 2000). The nature of hydrogels provides the attractive feature of facile and homogenous cell distribution within any defect size or shape prior to gelation. Furthermore, hydrogels allow good physical integration into the defect, potentially avoiding an open surgery procedure and facilitating the use of minimally invasive approaches for material and cell delivery (Hou et al., 2004). The hydrogel precursor loaded with growth factors and/or targeted cells can be injected into the wound site and experiences a solution-to-gelation transition (sol– gel) in situ due to physical or chemical stimuli (Tan et al., 2010). Cells are isolated from a small biopsy, expanded in vitro, and encapsulated in hydrogel precursors, which are subsequently transplanted into the patient by injection. The hydrogels provide the initial structural support and retain cells in the defective area for cell growth, metabolism and new ECM synthesis to restore the damaged tissue.

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Over the past decade, many methods have been employed for preparation of biodegradable hydrogels. Conventionally, biodegradable hydrogels have been synthesized by chemical crosslink methods such as crosslinking agents (e.g., glutaraldehyde, genipin, adipic dihydrazide and bis-sulfosuccinimidyl suberate) and free radical polymerization. Recent developments have utilized metalmediated ‘click’ chemistry methods to prepare hydrogels, the most common example being the copper (I)-catalysed reaction of azides with alkynes (Tan et al., 2011). However, these methods are limited by involvement of cytotoxic reagents, instability of the functional groups, possible side reactions and low coupling efficiency, which are unsuitable for GFs delivery and cell encapsulation. There is a continuing need to exploit simple, specific and highly efficient conjugation methods which are applicable to a broad class of biocompatible hydrogels with full preservation of bioactive function for regenerative medicine. Table 1. Biocompatible conjugation of hydrogels. Hydrogel Types

Physical hydrogels

Gelation Mechanism

Polymers

Affinity interaction

PEG/Hyaluronic acid

Nucleobase pairing (Hydrogen bonding)

Covalent hydrogels

PEG

Schiff-base reaction

Chitosan/Hyaluronic acid

Diels-Alder addition

Hyaluronic acid

Michael-type addition

PEG/Hyaluronic acid

Biocompatible hydrogels with potential applications in regenerative medicine can be classified into physical and covalent gels, according to their gelation mechanism (Tan et al., 2010). A variety of naturally- and synthetically-derived materials have been utilized to prepare hydrogels. Some reported biocompatible hydrogels systems are listed in the Table 1. The hydrogel network crosslinked by physical association between polymeric chains is the so-called physical gel, while the formation of a covalent gel takes place via covalent bonds between polymeric chains. To develop a suitable hydrogel as a cell carrier, the degradation rate and 79


Advances in Science

mechanical properties of the hydrogel must complement the tissue growth and natural ECM. In general, these properties can be fine-tuned through variations in the chemical structure and crosslinking density in hydrogels. For a given hydrogel system, activities of seeded cells can be regulated by attaching specific bioactive moieties to the polymer matrix backbone. Comprised of various ECM-like macromolecules and proteins, hydrogels control the tissue structure, regulating the function of the cells (Tan et al., 2010).

1- Materials for hydrogels 1-1- Natural materials Rapid Naturally-derived hydrogel forming polymers have frequently been used in regenerative medicine applications because they are either components of or have macromolecular properties similar to the natural ECM. However, natural polymers often undergo rapid degradation upon contact with body fluids or medium. Therefore, limitations of natural hydrogels have motivated approaches to modify these polymers as well as to utilize various synthetic polymers. Representative naturally derived polymers include proteins (e.g., collagen, gelatin and fibrin) and polysaccharides (e.g., chitosan, hyaluronic acid, chondroitin sulfate, agarose and alginate). Chitosan is a linear polysaccharide, which is a partially deacetylated derivative of chitin.

This

polycationic polysaccharide contains

glucosamine

and

N-

acetylglucosamine molecules, thus structurally similar to naturally occurring glycosaminoglycans

(GAGs).

Chitosan

is

considered

a

biodegradable

polysaccharide, which can be metabolized by human enzymes such as lysozyme. Chitosan has been investigated for a variety of tissue engineering applications in recent years due to its biocompatibility, biodegradability, low immunogenicity and cationic nature. However, unmodified chitosan can only be dissolved in acidic solutions due to its strong intermolecular hydrogen bonds, which limits its applications as an injectable hydrogel. Water-soluble chitosan derivatives support 80


Advances in Science

cell growth, and composites of chitosan and GAG or other bioactive proteins are able to create suitable biomimetic microenvironments for cell implantation. Hyaluronic acid (HA) is a naturally occurring non-sulfated GAG, consisting of multiple repeating disaccharide units of N-acetyl-D-glucosamine and Dglucuronic acid, that is widely distributed throughout the ECM of all connective tissues in human and other animals. HA plays an essential role in many biological processes such as tissue hydration, nutrient diffusion, proteoglycan organization, and cell differentiation. HA is especially prevalent during wound healing and in the synovial fluid of joints. HA is naturally degraded by hyaluronidase, which is ubiquitous in cells and in serum. Due to its good biocompatibility, biodegradability, as well as excellent gel-forming properties, HA and its derivatives have been widely explored as hydrogels for regenerative medicine.

1.2 Synthetic materials Synthetic polymers are appealing for hydrogels because their chemical and physical properties are typically more controllable and reproducible than those of natural polymers. Synthetic polymers can be reproducibly produced with specific block structures, molecular weights, and degradable linkages. Compared to natural hydrogels, synthetic hydrogels offer improved control of the matrix architecture and chemical composition, but tend to have lower biological activity. An appealing and effective strategy is to incorporate bioactive species such as cell growth factors, peptides and proteins into synthetic materials, resulting in biomimetic hydrogels with bioactive functions for optimal cell response. Synthetically-derived

materials

include

poly(ethylene

glycol)

(PEG),

poly(propylene fumarate) (PPF) and polypeptides, which are among the most widely used synthetic polymers for biodegradable hydrogels. Although many variations of synthetic polymers can form biocompatible and biodegradable hydrogels, PEG remains one of the most widely investigated systems. PEG is currently FDA-approved for several medical applications. Biodegradable PEG hydrogels can be obtained via copolymerization with other 81


Advances in Science

degradable polymers such as poly(lactic acid) (PLA), poly(glycolic acid) (PGA) and PPF. PEG hydrogels have been used as cell scaffolds, adhesive medical applications, and delivery vehicles with promising results. Particularly, the ability to control the crosslinking density provides the flexibility and tailorability to PEG-based hydrogels for cell encapsulation and tissue growth. PEG and the chemically similar poly(ethylene oxide) (PEO) are hydrophilic polymers that can be photocrosslinked by modifying each end of the polymer with either acrylates or methacrylates. Furthermore, many naturally occurring biopolymers, such as HA, fibrinogen, chitosan and heparin, are also generally examined in combination with PEG hydrogels.

2- Hydrogel systems 2-1- Physical crosslinking 2-1-1- Affinity interaction HIP peptide

LMWH

Backbone of polysaccharide Figure 1. Schematic of noncovalent assembly hyaluronic acid hydrogel via affinity interaction with flexible star PEG grafts.

Noncovalent assembly of polymeric materials via specific molecular recognition interactions such as protein-protein, peptide-peptide, and ionic crosslinking interactions has become increasingly prominent in the production of hydrogels. The ability of heparin and related GAGs to sequester and stabilize GFs has been exploited in the production of surfaces and covalently crosslinked hydrogels that 82


Advances in Science

can mediate cell proliferation and migration. The delivery system can rely on passive inclusion of HIP in hydrogel matrices via affinity interaction with heparin or covalent attachment of heparin to a drug delivery matrix. Recently, low molecular weight heparin (LMWH) and heparin interacting protein (HIP) or vascular endothelial growth factor (VEGF) have been transiently immobilized in a star PEG matrix via noncovalent interactions between the HIP/VEGF and covalently incorporated heparin. The studies suggest that the star PEG-heparin copolymer and its incorporation into hydrogel networks are capable of GF delivery. The assembly of noncovalently associated HA hydrogel was produced via the interaction

of

LMWH

modified

HA

derivative

and

peptide

CRPKAKAKAKAKDQTK, a GF sequence derived from the heparin-binding domain of HIP (Tan et al., 2012). The hydrogel precursor was designed with LMWH on HA backbone, which was accessible for conjugating with HIP in aqueous environment. The LMWH-functionalized HA (HA-LMWH) was employed via Michael addition of thiol functionalized HA to maleimidefunctionalized LMWH. The HIP-functionalized PEG (PEG-HIP) was employed via Michael addition of vinyl sulfone terminated four-arm PEG to HIP, which contains flexible star PEG grafts with HIP located at the PEG termini. The terminal LMWH of HA-LMWH was easily accessible to HIP in the aqueous solution, thus enabling affinity interactions at the hydrophilic PEG termini. Nanofibrous hydrogels were formed via the mixing of homogeneous, low viscosity solutions of HA-LMWH and PEG-HIP in PBS. Addition of a solution of PEGHIP to the HA-LMWH solution was pipeted to ensure homogeneity and immediately resulted in the formation of a self-supporting, viscoelastic hydrogel.

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Figure 2. (a) Microrheological characterization of the viscoelastic properties of HIP-mediated noncovalent assembly hydrogels. (b) SEM images characterized the morphologies of HIP mediated HA-LMWH/PEG-HIP hydrogel.

To study the rheological characteristics of the HIP mediated noncovalent assembly hydrogels, optical tweezer microrheological of polymer networks was shown in Figure 2a. The apparently low viscosity of HA-LMWH and PEG-HIP solutions exhibited storage moduli, G’(ω) ≈ 0.9~1.1 Pa, in excess of the loss moduli G”(ω) at low frequencies, indicating weak viscoelastic materials. Upon addition of PEG-HIP to the HA-LMWH solution, an increase in elastic modulus was observed, with statistically significant. HA-LMWH/PEG-HIP hydrogel with HA-LMWH concentration of 5 wt% resulted in elastic gels indicating a G’(ω) > 2.3 Pa. Compared the storage moduli of the HA-LMWH/PEG-HIP hydrogel with HA-LMWH and PEG-HIP solutions, it clearly demonstrated the effective crosslinking in the HA-LMWH/PEG-HIP hydrogel. Scanning electron microscope (SEM) image characterized the spongelike morphology of HIP mediated HA-LMWH/PEG-HIP hydrogel (Figure 2b). The hydrogel with the interconnected macrodomains comprising of the polymer and the pores originally occupied by water was observed. According to cross-sectional image, the HA-LMWH/PEG-HIP hydrogel structures displayed a continuous, porous and nano-fibrous structure after removing water by lyophilization, with the fiber sizes of 50~70 nm. 84


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The nano-fibrous structure was attributed to the noncovalent assembly mechanism, resembling other self-assembly macromolecular hydrogel system structures. The degradation properties of HA-LMWH/PEG-HIP hydrogel were monitored in PBS at 37oC as a function of incubation time. The HA-LMWH solution, and the other HIP mediated hydrogel without flexible PEG chains, e. g. HA-LMWH/HIP, were also investigated as control groups. As shown in Figure 3a, HA-LMWH/HIP was dismissed in 2 days, while HA-LMWH/PEG-HIP hydrogel lost its weight steadily up to 14 days. At day 14, the weight remaining ratio of the HALMWH/PEG-HIP hydrogel was 62%. Since swelling properties of hydrogel scaffolds are crucial for substance exchange during ASCs differentiation, the swelling kinetics of the HA-LMWH, HA-LMWH/HIP, and HA-LMWH/PEGHIP were investigated in PBS (Figure 3b). The results showed that both of the HA-LMWH and HA-LMWH/HIP dismissed in 2 days, which was consistent with the data of weight loss shown in Figure 3a. The swelling ratio of the HALMWH/PEG-HIP hydrogel significantly increased from 31.2 to 40.4 when incubated for 14 days. These results clearly demonstrated that the HIP alone was unable to interlink the HA-LMWH, while the PEG-HIP can serve as an elastic cross-link in noncovalently assembled nano-fibrous hydrogel networks, which was accessible for conjugating with LMWH in aqueous environment.

a

b

120

50

80 60 HA-LMWH HA-LMWH/HIP HA-LMWH/PEG-HIP

40 20 0

0

7 Time (days)

14

Swelling ratio

Hydrogel mass (%)

100

HA-LMWH/HIP HA-LMWH/PEG-HIP

40

30 0

7

14

Time (days)

Figure 3. (a) Degradation of hydrogels in PBS at 37째C with respect to weight loss. (b) Swelling kinetics of hydrogels in PBS as a function of incubation time.

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The potential application of the nano-fibrous HA hydrogel as an injectable scaffold in soft tissue engineering was studied by encapsulation behavior of human adipose-derived stem cells (ASCs). ASCs have great potential in the field of regenerative medicine. ASCs have been used in addition with various scaffolds and bioactive signaling molecules, such as GFs, to engineer human tissue. There is evidence that the ability of ASCs to grow and differentiate varies among physical structures and changes with bioactivities of scaffolds. Defining these variations in ASCs function and molecular mechanisms of adipogenesis will facilitate the development of cell-based therapies. The synthesis of polysaccharide hydrogels via specific heparin-peptide binding affinity and that might therefore permit manipulation of inductive properties for ASCs differentiation. The bioactivity of the HA-LMWH/PEG-HIP hydrogel was assessed in cellproliferation assays. The proliferation of ASCs in the presence of HA-LMWH, HA-LMWH/HIP, and HA-LMWH/PEG-HIP was monitored over time via DNA contents. The HA-LMWH and HA-LMWH/HIP did not cause significant increase in cell proliferation in vitro, while the HA-LMWH/PEG-HIP hydrogel showed a statistically significant increase in proliferation over the other samples as early as day 3. The cytocompatibility study showed the HIP mediated HA-LMWH/PEGHIP hydrogel was non-cytotoxic. Furthermore, the HIP mediated nano-fibrous hydrogels also induced differentiation of ASCs, and coupled with the cell proliferation data, suggested a novel potential mechanism for targeted differentiation of stem cells via the therapeutic affinity interaction cross-links.

2-1-2- Nucleobase pairing Self-assembly conjugation in polymeric materials has become increasingly prominent in the production of bioactive surface and matrix. A great variety of self-assembly have been reported for conjugation of peptides, proteins, DNA, oligonucleotides and nanoparticles. More recently, a biological hydrogel was selfassembled via the Watson-Crick base pairing of thymine and adenine functionalized star PEG (Tan et al., 2012). The biological self-assembly PEG 86


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hydrogel system is established by the Watson-Crick base pairing between thymine (T) and adenine (A) via the hydrogen bonding (Figure 4). Compared with linear PEG, multi-arm star-shaped PEG has the nature to induce stereo-complex formation, which shows promise in biomedically-relevant hydrogel systems. Firstly, thiol thymine (T-SH) and thiol adenine (A-SH) were synthesized respectively. Maleimide terminated four-arm PEG (PEG-Mal) was functionalized with either T-SH or A-SH functionalities as self-assembly precursors (referred as PEG-T and PEG-A) via the Michael-type addition. After dissolution and mixture of precursors in an aqueous environment, a stereo-PEG hydrogel network was self-assembled due to the formation of pairing complexes. 4-arm PEG-thymine

+

Self-assembly

4-arm PEG-adenine

Ideal network 3D gel

Figure 4. Schematic of biological self-assembly of ideal four-arm PEG hydrogel network via the intermolecular hydrogen bonding of the Watson-Crick base pairing between thymine and adenine functionalities.

Dynamic time sweep rheological experiments were conducted to monitor network evolution and evaluate viscoelastic behavior of this hydrogel (Figure 5a). Upon mixture of homogeneous PEG-T and PEG-A solutions (20 w/w%) at 37 oC, the storage modulus G′ was significantly increased and larger than the loss modulus G″ after 150 sec. The data indicate a final storage modulus G′ value of ~600 Pa after 500 sec, in excess of the loss moduli G”(ω), which resulted in a hydrogel with an typical viscoelastic property. The inset in Figure 5a illustrated the viscoelastic hydrogel after self-assembly in PBS. Two different molecular weights 87


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(10 KDa and 20 KDa) of four-arm PEG were employed for self-assembly in aqueous solution at 37oC. Three self-assembly hydrogels consisting of PEGT/PEG-A with molecular weights of 10 KDa or 20 KDa, e. g. PEG-10K/10K, PEG-10K/20K and PEG-20K/20K were prepared. The PEG-20K/20K hydrogel showed the highest storage modulus, which followed by the PEG-10K/20K and PEG-10K/10K hydrogels. Moreover, higher concentrations of polymer created a higher effective cross-link density and therefore leading to a tighter network. Incubation temperature has a significant influence on weight loss of self-assembly hydrogels. At a lower temperature, the hydrogel showed a slower weight loss rate than that at a higher temperature. At 50°C, the hydrogel showed a significantly faster weight losing rate than the others due to the dismission of hydrogen bonding, which totally dissolved in 9 days. Compared to at 50 oC, the hydrogels lost their weight steadily up to 18 and 24 days at 37°C and 20°C, respectively. The hydrogel at a lower temperature showed a lower swelling ratio rate than the one at a higher temperature, which was consistent with the data of weight loss shown in Figure 5b. Compared to the 20oC and 37oC, the swelling ratio of hydrogel significantly increased after 24h incubation at 50°C. At 37°C, swelling ratio of the hydrogel significantly increased from 17.3 to 26.8 after incubated for 14 days.

a

b o

100

Hydrogel Mass (%)

G' G" (Pa)

100 10 1 0.1

20 C o 37 C o 50 C

0

200

400 Time (s)

600

60

20 0

7

14

21

Time (day)

Figure 5. (a) Rheological characterization of the self-assembly hydrogel network. Inset image illustrates a formed hydrogel in PBS. (b) Degradation of hydrogel at different temperature in PBS.

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GFs that induce ASCs proliferation include insulin, dexamethasone (Dex) and VEGF. The biological hydrogel showed the highest efficiency loading of VEGF, and about 96% of VEGF could be entrapped within the hydrogel, but there was no significant difference compared with dexamethasone and insulin (Tan et al., 2012). The in vitro cumulative amount of GFs released from hydrogel at specific time points was quantified. The GFs incubated in PBS demonstrated a total cumulative release of approximately 68~89% over the 14-day time period, then totally released in 18~20 days. The release of dexamethasone showed a significantly faster rate than both of insulin and VEGF. The adhesion of ASCs to top surface of control hydrogel and GFs entraped hydrogels after culture for 24h was also characterized (Tan et al., 2012). Cell number on the GFs entraped hydrogel surfaces was greater than that of control hydrogel, but no significant difference was found. All hydrogels caused significant increase in cell proliferation in vitro after 7 days incubation, while the insulin loaded hydrogel showed a significant increase in cell proliferation over the control hydrogel. The potential application of this biological hydrogel as a cell scaffold in soft tissue engineering was also confirmed by encapsulation behavior of human ASCs. These unique characteristics of the biological hydrogel make it a promising candidate as an injectable scaffold for pharmaceutical and biomedical applications.

2-2- Chemical crosslinking 2-2-1- Schiff-base reaction The Schiff-base reaction has been utilized to prepare a polysaccharide hydrogel upon mixing, without employing any extraneous chemical crosslinking agents (Tan et al., 2009). The gelation is attributed to the Schiff-base reaction between amino groups of N-Succinyl-chitosan (S-CS) and aldehyde groups of oxidized HA (A-HA) (Figure 6). S-CS, a water soluble chitosan derivative, was synthesized via introduction of succinyl groups at the N-position of the glucosamine units of chitosan. HA can be oxidized, and the carbon-carbon bonds of the cis-diol groups in molecular chain are cleaved and generate reactive aldehyde functions, which 89


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can develop chemical crosslinking action with amino functions via the Schiff-base linkage.

=

= NH 2

+

N-succinyl-chitosan

CHO

aldehyde hyaluronic acid

CH N

Figure 6. Scheme of N-succinyl-chitosan and aldehyde hyaluronic acid composite hydrogel via the Schiffbase reaction.

Five different volume ratios of S-CS/A-HA composite hydrogels were prepared by the Schiff-base crosslinking reaction (–C=N–) between –NH2 of S-CS and – CHO of A-HA. Ratios of S-CS/A-HA did not significantly influence the gelation time, i.e. the gelation time for all five groups were within 1~4min. However, the gelation time was shorter when the ratios of S-CS/A-HA were 3/7~7/3. The crosssectional SEM images of the freeze-dried 3/7 and 7/3 composite hydrogels demonstrated that a higher content of S-CS results in the formation of smaller pore diameters and tighter network structure in composite hydrogels, which is likely due to the comparatively sufficient crosslinking.

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Table 2. Compressive modulus of S-CS/A-HA composite hydrogels at room temperature.

%(v/v) S-CS

%(v/v) A-HA

Compressive modulus (kPa)

30

70

12 ± 4

50

50

25 ± 3

70

30

28 ± 6

Compressive modulus is particularly important for cartilage tissue engineering. The compressive modulus of composite hydrogels were improved with increasing ratios of S-CS content from 30% to 70% (v/v) (Table 2). With increasing ratios of S-CS content from 30% to 70% (v/v), the compressive modulus of the composite hydrogels were improved correspondingly. The 5/5 and 7/3 hydrogels had significantly larger compressive modulus than the 3/7 hydrogel, which were 25 and 28 kPa, respectively, whereas no difference was found between them. However, measures must be taken in the future to further improve the mechanical strength of the present systems if they are used as the injectable scaffold for cartilage restoration, although many types of the hydrogels with similar strength have been diversely used for the same purpose.

a

b

120

*

40000

Cell Number

90

Hydrogel Mass (%)

50000

60 1:9 5:5 9:1

30

3:7 7:3

*

*

30000

20000

0 10000 0

7

14

21

28

Time (day)

TCP

3/7

5/5

7/3

9/1

S-CS/A-HA

Figure 7. (a) Degradation of S-CS/A-HA composite hydrogel in PBS at 37°C with respect to weight loss. (b) Number of bovine chondrocytes that adhered to the surface of S-CS/A-HA composite hydrogels versus control wells.

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Crosslinking density influences many of the macroscopic properties of hydrogels. In general, an increase in the crosslinking density results in a decrease in the water content and mass weight loss. The 3/7, 5/5 and 7/3 hydrogels showed a slower degradation rate than the 1/9 and 9/1 hydrogels in PBS at 37°C (Figure 7a), which is possibly due to the sufficient crosslinking and subsequent microstructure. Although some reports have indicated that Schiff-base crosslinking structure of HA is instable, interestingly, the S-CS/A-HA hydrogels were stable in PBS during the four weekt period, which indicated that the hydrolysis rate of S-CS/A-HA Schiff-base is slow under physiological conditions. Swelling properties of the as-prepared hydrogels are crucial for substance exchange when they are used as injectable scaffolds for biomedical applications. Both S-CS and A-HA have an abundant number of hydrophilic groups, such as hydroxyl, amino and carboxyl groups, which can easily produce hydration with water. The amount of S-CS and A-HA used in the hydrogel synthesis significantly affects the hydrogels’ swelling properties. The 1/9 and 9/1 hydrogels revealed less crosslinking, consequently increasing the exposure of polymer chains to water molecules, and significantly leading to water absorption enhancement and faster weight loss. Another important issue for injectable hydrogels is the encapsulation efficiency for the target cells. Since the microstructure and high water content are very similar to that of the extracellular matrix of natural cartilage, hydrogels may preserve the phenotype of chondrocytes. The bovine chondrocyte attachment study indicated that the hydrogels support cell adhesion. A significantly higher number of attached chondrocytes were observed on the 5/5 and 7/3 composite hydrogels than the 3/7 and 9/1 hydrogels (Figure 7b). Furthermore, encapsulation of chondrocytes demonstrated that the composite hydrogel promoted cell survival and the cells retained regular, chondrocytic spherical morphology. These preliminary studies indicate that the composite hydrogel supports chondrocyte adhesion and encapsulation, and may have potential uses in cartilage regeneration 92


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applications. As this process of hydrogel formation is simple, feasible, and usually performed under mild conditions without employing any extraneous toxic crosslinking agents, we believe that such a composite matrix will have potential applications in wound management, drug delivery, tissue engineering, and other related biomedical fields.

2-2-2- Diels-Alder addition The site-specific conjugation of biomolecules with other molecules or functional groups is important for subsequent biological investigations. The Diels-Alder reaction in aqueous environments, which involve a highly selective [4+2] cycloaddition reaction between a diene and a dienophile, is diverse in scope and efficient in reactivity, results in very high yields, produces no byproducts, and occurs under mild reaction conditions. Since the aqueous Diels-Alder chemistry occurs readily without the involvement of catalysts and coupling reagents, it provides a competitive alternative to metal-mediated ‘click’ chemistry. The compatibility of aqueous Diels-Alder chemistry with biomolecules has been exploited

elegantly

in

the

bioconjugation

of

protein,

peptides

and

oligonucleotides, which were site-specific without interference of the many functional groups present in polysaccharide backbone. Given the high specificity and efficiency, the success of aqueous Diels-Alder chemistry in biomolecule conjugation and immobilization may be extended to the synthesis of biodegradable hydrogels.

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[

Sol

]x

Functionalized hyaluronic acid derivatives

Aqueous

37oC

[

]x

Gel Figure 8. Schematic of hyaluronic acid hydrogel formation via the Diels-Alder reaction.

A HA derivative hydrogel with novel structures has been developed via aqueous Diels-Alder cycloaddition reaction of bioconjugation that specifically allows for biopharmaceuticals delivery (Tan et al., 2011). Hydrogel precursors were designed with furan groups (diene) on the outer PEG, corona which are accessible for reaction with maleimide (dienophile) functionalized HA derivative in aqueous environment at 37oC. For construction of maleimide-functionalized precursor, HA sodium was oxidized by sodium periodate, and the carbon-carbon bonds of the cis-diol groups in molecular chain are cleaved and generate reactive aldehyde functions, which conjugate with the maleimide containing molecule via the Schiff-base linkage. The furan-functionalized HA derivative contains a biodegradable backbone of polysaccharide and hydrophilic grafts of PEG with furan groups located at the PEG termini. The terminal furan groups are easily accessible to maleimide-functionalized HA derivative in the aqueous solution, thus enabling cycloaddition at the PEG termini. Both resulting maleimide- and 94


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furan-functionalized HA derivatives are water soluble at various pHs. Biodegradable hydrogels were generated by mixing maleimide- and furanfunctionalized HA derivative solutions (3 wt%) containing a 1:1 ratio of furan to maleimide functionalities at 37oC. Ultimately, this Diels-Alder gel composition affords a high water content, elasticity similar to many tissue matrices.

a

b

20 G' G''

12

Viscosity (Pa.S)

G', G'' (Pa)

15

15

10

5

9

6

3 0

0 0

500

1000

1500

2000

2500

0

500

1000

Time (s)

1500

2000

2500

Time (s)

Figure 9. (a) Rheology was used to monitor dynamic network formation and indicates gelation within minutes and complete reaction occurring in less than 1 h at 37°C. (b) Viscosity of polysaccharide mixture increased as a function of polymerization time.

To study the mechanical characteristics of the hydrogel, dynamic time sweep rheological experiments were conducted to monitor network evolution and evaluate viscoelastic behavior of the polysaccharide hydrogel. As shown in Figure 9a, the rheological analysis showed that the storage modulus G′ was significantly increased and larger than the loss modulus G″ after 1300 sec, an indication of an elastic rather than viscous material. Furthermore, the data indicate a final storage modulus G′ value of 18.1 ± 0.7 Pa, signifying a structurally robust network that maintains its 3D shape with loading. The value of the storage modulus G′ exceeds that of the loss modulus G″ indicating the formation of a strong and elastic hydrogel. Such rheological behavior is characteristic of solid like gel material, hence, it may be suitable for applications that request a soft characteristics. The kinetics of hydrogel formation were also measured by monitoring the viscosity as a function of reaction time (Figure 9b). The transition from liquid-like behavior to elastic gel-like behavior occurs which was observed after approximately 26 min. 95


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After incubation over 33 min, the viscosity values increased rapidly and then leveled off after 40 min, indicating that the crosslinking of polysaccharide was essentially completed. The degradation properties of the hydrogel were monitored as a function of incubation time in PBS at 37oC (Figure 10a). The hyaluronic acid derivative hydrogel lost its weight steadily up to 21 days. At day 21, the weight remaining ratio of hydrogel was 43%. Since hyaluronic acid is an enzymatically degradable polysaccharide, the enzymatic resistance of hydrogels was also investigated in a hyaluronidase solution (100 U/ml PBS) as a function of incubation time. The hyaluronidase resulted in a significant mass loss, and the hydrogels were degraded in 11 days. Swelling properties of the as-prepared hydrogel are crucial for substance exchange when used for biomedical applications. To examine responses to external stimuli, the swelling kinetics of the hydrogels were investigated in response to step changes in incubation temperature in PBS. The results showed that the gel underwent significant volume changes with temperature, with the gel swelling when the temperature was increased from 20oC to 80oC. The swelling ratio of the gel linearly increased from 28.4 to 43.7, not fragmenting at any of these temperatures. A cooling process was also performed by temperature-induced changes, with temperature values < 60oC leading to rapid gel shrinking.

a

b 100

PBS 100U hyaluronidase/PBS

80 Cumulative release (%)

80 Hydrogel mass (%)

100

60 40 20 0

60

40 lysozyme insulin

20 0

0

5

10

15

20

0

Time (days)

4

8

12

16

20

Time (days)

Figure 10. (a) Degradation of hydrogel in PBS with or without 100 U/ml hyaluronidase at 37째C with respect to weight loss. (b) Cumulative release profiles of insulin and lysozyme from Diels-Alder hydrogel in PBS at 37oC.

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Delivery of drugs or proteins in a biodegradable hydrogel network is of great interest for the development of biosensors and medical materials. Two model proteins, insulin (negatively charged) and lysozyme (positively charged) were encapsulated into the hydrogel. To evaluate the ability of effectively deliver drugs, the in vitro release behavior of insulin and lysozyme from the hydrogel were examined (Figure 10b). Approximately 22.8% of insulin was seen to be released from the hydrogel after 24 h of incubation. However, in the lysozymeencapsulated hydrogel, there was a sustained release of lysozyme and also low initial burst release (~3.4% at the first day). In comparison to the high burst release (first 3 days, 42-56%) and short duration (9 days, up to 70-85% insulin release) from hydrogels, the polysaccharide hydrogel demonstrated a decrease in the burst release to a low level (~17% at the first three days) and increase the sustained release to a long duration (21 days, up to 60% lysozyme release). The release of insulin and lysozyme detected from hydrogel exhibited different kinetic profiles, demonstrating that the charges of protein could also be taken as an adjustable factor of release profiles. As expected, the charges of encapsulated proteins influenced gelation formation due to the electrostatic interactions. Cytocompatibility of hydrogel and the bioactivity of insulin released from gel were assessed in cell-proliferation assays. ASCs proliferated on the insulin loaded hydrogels after 5 days incubation, and the DNA content of insulin loaded hydrogels was significantly greater than that of pure hydrogels, whereas no difference was found between hydrogels without insulin. The ASCs (>95%) were shown to survive in pure and insulin loaded polysaccharide derivative hydrogels. Roundly shaped ASCs were distributed on the pure and insulin loaded polysaccharide derivative hydrogels, and migrated into the gels, indicative of the highly hydrophilic nature of the gels. The ASCs residing in the pure and insulin loaded polysaccharide derivative hydrogels were further observed after 5 days culture, which possessed normal spherical morphologies. Cytocompatibility study showed the formed hydrogels were non-cytotoxic and preserved the viability of the entrapped cells. This hydrogel can be used as a potential delivery system as it 97


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resemble the ECM of tissues comprised of GAGs and can create a more biomimetic microenvironment with excellent biocompatibility.

Conclusion Hydrogels are promising substrates for regenerative medicine with the advantage that drugs and cells can be readily integrated into the gelling matrix. Many efforts have been developed to improve hydrogels and thus, support the development of more natural and functional tissues. In this chapter, we discussed the design of biodegradable and biocompatible hydrogel systems to be potentially used in regenerative medicine applications. These biodegradable hydrogels can serve for therapeutically effective platforms for both in vitro and in vivo applications. The success of tissue constructs is highly dependent on the design of the hydrogel scaffolds including physical, chemical and biological properties. An ideal hydrogel would potentially mimic many roles of ECM found in tissues, resulting in the coexistence of both physical and chemical gels. Hydrogel development will likely have a significant impact on the advancement of regenerative medicine. However, current hydrogel systems are unable to meet all the design parameters simultaneously (e.g., degradation, biocompatibility or mechanical properties). Novel crosslinking methods should be developed, both to enhance the material biocompatibility as well as control the mechanical properties. In addition, cell induction ligands such as growth factors and genes can be incorporated into hydrogels such that specific signals could be delivered in an appropriate spatial and temporal manner.

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References Liu X., Ma L., Mao Z., Gao C. (2011). Chitosan-based biomaterials for tissue repair and regeneration, Advances in Polymer Science, 244, 81–128. Langer R., Vacanti J.P. (1993). Tissue engineering, Science, 260, 920–926. Tan H., Wu J., Lao L., Gao C. (2009). Gelatin/chitosan/hyaluronan scaffold integrated with PLGA microspheres for cartilage tissue engineering, Acta Biomaterialia, 5, 328–337. Tan H., Gong Y., Lao L., Mao Z., Gao C. (2007). Gelatin/chitosan/hyaluronan ternary complex scaffold containing basic fibroblast growth factor for cartilage tissue engineering, Journal of Materials Science: Materials in Medicine, 18, 1961– 1968. Drury J.L., Mooney D.J. (2003). Hydrogels for tissue engineering: Scaffold design variables and applications, Biomaterials, 24, 4337–4351. Tememoff J.S., Mikos A.G. (2000). Injectable biodegradable materials for orthopedic tissue engineering, Biomaterials, 21, 2405–2412. Hou Q.P., De Bank P.A., Shakesheff K.M. (2004). Injectable scaffolds for tissue regeneration, Journal of Materials Chemistry, 14, 1915–1923. Nuttelman C.R., Rice M.A., Rydholm A.E., Salinas C.N., Shah D.N., Anseth K.S. (2008). Macromolecular monomers for the synthesis of hydrogel niches and their application in cell encapsulation and tissue engineering, Progress in Polymer Science, 33, 167–179. Tan H., Marra K.G. (2010). Injectable, biodegradable hydrogels for tissue engineering applications, Materials, 3, 1746–1767. Tan H., Ramirez C.M., Miljkovic N., Li H., Rubin J.P.; Marra K.G. (2009). Thermosensitive injectable hyaluronic acid hydrogel for adipose tissue engineering, Biomaterials, 30, 6844–6853. Tan H., Luan H., Hu Y., Hu X. (2013). Covalently crosslinked chitosanpoly(ethylene glycol) hybrid hydrogels to deliver insulin for adipose-derived stem cells encapsulation, Macromolecular Research, DOI: 10.1007/s13233-013-10238. 99


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Tan H., Shen Q., Jia X., Yuan Z., Xiong D. (2012). Injectable nano-hybrid scaffold

for

biopharmaceuticals

delivery

and

soft

tissue

engineering,

Macromolecular Rapid Communications, 33, 2015–2022. Tan H., Wan L., Wu J., Gao C. (2008). Microscale control over collagen gradient on poly(L-lactide) membrane surface for manipulating chondrocyte distribution, Colloids and Surfaces B: Biointerfaces, 67, 210–215. Tan H., Lao L., Wu J., Gong Y., Gao C. (2008). Biomimetic modification of chitosan with covalently grafted lactose and blended heparin for improvement of in vitro cellular interaction, Polymers for Advanced Technologies, 19, 15–23. Tan H., Wu J., Huang D., Gao C. (2010). The design of biodegradable microcarriers for induced cell aggregation, Macromolecular Bioscience, 10, 156– 163. Tan H., DeFail A.J., Rubin J.P., Chu C.R., Marra K.G. (2009). Novel multi-arm PEG-based hydrogels for tissue engineering, Journal of Biomedical Materials Research, 92A, 979–987. Tan H., Huang D., Lao L., Gao C. (2009). RGD modified PLGA/gelatin microspheres as microcarriers for chondrocyte delivery, Journal of Biomedical Materials Research, 91B, 228–238. Tan H., Li H., Rubin J.P., Marra K.G. (2011). Controlled gelation and degradation rates of injectable hyaluronic acid-based hydrogels through a double crosslinking strategy, Journal of Tissue Engineering and Regenerative Medicine, 5, 790–797. Tan H., Zhou Q., Qi H., Zhu D., Ma X., Xiong D. (2012). Heparin interacting protein mediated assembly of nano-fibrous hydrogel scaffolds for guided stem cell differentiation, Macromolecular Bioscience, 12, 621–627. Tan H., Xiao C., Sun J., Xiong D., Hu X. (2012). Biological self-assembly of injectable hydrogel as cell scaffold via specific nucleobase pairing, Chemical Communications, 48, 10289–10291. Tan H., Chu C.R., Payne K.A., Marra K.G. (2009). Injectable in situ forming biodegradable chitosan-hyaluronic acid based hydrogels for cartilage tissue engineering, Biomaterials, 30, 2499–2506. 100


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Sun J., Xiao C., Tan H., Hu X. (2013). Covalently crosslinked hyaluronic acidchitosan hydrogel containing dexamethasone as an injectable scaffold for soft tissue engineering, Journal of Applied Polymer Science, DOI: 10.1002/app.38779. Tan H., Rubin J.P., Marra K.G. (2010). Injectable in situ forming biodegradable chitosan-hyaluronic acid based hydrogels for adipose tissue regeneration, Organogenesis, 6, 173â&#x20AC;&#x201C;180. Tan H., Hu X. (2012). Injectable in situ forming glucose-responsive dextran-based hydrogels to deliver adipogenic factor for adipose tissue engineering, Journal of Applied Polymer Science, 126, E180-E187. Tan H., Rubin J.P., Marra K.G. (2011). Direct synthesis of biodegradable polysaccharide derivative hydrogels through aqueous Diels-Alder chemistry, Macromolecular Rapid Communications, 32, 905â&#x20AC;&#x201C;911.

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4 Four-element Principle of Organic/Polymer pi-Semiconductors Ling-Hai XIE a, Wei Huanga, b a Center for Molecular Systems and Organic Devices (CMSOD), Key Laboratory for Organic Electronics & Information Displays (KLOEID) and Institute of Advanced Materials (IAM), Nanjing University of Posts and Telecommunications (NUPT), Nanjing 210046, China b Jiangsu-Singapore Joint Research Center for Organic/Bio- Electronics & Information Displays and Institute of Advanced Materials, Nanjing University of Technology, Nanjing 211816, China E-mail :

O

iamlhxie@njupt.edu.cn

iamwhuang@njupt.edu.cn

rganic/polymers semiconductors are pi-systems that are categorized into πconjugated organics and π-stacked organics. From the bottom-up view of organic devices, we here demonstrate the four-element principle of organic/polymer π-semiconductors according to our research results using fluorene-based oligomers/ polymers and stacked polymers as typical models. We highlighted the role of four nodes at the molecular-scale that include electronic structures, steric hindrance, conformation and topology, as well as supramolecular bonds (or called as noncovalent interactions) on the thin-films and organic devices. The four-element chemical blueprints that have been described by us are complementary to the energy diagram and morphology regimes in organic devices, allowing for more deeply understanding relationships between molecule systems and organic devices. Four-element analysis and designs could be further extracted into the matter-energy-information-awareness scenario that is one creative epistemology for not only science, but also philosophy, technology, economy and social society.

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Introduction Plastic electronics boost several high-tech industries areas that include flat panel displays, solid light sources, sustainable energy, semiconductors, and RFID and others (Xie, 2012). Polymer semiconductor-based organic thin-film devices afford a series of advanced alternatives to these areas by means of the polymer features of mechanical, flexible, large area, low-cost fabrication (Burroughes, 1990; Forrest, 2004). Until now, organic devices that have been extensively investigated include polymer light-emitting devices (PLEDs), polymer solar cells (PSCs), transistors, lasers, memories, photodectors, sensors and actuators. This area face challenges not only from the unprecedented exploration of functions such as spintronics, terahertz (THz), quantum cellular automata (QCA) and mechatronics as well as robotics, but also from the commercialization of organic devices with merits of both high-performance and low-cost as well as long lifetimes that strongly depend on the progress in the complex relationships between molecular structures and device parameters.

Figure 1. (top) the trans-polyacetylene, (bottom) two basic p-type and n-type π-segments that are combined by two manners of type I: π-conjugated configuration and type II: π- stacked configuration.

Polyacetylene is the basic models of organic/polymer semiconductors (Figure 1) (Chiang, 1977; Shirakawa, 1977). From the chemical structure point of view, organic/ polymer semiconductors are the transfer and transport π-channels for 104


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charge and carriers such as electrons, holes, polarons, bipolarons and so on. The essence of organic/ polymer semiconductors is the π-systems where π-orbital make it possibility to delivery electrons and to generate the electrical currents. Basically, on the one hand, the π-segments have been divided into two basic types that are n-type and p-type according to the electronic features (Figure 1). To design the idea fundamental n-type or p-type organic/polymer semiconductors are of crucial importance to improve the device performances. On the other hand, πorbitals can be arranged into two extreme manners that are parallel π-π conjugation style via covalent bonds and head-to-head π-π stacking conjugation style via noncovalent interactions (Figure 1). As a result, polymer πsemiconductors are categorized into π-conjugated polymers and π-stacked polymers (Xie, 2012). One way to create new-concept organic/polymer semiconductors is to make the above four kinds of π molecular segments, including p-type, n-type, π-conjugated, and π-stacked segments, transform into new elementary π-segments and to organize them into high-level complex structures by means of covalent bonds or noncovalent interactions. We demonstrated a π-stacked and conjugated system based on poly(N-vinylcarbazole) postfunctionalized with terfluorene for stable deep-blue hole-transporting materials and multifunctional white host materials (Xie, 2009). On the other hand, it can be predicted that δ-semiconductors that consists of low-dimensional polymetal chains at the molecular scale will be another key trend for semiconductor science and technology beyond organic π-semiconductors.

Figure 2. The four-element scenarios in organic devices, including device performances, energy diagram, morphology regime, and molecular structures.

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Organic devices could be commercialized by virtue of the intriguing function, excellent performance, good stability, low-cost fabrication procedures and other individual quality. Organic semiconductor science have task and mission to uncover the deep relationships for the guideline of the technology, industries, economy and societies in organic/ plastic electronics. The self-contained four nodes are indispensable to the complete systems in organic electronics (Figure 2). Firstly, the prerequisite condition for this area is to fabricate and evaluate the prototype organic devices such as organic light-emitting devices, polymer solar cells, organic lasers, nonvolatile organic memory, and others. Secondly, the energy diagrams that are directive to the performance of organic electronic devices is described by the series of parameters such as work functions of electrodes, femi energy levels, the HOMO and LUOMO of active layers, energy levels of dopants and so on. However, energy diagram is not enough informative to fully understand the organic devices. Thirdly, organic devices are determined by the organic semiconducting thin films. The morphology regime is another important blueprint to organic devices, especially stability and lifetime, which is strongly impacted by fabrication procedures. Organic semiconductors are molecular semiconductors that have intriguing supramolecular interfaces between molecules. Polymorphism makes the molecule-device relations much more complex. Organic semiconducting thin films involve not only supramolecular and molecular scales, but also meso-scale and nanoscale superstructures. Hierarchical features make the complicated relationships between organic semiconductors and organic devices, meanwhile complex organic semiconductors become the trends in the future. Finally, organic devices have long-range correlation with molecular structures via the energy diagram and morphology regimes. Covalent structures, especially at the molecular scale, are of ultimate importance to improve the performance of organic devices. It is very necessary to set up the chemical scenario that benefit for the guideline of molecular design to challenge the thorny issues. The chemical scenario are emergent to be set up for the chemists, especially organic chemists entry into this areas. The basic question is which 106


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molecular segments should be screened into the organic/ polymers molecules that strongly depend on the knowledge from the experiment observations. These challenges have been discussed extensively in the literature. One example is that specific electronic structure groups may be selected according to doping or blending host-guest experiments (Lin, 2011; Liu, 2010). However, the transformation from observation and literatures into molecular design is not easily that suffer from a screening process, that make a reservoir with many data that can be set up into inter-relative social networks by logic analysis, terminally to form the principle to guide the molecular designs. We have proposed four-element principle that is extracted from a lot of practices by the Xie-Huang group by means of the self-similar epistemology on the template of the discipline of physical organic chemistry (Xie, 2010). Four nodes include electronic structures, steric hindrance, conformation and topology, as well as supramolecular interactions (or called as noncovalent bonds) (Figure 3). The four element principles of molecular design are favorable for guiding chemists how to organize or use the molecular engineering tools. Four element nodes at molecular scale make it possible to set up systemically chemical scenario that will bridge between physics and chemists. In this chapter, four-element scenario and molecular design of organic/polymer semiconductors in the framework of the four-element principle will be discussed and demonstrated.

Figure 3. Four-element chemical scenario in organic semiconductors. Reprinted with permission from ref (Xie, 2010). Copyright 2010, Editorial office of Acta Physico-Chimica Sinica.

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Molecular designs are a part of molecular engineering that includes two conceptual levels. One is the isolation process of chemical systems that are boundary from the ever continuous phase, such as nanoparticles and metal complexes; another is the integration and networks of organic molecules via covalent engineering process to produce the new functionalities or external outstanding properties. Organic synthesis affords flexible tools to control chargecarriers behavior through the combination and reorganization of Ď&#x20AC;-orbitals, resulting in dramatically different optoelectronic properties as well as morphologies for semiconducting polymer films applied in electronic devices. Molecular design of Ď&#x20AC;-conjugated polymers at the molecular scale that make powerful tools to access stable and high-performance polymer devices over any other methods for organic devices is not readily feasible for inorganic semiconductors. From top-down point of view, molecular designs can be directed by device functionality, performances, energy diagram and morphology regimes such as assembling building blocks and superstructures. As a result, several key ways can be established such as function-directed molecular design, performancedirected molecular design, electronic structure-directed molecular design, morphology-directed molecular design, and superstructure-directed molecular design. On the other hand, in the framework of the self-contained four-element devices, from bottom-up point of view, top-down architectures are dominated by four-element designs that will be detailedly discussed as follows.

Figure 4. Energy diagram levels of organic diode devices.

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1-1- Electronic structure analysis and design Device functions and performances make the first driving force to design electronic

structure

in

organic/polymer

semiconductors.

For

organic

semiconductors, electronic structures of organic semiconductors are the first key element to dominate device behaviors from the physical energy diagram of organic devices. Electronic structures such as the Highest Occupied Molecular Orbital (HOMO), Lowest Unoccupied Molecular Orbital (LUMO) and band gap at the molecular scale are close to the various electronic processes such as charge transfer and exciton behaviors such as energy transfer, dominating the device performances. The arrangement of n/p or D/A building blocks determine the electron and exciton behaviors such as injection, transport, transfer, dissociation, freedom road, binding energy, and many parameters at the physical levels. Therefore, the p-n chemical design is always the first tool to explore the new device functions and tune energy diagrams at the large scales. Frontier molecular orbital theory is suitable for not only organic reactions, but also organic optoelectronic materials. When electron-withdrawing molecular building blocks and electron-rich molecular building blocks were condensed into one molecule or polymer chains via the organic synthesis, their energy level reorganize and split to form new HOMO and LUMO energy levels. Huang group has proposed a p-n (DA) molecular design for polymer band-gap engineering (Yu, 1998; Huang, 1998; Huang, 1999). p-n molecular design that are also called as D-A approach can trace back to the molecular orbital theory in organic chemistry. The p-n designs are essential to organic devices that resemble to p-n junction design in inorganic devices.

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Figure 5. (top) p-n copolymerization and chemical structures of the modeling polymers. (bottom) Absorption and fluorescence spectra of polymers A-D as films and Cyclic voltammograms of polymers E-G. Reprinted with permission from ref (Yu, 1998). Copyright 1998, the American Chemical Society.

For the basic organic diodes, a typical energy diagram is shown in Figure 1. Electronic structure designs are based on the requirement of energy diagram in organic devices. Initially, for the polymer light-emitting devices (PLED), the basic problem was to realize the red-green-blue (RGB) three primary colors and to balance the injection between electron and hole. One effective strategy is p-n copolymerization strategy involves the alteration of the electronic behavior of Ď&#x20AC;conjugated polymers. In plastic electronics, we first demonstrated the p-n copolymerization to tune the color and balance between hole and charge of the carriers from the anode and cathode (Yu, 1998; Huang, 1998; Huang, 1999). The modeling conjugated copolymers A-G was shown in Figure 5, in which the electron-rich oligothiophenes acting as p-dope type blocks, while the Ď&#x20AC;-deficient di(1,3,4-oxadiazole) phenylenes acting as n-dope type blocks. The length of the pdopable segments in copolymers A-G can be altered by the thiophene ring number from 1 to 3. As a result, in the photoluminescence spectra, the emissive colors are blue, green, and orange for corresponding copolymers A-D with one, two, and 110


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three thiophene rings in the p-dopable segments (Figure 5). In cyclic voltammetry, for copolymer E-G, the oxidation peak potential is gradually reduced with increasing the thiophene ring number in the oligothiophene segments from 1 to 3 from 1.95 V for E, to 1.64 V for F, to 1.38 V for G. The oxidation potentials of copolymer G are close to those of hole-injection favorable polythiophenes. The pn molecular design affords the possibility of tuning redox behavior (related to the HOMO and LUMO of polymers) and emissive color of the resultant polymers.

Figure 6. Three basic energy level arrangements of p-type (donor) and n-type (acceptor) segments.

There are three basic types of energy level diagrams in p/n doping organic semiconductors, as shown in Figure 6. For type I, there is seldom overlap of energy level between donors and acceptors. In this case, ground-state electron transfer occurs. The bulk conductivity of conducting polymers may be enhanced by chemical doping. Using this technique, hole or electron carrier injection have been effectively enhanced by either and molybdenum p-dopants such as tetrafluorotetra cyanoquinodimethane or as n-dopants such as cobaltocene. For type II, there is a partial overlap of energy levels between donors and acceptors. In this case, the p-n molecular design causes the large intrachain charge transporting and Frรถster energy transfer features. Hybridization generates the new HOMO and LUMO with a lower band gap. Mesomerism leads to the improved quinoid states or de-aromaticity, thereby facilitating charge transfer. As a result, these types of D-A conjugated copolymers exhibit a strong ground-state dipole due to intramolecular charge transfer. Photo-induced electron transfer is readily observed in the excited state. Charge transfer (CT) on the polymer chains could reduce the 111


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probability of excited-state trapping at unstable sites, facilitating the device stability. The low band gap Ď&#x20AC;-conjugated copolymers of type II are favorable for ambipolar charge-transport behavior. The type of memory devices are determined by the depth of trap that can be tuned by CT states. CT complexes impart the wide light absorption of donor-acceptor copolymers that can improve the efficiency of polymer solar cells. For type III, there is an inclusion relation between donors and acceptors. When the energy structure with type III occurs between molecular segments, it has been proven that polymer semiconductors facilitate FRET more easily than charge transfer. Copolymers of the type III are very useful in designing red and white light-emitting polymers (LEPs). Exciton confinement could improve the external quantum efficiency (EQE) of PLEDs, in which efficient energy transfer and exciton trapping occur from segments with wider energy gaps to segments with narrow band gaps.

Figure 7. Polyfluorene-based semiconductors combined with various periodic table elements for organic electronics. Reprinted with permission from Ref (Xie, 2012). Copyright 2013, Elsevier.

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The powerful p-n molecular designs to tailor the electronic structures make polyfluorenes entry into various organic devices such as polymer light-emitting devices, polymer lasers, polymer solar cells, memory cells, sensors and photodetectors. Electron-donating groups and electron-withdrawing groups have been introduced into polymer films through substitution, copolymerization, endcapping and other techniques. The introduction of atoms with an electronegativity different from carbon and their molecular segments with large polar effects provides an effective method to tune the optoelectronic properties of semiconducting polymer films. The introduction of periodic table elements into polyfluorenes makes the large change of bandgaps due to the different electronegativities and dipoles of molecular segments (Figure 7). Controllable energy transfer make a chance to RGB primary color and white emission. Intrachain energy-transfer polyfluorenes are promising red lightemitting materials with high-efficiency. Mutli-component copolyfluorenes are potentially commercialized white-lightemitting materials.

1-2- Steric hindrance analysis and design Steric hindrance enable finely tune optoelectronic properties and morphology of organic/polymer semiconductors by actuating three other elements, including electronic structures, conformations, and noncovalent interactions, in turn providing a unique tool to improve device performance and stability. Steric hindrances can occur at the intrachain and interchain in polymer semiconductors. The intrachain steric configuration can actuate dihedral angles between the adjacent monomer units. Its set effect is the basis to the conformation effects that will mention in the next section, resulting in small change of electronic structures. On the other hand, steric hindrances can increase the interchain distance, reduce the packing capability, and suppress aggregation. In contrast with electronic 113


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structure design, most of the sterically hindered groups do not alter the HOMO and LUMO energy level significantly when it is introduced into Ď&#x20AC;-conjugated polymer chains. The softest steric hindrances come from pendant groups of various alkyl side chains that act as the spacer role to segregate the conjugated rigid chains. The remote steric hindrance is very ideal steric hindrance models without change of dihedral angle of segment monomers in polyfluorenes. Bulky groups are favorable for the improvements of glassy transition temperature (Tg) and color stability in amorphous light-emitting polymers. Steric hindrance design will make new bulky groups into organic/polymer semiconductors to understand the relations between molecular structures and devices and to improve device entity. Aromatics containing sp3 carbon are state-of-the-art bulky groups, such as spirofluorenes, arylfluorenes, and others such as iptycenes. These bulky pendant groups have been demonstrated to be effective morphological stabilizers for OLED or PLEDs (Hou, 2009). Steric hindrance designs have been also reported by other groups. Dendron substitution is an extreme example of steric hindrance that exploits the self-shield and self-packed effect of polymer chains. However, the large building blocks into organic semiconductor make the longer distance among Ď&#x20AC;-segments, and disorder transport pathway, resulting in damaging device performance such as higher turn-on voltage and so on. Therefore, it should take care to choice the suitable bulky groups that do not damage the device performance, but improve the morphological stability.

Figure 8. Chemical structures of PFO and PSBFs,together with varous spirocyclic aromatic hydrocarbons (spiro-arenes) building blocks such as SBF, SFX, SFDBX, SFDBA and SFDBAO.

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Xie-Huang group carried out implementations of steric hindrance design to demonstrate two outstanding sterically hindered groups that are spirofluorenes in conjugated polymers for PLEDs (Yu, 2000) and arylfluorenyl moieties in stacked polymers for nonvolatile flash memory (Xie, 2008). For PLEDs, the device stability is the issues for the commercialization. One key point is the color and spectral stability that includes the color purity and efficiency keeping during the long-term operations. Polyfluorenes as the most potential blue light-emitting materials were hindered by the issues of low-energy green emission bands that are also called as g-bands. The origin that is still debating stemmed from either aggregates/excimers or ketone defects. In this situation, we proposed the spirofuctionalized conjugated polymers to improve morphological stability by synthesis through Suzuki coupling reaction (Yu, 2000). This design of highly steric spiro-structure with dihedral angle of 90ď&#x201A;° between two fluorene units can separation the interchain interaction, suppress aggregation and crystallization, while maintain the backbone structure of polyfluorenes so as to keep the electronic features with the same bandgap and emissive peak. Amorphous polymer thin films with high glassy transition temperature lower the tendency of crystallization. PSBFs with the R = C6H13, are only soluble in hot solvent; R = C8H17 (PDOPDSF), Mw > 30, 000, completely soluble in chlorobenzene. PSBFs gave optically high quality films with narrower emissive spectra and improved thermal stability. As result, PSBFs have been high Tg and less the g-band than PFO under the same annealing condition. Along with this design, we explore a series of spirofluorene molecular segments via concise one-pot routes, including SFX (Xie, 2006), SFDBX (Liu, 2009), SFDBA, and SFDBAO (Lin, 2012) (Figure 8). Polyfluorenes bearing with these spirocyclic aromatic hydrocarbons (SAH) (Xie, 2010) will be the start-of-the-art models for the investigation of structure-performance relationships.

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Figure 9. Chemical structures of hindrance-functionized stacked polymers, including PVK-PF, PPFVK, and PPFS, together with the proposed conformational change of shuttle motion for nonvolatile flash memory. Reprinted with permission from ref (Xie, 2008). Copyright 2008, the American Chemical Society.

In comparison with conjugated polymers Ď&#x20AC;-stacked polymers have wider bandgap than dynamic conformations, which make a chance to organic devices. PVK and PS are typical Ď&#x20AC;-stacked polymers. Beyond the hole transporting materials by PVK, the PVK can act as the electro phosphorescent host materials for blue color PLEDs owing to the high triplet energy levels. The stability and electrontransporting limitation are the issues. The supramolecular characteristics of stacked polymers with the dynamitic conformations are one entry to polymer nonvolatile memory devices. The controllable conformations should be made it in order to improve the low-state stability in flash memory. Facing on the issues, we proposed hindrance functionalization of supramolecular polymer semiconductors (SPSs) (Xie, 2008). Phenylfluorenyl moieties (PFMs) have a sp3 carbon to link the active conjugated polymer or organic semiconductors, so in general do not play a role in changing electronic structures such as the typical HOMO and LUMO. Photophysical and electrochemical studies show that PVPFK has the ET level of 2.80 eV with the better electron-transporting ability than PVK. Proof-ofconcept PPLEDs have been fabricated using PVPFK as host and FIrpic as blue guest. (Yin, 2011). Our device exhibited higher luminance efficiency than PVK device in the whole measuring voltages range from 0 to 21 V. The maximum 116


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luminance efficiencies were 14.2 and 13.3 cd/A for PVPFK and PVK device, respectively, and the maximum power efficiencies were 2.8 and 2.4 lm/W for PVPFK and PVK device. Recently, we use the approach in polystyrenes to construct PPFS for polymer electroluminescent hosts and nonvolatile memories (Yin, 2013). Dual terminal memory devices using PPFS as active layers exhibited nonvolatile flash memory with ON/OFF current ratio of up to 104 under a read voltage of 1.5 V, which are dramatically different from insulate PS without any memory performance. Polymer light-emitting devices (PLEDs) using PPFS as host and FIrpic as guest exhibited blue emission with CIE chromaticity coordinates (x, y) of (0.146, 0.349), maximum luminance of 793 cd/m2, maximum luminance efficiency of 3.0 cd/A and the maximum power efficiency of 0.54 lm/W. Hindrance-functionalization open a way to transform functional Ď&#x20AC;-stacked macromolecules

into

advanced

polymer

semiconductors.

SAH

and

phenylfluorenyl moieties (PFMs) are two fundamental bulky groups with large steric hindrance for organic electronics.

1-3- Conformation/topology analysis and design Beyond steric hindrance, conformation and topology analysis is another key node for the molecular design of organic/polymer semiconductors. On the one hand, they dominate the intrachain carrier-transport channels, on the other hand are closely related with the morphology regimes that involves chain packing, textile and others. Conjugated polymers have more rigid conformation than stacked polymers and other vinyl polymers. Conformation designs offer a chance not only to improve the morphological stability, but also to tune optoelectronic properties. Linear Ď&#x20AC;-conjugated polymers are regarded as one-dimensional carrier-transport channels with low-dimensional characteristics. In general, linear and planar conformations afford high charge mobilities in conjugated polymers. The planarizations of conformations support the electron delocalization that is determined by the effective conjugation length. Toward this direct, one molecular design is ladder-type Ď&#x20AC;-conjugated polymers that possess high rigid backbone and 117


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planar conformation with the feature of large Stokes shifts, benefiting the improvement of charge mobility than precursor conjugated polymers without fuse. Another conformational molecular design allows for the construction of zigzagged or kinked conformations by means of isomerization with shorter effective conjugation lengths in conjugated polymers that than those of their linear counterparts (Liu, 2010; Wang, 2011). The curve conformations will present many intriguing features to deeply get insight into the relations and afford new tuning methods. Rod-coil copolymers and rod-rod copolymers belong to conformation design. They have two different block segments that exhibit versatile supramolecular assembly behaviors. The aims to control conformational change are very severe challenges, but are very important to extend the functions and improve the performances of conjugated polymers. One appealing examples is the beta-conformation chain in polyfluorenes that have several advantages such as the stability for PLEDs, the higher luminescent efficiency and low laser threshold for organic lasers over the intrinsic alpha-polyfluorenes. Conformational change also afford new chance to fabricate the nanopatterned thin films and create the electrically bistable systems at the molecular or chain scale (Xie, 2008). In addition, it is important to seek very ideal models to understand the conformation effects on the devices. Beyond conformation control, topology is the spatial arrangement of active Ď&#x20AC;-molecular bricks that enable to direct complex organic semiconductors. In fact, an intrinsic topology design is from organic semiconductors to polymer semiconductors. Several topology molecular designs in polymer semiconductors, such as multi-armed macromolecules, hyperbranched polymers and dendrimers, have also been developed. Macrocycles make it possible to construct organic/polymer semiconductors more complex. One advanced topology design is to introduce catenane and rotaxane into organic/polymer backbones.

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Figure 10. The chemical structures of diarylfluorene-based tolophogy designs: (a) H shaped molecule, (b) multi-armed molecule, (c) chair-shaped molecule, (d) soluable covalent organic framework dendrimer, and (e) frame[n]arenes.

In the framework of morphology-directed molecular design in organic devices, our group makes conformation and topology design that focus on the diarylfluorene-based complex organic semiconductors. We first set up the FriedelCraft protocol of fluorenol as a tool to construct complex diarylfluorenes (Xie, 2006). This protocol has several features such as high yields, metal-free, room temperature, C-H activation and functionalization, and click-like style and so on. Friedel- Craft protocol offers a flexible and elegant platform for polymerization (Liu, 2011), post-functionalization, and others. One kind of π-conjugationinterrupted polymer that contains tetrahedral sp3 carbons on the polymer backbone have been designed and synthesized by Friedel- Craft protocol. The π-interrupted polymers exhibit an amorphous phase with higher Tg temperatures, high spectral and thermal stability that is well suited for applications of wide-bandgap violetblue materials or electro-phosphorescent host materials. Hyperbranched πconjugation-interrupted polymers have been modified by the large π-aromatics with high mobility, which give obviously low turn on voltages in PLEDs (Liu, 2009). Ultimately, topology designs make new-generation organic framework semiconductors and organic nanomolecular semiconductors that have been 119


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proposed by us

(Zhang, 2013). Soluble covalent organic frameworks can

integrate multiple types of π-segments into one entity that lead to dramatically distinguishing properties. Frame molecules possess many porous superstructures that are suitable for applications in either biosensors or the detection of explosives. We designed asymmetric chair-shaped opening frames toward multicomponent organic semiconductors (Figure 10c). Framework design have been achieved by the Friedel-Crafts protocol using the starting H-shaped structures (Xie, 2006), three-dimensional nano-sized multi-armed structures (Figure 10c) (Zhao, 2008) and three-dimensional dendrimers (Figure 10d). We are challenging the frame[n]arenes (Figure 10e) by the fluorenol and Friedel-Crafts protocols.

1-4- Supramolecular analysis and design The weaker and reversible noncovalent interactions always exist between molecules. The noncovalent interactions between polymer chains are more complex than small molecules. Supramolecular interactions and non-covalent bands that serve as molecular-scale interfaces determine the self-assembly behavior, molecular nanostructures, phase separations, and domain morphologies. For organic/ polymer semiconductors, polymer chains that consist of active πsegments enable to be reorganized into different morphology of polymer thin films through weak supramolecular interactions, in turn resulting in the alteration of the current-voltage characteristics of devices. In general, noncovalent interactions include π-π stacking interactions, hydrogen bonds, ionic bonds, van der Waals forces, and hydrophobic interactions. Supramolecular π-π stacking interactions are closely related with electron-hopping that play a unique role in Organic/ polymer semiconductors and electronics. The π-π stacking motifs of πsegments either edge-on or face-on arrangements on the surface dramatically influence the mobility of transistors. Supramolecular organic semiconductors have been proposed by us to explore the conformational change or switches of 120


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intramolecular supramolecular systems such as π-stacked polymers for the resistive memory devices (Xie, 2008). Main-chain-type supramolecular polymers with dynamic linkages of noncovalent bonds enable to organize the active πsegments into the unique fantastic arrangements and supramolecular p-n heterojunctions. Conjugated polymers without any alkyl side chain have strong ππ interactions, resulting in the insolubility. Soluble polymer semiconductors are thanks to the van der Waals forces from alkyl side groups. Ionic bonds such as quaternary ammonium salts make conjugated polymers water soluble into conjugated polyelectrolytes with the potential applications such as biosensors and optoelectronic devices. Hydrogen bonds lead to the generation of conjugated polymer amphiphiles that are favorable for well-defined, hierarchically selfassembled architectures. Thus, supramolecular approaches that are of utmost importance offer big room and versatile toolbox to uncover the relation between molecules and devices, optimize the desirable device performance, and impart new functions. Ultimately, supramolecular nanoelectronics will be one intermediate field between molecular electronics and organic thin-film electronics.

Figure 11. The research route diagram of supramolecular conjugated polymers and thin-films.

It is emergent to explore supramolecular thin-film electronics from the view of device commercialization. One reason is the complex relationship between thin films without well-defined supramolecular interactions and devices. Furthermore, supramolecular analysis are favorable for clarifying how supramolecular 121


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interactions work in printing inks, thin films and devices. Morphology will become controllable in the concrete film-forming procedures by means of supramolecular

design.

Supramolecular

functionalization

of

polymer

semiconductors allow for the state-of-the-art models to understand the detailed mechanism of optoelectronic properties. In this context, we focus on a series of supramolecular design of polyfluorenes to understand the evolution of morphology from precursor structures to beta phase, to carry out tuning emission colors by coordinative interactions and hydrogen bonds, to clarify the origin of debated g-band (low-energy green emission bands) on the blueprints (Figure 11). In order to obtain the ideal prototype models that rule out disturbing from other nodes, we designed diazafluorenes without disturbing configuration and backbones that have absolute advantages in terms of the deep investigations of the structure-function relationship over any other molecular segments due to the extreme similarity with fluorene except for two carbon atoms that are replaced by two nitrogen atoms. diazafluorenes as the model have been synthesized by our group (Zhao, 2011, Li, 2012). 4,5-diazafluorenes by the 2,7 postion linkage into polyfluorenes cause very small changes of steric hindrance and conformation, and limited change of electronic structure. In contrast, diazafluorenes impart polyfluorenes the coordination of metal ion, N…H-C weak interactions, and pH stimuli-responsive feature, making them become supramolecular polymer semiconductors with solvent-dependent color and emissions in solutions (Li, 2013). In comparison to poly(9,9-dioctylfluoren-2,7-yl) (PFO), the copolymers in the thin films cannot form beta phase with a configuration of a “planar zigzag” or 21 helix conformations that indicate that driving force from octyl side chains face the competition with N…H-C weak interactions. Copolymer PDAF8-co-F8 can form the organogels but require different solvents. Diazafluorene-based supramolecular conjugated polymers are potential candidates for organic thin-film devices. In order to set up the morphology scenarios and to clarify polyphorphism and g-bands of PFO, fluorenol have been selected by us as a unique hydrogen bond motif that is one of most investigated supramolecular interactions. We 122


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design polyfluorenol as molecular models of supramolecular conjugated polymers (SCPs) to investigate the external conditions and molecular weight effects on morphology, gels, and nanostructures (Lin, 2013). Hydrogen bonded PPFOHs have the big effects of solvent and molecular weight (Mn) on morphology and optoelectronic properties of supramolecular thin films. Solvents have been categorized according to whether organogels can form or not. Organogels afford robust criterion for the choice of solvents to spin coat thin films with different aggregates. Color-tunable supramolecular thin films have been achieved by the solvents and other external conditions for supramolecular light-emitting devices.

Figure 12. The models of supramolecular steric hindrance from the observation of the single crystals of spirofluorenethiophene-S,S-dioxides. Reprinted with permission from ref (Xie, 2007). Copyright 2007, the American Chemical Society.

It should be noted that the four-element principle is an ideal model. In fact, they interplay with each other and together to make the change of organic/polymer semiconductors. When one periodic table atom is incorporated into polymer semiconductors, it will alter not only the electronic structures, but also steric hindrance, conformation, topology, and even supramolecular interactions. Selfconsistent interplay between four elements generates the complex principles that afford power approach to challenge the issue that exist in organic electronics. To combine the electronic structure design and steric hindrance design have been more useful design to resolve the material issues. For example, polyfluorenes with side and/ or pendant chains that serve as both electronic structural groups and 123


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steric hindered groups make successful design stable color-purified highefficiency blue LEPs with gradient electronic structures. The π-substituent groups are multifunctional to design sterically hindered donors or acceptors, which represents a powerful tool for electroluminescent materials design. White PLEDs with high PL efficiency have been achieved by integrating a series of pendent segments with blue-, green- and red light-emitting. Dendrimer core-shell systems are FRET scaffold by means of both electronic structure and conformation design that suppress luminescence-quenching processes and achieve high QE. Integration of electronic structures with supramolecular design belongs to such multiple designs that will create donor-acceptor supramolecular conjugated polymers for organic electronics. One creative strategy is supramolecular steric hindrance (SSH) that affords a new design strategy for organic semiconductors and materials. In our work, we observed SSH phenomena from the spiro-molecules that belong to the bulky groups but can form well-defined noncolavent weak interactions (Xie, 2007). Spirofluorenethiophene-S,S-dioxides have two types of dimeric arrangements within molecule crystals through π-π stacking of the antiparallel interactions of two fluorene planes and C-H…O/C-H…S hydrogen bonding of intermolecular interactions (Figure 12). This result inspired us to design the supramolecular hindrance motifs to balance the color of light emitted and current efficiency (Xie, 2010; Gu, 2011). SSH afford a tool to design intramolecular π-stacked molecules for supramolecular spintronics. Polyretaxane is another impressive example as a molecular wire with multilevel design of steric, topology and supramolecular feature. Encapsulation play roles in shield effects that are assembly by supramolecular interactions, resulting in the stability of polymer chains against thermal and/or photo-oxidation for the color stability of polyfluorenes.

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2- Summaries and Outlook To summarize, we extracted the four nodes that include electronic structure, steric hindrance, conformation and topology as well as supramolecular interaction by means of self-similar epistemology in order to establish the chemical scenario of organic devices. Attempts to clarify the basic relations between four-element molecular structures, energy diagrams, morphology regime, and device performances have been made according to our groupâ&#x20AC;&#x2122;s works and others. The four elements have been combined with supplementary elements to deeply and clearly in the framework of four-element analysis and design, we carried out the molecular designs of p-n copolymerization, hindrance-functionalization in conjugated polymers and stacked polymers, soluble covalent organic frameworks via the Friedel-Crafts protocols that established by us, and fluorenol-based supramolecular functionalization. The p-n conjugated copolymers have tunable RGB emissive color and redox electrochemical properties for PLEDs. Stacked polymers with the feature of conformational change were effectively tuned by the steric design and have been successfully explored to achieve the highperformance nonvolatile flash memory. Supramolecular conjugated polymers (SCPs) have been demonstrated with the very big solvent and molecular weight effect on morphology and optoelectronic properties. State-of-the-art SCP models support the opinion of aggregates to cause the g-bands in supramolecular thin films. Outlook: Four-element designs will afford series of models to uncover accurate and precise scenario for the origin of g-bands in polyfluorene based PLED. Fourelement scenario and principle will guide the organic chemists to select molecular segments into organic/polymer semiconductors for plastic electronics and their technology, engineering as well as industries. Molecular frame and network design make an opportunity to create new-generation organic framework semiconductors. Four-element polymer semiconductors will become advanced optoelectronic materials to achieve stable high-performance polymer lightemitting devices, polymer solar cells, thin film transistors, and polymer memories. 125


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On the one hand, evolution of four-element design direct us beyond organic/polymer molecular semiconductors to entry the realm of nanoscale and nanomolecular semiconductors. Soluble organic nanoscale frameworks and networks that open a door to organic frame chemistry and materials will update smart and awareness materials for control devices and robotics. A perspective challenge for four-element principle is to build up database by artificial intelligence techniques such as social networks. This chapter about four-element analysis and design can be further extracted into the four-element scenario of matter-energy-information-awareness that is a general law of the universe, which hopefully is interesting for not only scientists but also philosophers and educators.

Acknowledgement The project was supported by the National Key Basic Research Program of China (973) (2009CB930600), National Natural Science Foundation of China (21274064, 21144004, 60876010, 61177029, 20774043, 20704023, 20974046), The Program for New Century Excellent Talents in University (NCET-11-0992), Key Project of the Ministry of Education of China (707032, 208050), Natural Science Foundation of Jiangsu Province (SBK201122680, BK2011761, BK2008053, SJ209003). L. H. X. thanks Jiangsu Overseas Research &Training Program for University Prominent Young & Middle-aged Teachers and Presidents.

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Chapter

5 Integrated Silicon Photonics Applied for Optical Interconnects Zhou Fang, Ce Zhou Zhao Department of Electrical and Electronic Engineering, Xi’an Jiaotong-Liverpool University, No. 111 Ren’ai Road, Suzhou, Jiangsu, 215123, China. E-mail :

Zhou.Fang07@xjtlu.edu.cn

cezhou.zhao@xjtlu.edu.cn

A

lthough copper interconnects have made essential contribution to the continuation of Moore‟s Law, they now barely fill the increasing bandwidth requirement in current computing systems. Optical technology, however, is one potential approach to revolutionize short-reach interconnects. Silicon photonics has been regarded as one leading candidate technology because of its lowcost and CMOS-compatible fabrication processes. Based on this perspective, great progress has been made in silicon photonics these years, and the corresponding devices for commercial applications have gradually been development in industry as well. This chapter provides an overview of the progress in silicon photonics for optical interconnect purposes. Crucial breakthrough in industry will also be reviewed at the end.

Introduction Through the Moore‟s Law, which successfully forecasts that the number of transistors on integrated circuits doubles approximately every two years, the semiconductor industry has adopted an economic imperative for improving computing performance [1]. Currently, Moore‟s Law still continues after Intel developed Ivy Bridge microprocessors by employing the tri-gate (3D) transistors manufactured at 22 nm [2, 3]. The major limit comes from the inherent trade-offs 133


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between clock frequency scaling for higher computing performance and the resulted power dissipation. To address this problem, multicore processors have been introduced as one effective way to further improve the performance. Instead of using one large core, several smaller cores are replicated on the same chip and the performance gain can come from parallel code execution using multiple threads across each core [4]. Furthermore, by handling more programs in parallel, the latency of inter-processor communication can also be reduced. However, further exploring multicore technique also brings a new challenge about the

on-chip

interconnects.

Traditional

interconnects

are

based

on

a

copper/dielectric paradigm, and reduction in interconnect width is always expected to increase device density on chips. Although the copper-based interconnects are now widely applied because of its low cost, it has at least three inherent problems. The first one is the limitation on integration density. Electrical connection cannot go through one another without insulation. The second problem focuses on its power dissipation. When scaling the metal interconnects, the resistance of the wires becomes significantly large. As a result, interconnects currently draw far more power than the transistors themselves, and they produce more heat as well which is harmful for the reliability of the processors. Finally, metal-based interconnects also have their limitation on speed because of the inherent RC time delay, or specifically the capacitive rise time of the interconnects, which is also one fundamental source of lag in todayâ&#x20AC;&#x;s microelectronics. Therefore, ITRS foresees that copper interconnects cannot scale to meet the communication challenge, and will be supplanted by an alternative technology by 2020 [5]. It is now widely recognized that photonic interconnects can offer natural solutions for the bottleneck of high-speed interconnect on chips [4, 6, 7], since it is the only interconnect technology which has demonstrated terabit per second single channel throughout [7]. One simple configuration of photonic interconnects is described in Figure 1.

The lasers on the left-hand side are used to supply the light at different

wavelength in different channels, acting as the â&#x20AC;&#x153;carriersâ&#x20AC;?. Electronic signals from 134


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the outside are added to the carriers by using the modulators to change the properties of the carriers. The optical signals are then combined together through the multiplexers, and transport through one single waveguide before being demultiplexed at the receiving end. The electronic signals are finally extracted from the optical signals by using the detectors. The basic principle is similar to the one used in wireless communication.

Figure 1 The scheme of the basic configuration of photonic interconnects

Instead of the RC time constant, the Nyquist frequency becomes the limit bandwidth factor in the optical interconnects descried above. Fortunately, the optical frequency is of the order of several hundred THz, so the Nyquist limit is of no concern from the foreseeable future. Furthermore, unlike electrons, photons have no rest mass and charge, which means that they can travel at velocity of light without the interference with electromagnetic field. This provides the interconnect network an enormous bandwidth increase, immunity to electromagnetic noise, low power consumption and reduced immunity to temperature variations. In addition, wavelength division multiplexing allows numbers of information channels propagating through one single waveguide, which provides the potential for high-density integration. 135


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Among a variety of optical solutions, silicon photonics links are especially attractive for several reasons: (1) Si provides a wide band infrared transparency from the wavelength of 1100 nm to 7000 nm approximately, which is far from being limited to the near-infrared (IR) communication band of 1300-1550 nm; (2) as silicon is considered as the workhorse of semiconductors industry, integration of photonic and electronic devices requires high compatibility to reduce the cost. On this aspect, silicon photonics naturally offers an excellent platform that provides this

compatibility with

standard

complementary metal

oxide

semiconductor (CMOS) fabrication process; (3) some excellent optical properties, like large optical damage threshold and thermal conductivity, also appear in Si. All of these reasons select Si as a remarkable candidate for photonics. This chapter aims to evaluate the potential impacts of silicon photonics in semiconductor industry. Previous reviews on silicon photonics can be found in [8â&#x20AC;&#x201C;10]. Here, after providing a brief introduction of the history, we will mainly focus on the application of silicon photonics for optical interconnect purposes. The latest progress in different building blocks, including passive devices, modulators, detectors and light sources will be reviewed in detail. In the end, some outstanding breakthroughs in commercialized devices developed in industry will also be discussed.

1- History of Silicon Photonics The origins of silicon photonics can be tracked back to 1986 when Soref and Larenzo published the pioneer work about the properties of silicon-on-insulator (SOI) structure for light confinement in near infrared [11]. The work also pointed out the advantages of optics when integrated with electronics. One year later, the second paper on exploring the electro-optical effect in silicon, which evaluated the possibility to change the refractive index of in silicon waveguide through plasma dispersion, was published, and this actually evaluated the possibility to realize modulation, switching and routing on chips optically without optoelectronic conversion [12]. The early work then stimulated activities that resulted in 136


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substantial progress to build the first generation of silicon photonics prior to 2000, mostly in passive devices [13–18]. At that time, the primary application was viewed to lay in telecommunications, serving as wavelength division multiplexing or optical switching in fiber-optic network. At the end of 20 century, silicon-based waveguides had been shrunk in size from more than 100 μm2 typical waveguides in 1980s to 5 μm2 size of rib waveguides by taking the advantage of the large refractive index contrast between Si (n = 3.45) and SiO2 (n = 1.45) [19]. Meanwhile, the emerging of techniques for hereogrowth of germanium on silicon also allowed the development of high-speed CMOS compatible optical detectors [20]. A review about the history of silicon photonics in 1980s and 1990s can be found in [21]. At the beginning of 21st century, silicon photonics started to boom and many important breakthroughs have been obtained. Luminescence was observed in Si using quantum confinement [22–24]. There are also some progress in modulators [25] and microcavities [26]. Key evolution, however, actually happens after 2004. In 2004, the first International Conference on Group IV Photonics (GFP), which is also the first ongoing global meeting solely to silicon photonic, was convened initially in Hong Kong. The year also witnessed to the ramping up of the investment from industry and government on silicon photonics. For example, the European Union started its Silicon Heterostructure Intersubband Emitter (SHINE) program to develop SiGe/Si quantum-cascade structures emitting at a wavelength in the range of 8-120 μm [27]. Defense Advanced Research Projects Agency (DARPA) also launched its four-year project to develop optoelectronic integrated circuits (OEIC) based on 1550 nm silicon photonics with on electronic and photonic integrated circuits (EPIC) in silicon [27]. In addition, the year 2004 also saw the publication of textbooks which specifically focused on silicon photonics for the first time [28, 29]. As for the research progress, crucial components, including, wavelength converter in silicon using Raman induced four-wave mixing (FWM) [30], all-optical switches [31, 32] and the first silicon-based Raman laser [33], were all realized by UCLA in 2004. Liu and Paniccia in Intel 137


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announced an optical device fabricated wholly in silicon using CMOS-compatible process, which can modulate the embedded optical signals at the speed of over 1 GHz [34]. In the same year, IBM also reported its low-lose silicon waveguide [35] and a 29 GHz SiGe photodetector [36]. These achievements gave the demonstration about the potential of silicon photonics in reducing the excessive power dissipation in microelectronic circuits. Therefore, 2004 is usually regarded as a turning point where silicon photonics stepped to its second generation. Afterwards, the development of silicon photonics for optical interconnect purpose was moving forwards step by step. In 2005, the modulation speed of silicon modulators was improved to 10 Gb/s individually by Intel [37] and Luxtera [38]. Simultaneously, Cornell University also reported a silicon modulator based on a ring resonator with a modulation speed of 1.5 Gb/s [39]. The bandwidth of SiGe photodetectors were also extended to 39 GHz [40]. In addition, the first continuous-wave (CW) Si Raman laser was reported by Intel in the same year [41]. In 2006, linear electro-optic (Pockels) effect was induced to silicon by applying strain to the crystal structure [42]. A hybrid AlGaInAs-silicon evanescent laser which can be electrically pumped was also invented by the UCSB and Intel [43]. Meanwhile, Foster and his colleagues from Cornell reported a broadband amplifier based on Raman gain [44]. Besides, multistage high-order microringresonator filters were intensively investigated by MIT, Intel and IBM as an effective method to achieve extremely high extinction ratio. The year 2007 was a fruitful period for the development of silicon photonics. In this year, Intel achieved a new milestone by developing a 40 Gb/s silicon optical modulator. Instead of using carrier injection, the operation of this modulator was based on carrier depletion of a p-n diode embedded inside a silicon-on-insulator waveguide [45, 46]. In the same year, Intel also developed a group of active devices operating at 40 Gb/s, including: (1) an electrically pumped mode-locked silicon evanescent laser on silicon for pulse generation [47, 48], (2) a fast SiGe waveguide-based photodetector [49], (3) an on-chip Raman amplifier [50], and (4) 138


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a high-speed wavelength converter based on four-wave mixing [51]. All of these breakthroughs from Intel provided solid evidence on the feasibility of silicon photonics for high-speed optical interconnects. In addition to Intel‟s achievement, Luxtera launched its world's first four-channel 10 Gb/s monolithic optical receiver using Ge waveguide photodetectors monolithically integrated in the 130 nm CMOS process [52], while IBM demonstrated an 10 bits optical buffer operating at 20 Gb/s by using 100 cascaded ring resonators [53]. Since most devices in silicon photonics were achieved with excellent performance in the year 2007, one of the priorities in the development of silicon photonics then shifted to the lowcost integration of these photonic devices together onto the CMOS electronic chips. In 2008, DARPA launched the Ultra-performance Nanophotonic Intrachip Communication (UNIC) project, aiming to achieve the optical interconnect with high performance for highly compact supercomputer systems [54]. Intel demonstrated the world‟s first cascaded silicon Raman laser [55] and a Ge/Si avalanche photodiodes (APD) with 340 GHz gain–bandwidth product [56]. For Raman lasers, the output beam is always at a longer wavelength than the pump. Therefore, one Raman laser can itself act as a pump to generate Raman lasing at an even longer wavelength. This provides an effective way to generate the beam with long-wavelength ranges. In the same year, Kotura developed an UltraVOA array as the first example of a successful silicon photonics-based product, which provides simple current-controlled optical attenuation (0-40 dB) and enables ultrafast (300 ns) power management in optical networks [19]. In 2009, Koos and his colleagues achieved a record high silicon photonic optical signal processing by developing a 4-mm-long silicon–organic hybrid waveguide with a record nonlinearity coefficient of γ = 105 W-1km-1. By embedding a thirdorder nonlinear optical polymer in the nano-slot of a silicon strip waveguide, it allowed the demultiplexing of a 170 Gb/s data signal at λ = 1556 nm, converting it to 42.7 Gb/s streams at λ = 1540 nm [57]. Jalali from UCLA also achieved a milestone work on non-linear optical properties of silicon based on their 139


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observation of coherent anti-Stokes Raman scattering in silicon [58]. Slow light was also a hot topic in this year as a new method to enhance the third-order nonlinear optical response of the bulk crystal silicon [59–61]. Example was achieved by employing 2D photonic crystal built in a silicon membrane as a vehicle for the generation of green light [61]. In addition, to overcome the poor emitting efficiency in Si, people also started to investigated the lasing in Ge in this year [62–66], and 1590 nm multi-line lasing using 1060 nm pumping of an Ge-onSi channel-waveguide Fabry-Pérot resonator was realized [64]. In 2010, Intel achieved a milestone work for silicon-based optical interconnects by building a 50 Gb/s silicon photonics link using silicon manufacturing methodologies [67]. Assefa from IBM demonstrated an effective method to overcome the large amplification noise while keeping a high signal gain [68]. By applying strong non-uniform electrical field to suppress the region of impact ionization, a dramatic reduction of amplification noise by over 70% was realized. Non-linear effects in silicon was widely regarded as a hot topic in silicon photonics in this year as well, as they can potentially offer silicon excellent performance in signal detection with unprecedented sensitivities, broadband electro-optic modulation, and even lasing and amplification [69–71]. For example, IBM reported the first CMOS-compatible source monolithically integrated on Si by creating an optical parametric oscillator formed by a Si3N4 ring resonator [72]. Error-free all-optical demultiplexing was also realized at 160 Gb/s based on FWM in silicon [73]. In addition, new materials, such as organic groups or graphene, were also proposed to incorporate into silicon photonics [74, 75]. It can be seen that the booming in silicon photonics really happened in the past ten years. It includes the invention of new structures, new phenomena in existing materials, the incorporation of new materials, or more importantly, significant progress in industry. The history above has provided the solid evidence that the optical interconnects based on silicon photonics is moving into practice step by step, and incorporating the photonic circuits into today‟s CMOS microelectronics is no longer a dream now. In the following paragraphs, we will continue to 140


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explore the potentials of silicon photonics in optical interconnects in detail by reviewing the characteristics and the latest techniques in different parts of devices in silicon photonics.

2- Passive Devices 2-1- Waveguides The silicon waveguides are the channels through which optical signals can be transferred from one point to another, and they are the bases of all silicon photonic circuits. As described in introduction, silicon is transparent for the wavelength larger than 1100 nm. The reason is that silicon has a bandgap of 1.12 eV, corresponding to its optical absorption band edge at 1100 nm. For wavelength smaller than 1100 nm, the energy of the photons is larger than its bandgap, so silicon exhibits high absorption. In this case, silicon becomes an effective material for photodetectors or CMOS imaging. In contrast, for wavelength larger than 1100 nm, including the communication bands of 1300-1550 nm, ideally silicon is transparent, and is suitable for the use of making waveguides. Before Silicon-on-insulator (SOI) is developed, one of the most commonly applied waveguide materials was silica. The optical confinement in silica waveguides can be achieved by doping Ge, Ti or P into silica via chemical vapor deposition (CVD) or flame hydrolysis (FHD), so that a contrast in refractive index between the core and the cladding can be created (typically from 0.1% to 0.75%) [9, 13]. Although this provides more flexibility in adjusting the refractive index contrast by changing the doping density, the resulted contrast is still too low to provide strong confinement in silica waveguides. This means that both thick cladding layer (about 50 Îźm) and wide spacing between waveguides are required. Therefore, as shown in

Figure 2,

silicon waveguides are not truly compatible with

electronic IC technology. Similar problems were also found in some other approaches, such as Si3N4 [76], silicon oxynitride on oxide [77], or differently doped silicon [78]. 141


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Figure 2 Comparison of the cross-sections of a CMOS chip, a SOI waveguide, a silica waveguide, and a silica single-mode optical fiber.

Fortunately, after SOI structure was developed, it provided a natural planar waveguide because of its large refractive index contrast (3.5 and 1.46 for Si and SiO2 respectively in 1550 nm wavelength). The resulted strong optical confinement gives the chance to reduce the size of the waveguides, and thin cladding layers can then be used making it compatible with the electronic IC technology (Figure 2). The fabrication of SOI has been achieved by ELTRAN, SIMOX and BESOI, while the most popular one is SmartCut. The details have been summarized in [79]. To minimize the dispersion and increase the data transmission rates, single-mode waveguides are usually preferred in silicon photonics. The calculation of propagation modes in SOI waveguides has been summarized by Reed and Knights in [28]. Briefly, the result can be expressed as:

(1)

where θc is the critical angle for total internal reflection, h is the silicon film thickness, n is the refractive index and k0 the wave vector. It can be seen that for h=3μm, for example, the number of supported modes is still larger than 10. Fortunately, single-mode propagation in this case can be achieved by employing the rib structure (Figure 3a). Although the rib waveguide may be multimode when the film thickness is around several micrometers, the higher order modes “leak” 142


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into the surrounding slab regions when propagating along the waveguides, resulting in an effective single-mode propagation in the rib region [13]. With the development of the processing techniques, silicon films as thin as 220 nm have been intensively applied in recent work [80]. The single-mode transition in this case can be naturally guaranteed. This evaluation in thickness also brings several advantages. As waveguides in sub-micron scale can provide strong optical confinement, the bend radii can be very small, which will effectively reduce the footprints of the devices. In addition, smaller waveguides also make the depletion in modulators much easier. Vertical SOI rib waveguides can be fabricated by the reactive ion etching (RIE). Surface roughness, however, then becomes one dominated challenges when the devices are downsized. Particularly, sidewall roughness, introduced by imperfect etching, will result in scattering at the interface between waveguide core and cladding, forming a major source of propagation loss [81]. The relationship between light transmission loss and sidewall roughness of Si/SiO2 on SOI substrate has been modeled by Lee in [82]. The scattering loss is typically about 0.2-3.0 dB/cm, decided by both sidewall roughness and the size of waveguide. This topic has been analyzed in detail by Yap in [83] as well. Thermal oxidation is a common method to reduce the surface roughness of the waveguides [84]. Usually higher temperature and longer oxidation time are preferred to improve the interface roughness more distinctly, and previous work has shown that 1000 â&#x2014;ŚC is enough to eliminate the profile distortion within 1 hour [85, 86]. The disadvantage of this method is that that this process consumes more silicon, and it also induces residual stress in silicon. High-temperature hydrogen annealing was then developed as an alternative way to reduce sidewall roughness. In hydrogen ambient, the surface mobility of silicon atoms is enhanced, and the migrating atoms in this case can smooth out the surface roughness [87]. Improvement of sidewall roughness was also reported by using surface encapsulation with SiNx [88]. Besides these common techniques introduced above, we also review one novel technique for surface roughness reduction by 143


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using self-perfection by liquefaction (SPEL) [89]. An XeCl excimer laser was used to selectively melt the surface layer of the waveguide, and the average sidewall roughness was reduced from 13 to 3 nm (Figure 3, b) and c)). The major advantage is that this method does not require harsh processing condition, such as high temperature, and the processing is selective. This means that the objective area can be processed without affecting other components on the same chip. Therefore, it provides a good approach for the processing in laboratory. a)

c)

b)

Figure 3 a) The schematic diagram of an SOI ridge waveguide. A 4 Îźm wide Si waveguide on SiO2 with rough sidewalls b) before and c) after exposure under an XeCl excimer laser. Figure b) and c) are taken from [89].

2-2- Couplers As standard single-mode fibers (SMFs) keep being the preferred method for the transmission of optical signals, coupling of optical signals between the fiber and the silicon waveguide is crucial for the communication between the outside world and the photonic chips. Exploring efficient coupling seems to be a conceptually trivial topic, but in practice is a tough problem. As shown previously in

Figure 2,

SMFs used in optical telecommunication typically have a core diameter and mode 144


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size of about 10 μm, compared with less than 1 μm for that of silicon waveguides. Combined with the large difference in refractive index between silica core of SMFs and silicon, the resulted large mode mismatch will cause a big loss in optical power. A variety of techniques exist for performing the coupling task, including prism coupling, end-fire coupling, butt coupling, taper coupling and grating coupling, and the schemes of them are shown in

Figure 4.

Prism coupling introduces an input

beam through the surface of a waveguide by a prime. A specific angle is required to ensure phase matching. Although this method is relatively simple, it is not suitable for SOI waveguide with a surface cladding, as the prism coupling can damage the surface of the waveguide. The principle of end-fire and butt coupling are very similar, as both of them contain a simple process of shining light onto the end of the waveguide. Different from the butt coupling, where the fiber and waveguide are “butted” together directly, the end-fire coupling incorporates a lens to focus the input beam from the fiber into the end face of the waveguide core. However, all of the three coupling techniques listed above become much more difficult as the waveguide dimensions are reduced. This leaves taper and grating coupling the only monolithically integrated means for efficient coupling. Both structures are described in this section, which actually exhibit very complementary performances.

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a) Prism

b) Butt

Input beam

Input beam Waveguide

Waveguide Input beam

c) Lens

d) Taper To optical fiber

Waveguide Lens e) Inverted taper To optical fiber

f) Grating

Polymer Input beam

Waveguide

Figure 4. The schematic diagram of a) prism coupling, b) end-fire coupling, c) butt coupling, d) 3D taper coupling, e) inverted taper coupling and f) grating coupling.

Several different structures have been explored for taper coupling, and the simplest one is the three-dimensional (3D) taper. As shown in Figure 4 d), the coupler starts with a large cross-sectional waveguide area comparable with the optical fibers at the edge of the chips. The cross-sectional area is then gradually reduced, finally to the size of standard rib waveguides integrated on chips. The aim of using the tapers is to reduce the waveguide dimensions in a smooth, lossless transition. The angle of the taper is typically very small to achieve this smooth transition, and it should be produced with very low surface roughness. The advantage of the 3D taper is that it relaxes the alignment tolerances between the input fibers and the waveguides on chip. However, it requires differential etching rate along the length of the taper because of its 3D structure, which makes it current not suitable for commercial purposes. To overcome this problem, Knights and his colleague developed a non-vertical taper, in which a wedgeshaped taper is fabricated on the top of the waveguide so that the mode initially 146


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occupies both the rib and the wedge when launched from the fiber at the edge. Coupling loss of less than 0.5 dB/facet was reported in this structure [90]. Another elegant solution is to employing the inverted taper (Figure 4 e)), or sometimes called spot-size converter. This idea was actually first proposed as a mode converter between a Si3N4 waveguide and a silica waveguide to facilitate laser to fiber coupling [91], and was then extended to fiber-waveguide coupling, for which low loss and broadband operation were demonstrated [92, 93]. A silicon waveguide was tapered to a width below that required supporting a propagating mode at the input facet, through which the fiber mode is coupled evanescently to the taper. As the taper adiabatically widens, the fiber mode is gradually more confined and eventually match the single mode of the waveguide. The silicon waveguide is tapered to a width well below that required to support a propagating mode. The fiber mode is coupled evanescently to the silicon taper and becomes progressively more confined as the taper adiabatically widens to the final singlemode strip waveguide. The taper is covered with polymer layer to avoid large leakage of the evanescent mode at the edge by confining the light into the polymer waveguide. Coupling loss of less than 1 dB per connection has been achieved in this structure [94]. Alternatively, grating-based device provide another approach of transferring power from fibers to small waveguides. Unlike the taper coupling structure, in which the fiber is coupled to the waveguide horizontally, grating coupling can be used to couple light from a fiber perpendicular to the surface into the planar waveguide (Figure 4 f)). Grating coupler can actually be regarded as a waveguide, of which the parameters change periodically. According to the phase-matching condition, the grating period Λ at a given wavelength λ and a coupling angle θ with respect to the vertical should satisfied: ( )

147

(2)


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where k = 2π/λ is the modulus of the out-coupled wave vector, p = ±1, ±2, ±3… is the diffraction order, β = (2π/λ)neff is the real part of the propagation constant, and neff is the mean effective index along one grating period. The first experimental report of grating coupling was from M. L. Dakss and L Kuhn, who made a grating from photoresistor with coupling efficiency about 40% [95]. Grating for 220 nm SOI was then demonstrated in 2004 which had the coupling efficiency of 25% [96], and the efficiency can be improved to over 69% using the backside reflector to suppress the radiation toward the substrate [97]. Compare with the taper-based couplers, grating couplers allow light coupling without the need for dicing and polishing the chip edge, which also makes waferscale testing of nanophotonic circuits possible. Furthermore, alignment tolerance is also relatively large for grating couplers (typically 1 dB additional loss for alignment error of ±1 μm), hence reducing the packaging cost. However, the disadvantage is their high sensitivity to the operating wavelength and to the polarization state. Therefore, both of these two coupling methods are still being explored today, and they may be used complementarily on the same chip in different cases in future.

2-3- Ring resonators Resonant cavity is another essential passive device in silicon-based optical interconnects for their application for wavelength division multiplexing (WDM) and demultiplexing. Similar to the acoustic analogue the tuning fork, the optical resonant cavity provides an enhancement for optical waveforms within a sizedependent resonant frequency spectrum. In silicon photonics, several different structures, including reflective facets (Fabry-Pérot cavities), silicon photonic crystals and whispering gallery, can be used to form the resonance cavities. Reflective facets act as a major component of lasers, surrounding the gain medium and providing feedback of the laser light two mirrors. However, the major disadvantage is that they need a non-planar fabrication process, hence not suitable for on-chip applications. Silicon photonic crystal contains a relatively complex 148


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structure with an area of porous structure to form a 2D diffractive grating. The recent applications for it mainly focus on manipulating the flow of light using slow light, in which the group velocity of the optical pulses is slowed down by the interaction with the surrounding medium to make the light more controllable. a)

b)

c)

e)

d)

Figure 5 The schematic diagram of a) reflective facets, b) silicon photonic crystal, whispering gallery including: c) microring, d) microdisk and e) microsphere.

Whispering gallery is a group of filtering schemes formed by a waveguide bus together with a “gallery” device, including micro-ring, disk, sphere and so on. With suitable design on device size and geometry, the light at specific wavelength can be coupled from the waveguides to the “gallery” devices. Microring-type is the most popular one for filtering purpose because of its inherent 2D geometry and small size. It offers an approach to fabricate the resonant cavities on chip with standard photolithographic techniques. The exchange of optical power between a waveguide and a resonator in the geometry shown in

Figure 5

c) has been analyzed

in detail by Yariv in [98, 99]. We use a1 and b1 to represent the input and the output wave along the waveguide bus, and a2 and b2 for the ones coupled and returned through the ring. Under the conditions that a single unidirectional mode of the resonator is excited and that the coupling is lossless, the coupling can be described by means of two constants t and k in a unitary scattering matrix: | |

|

| |

| |

149

|| |

(3) (4)


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where t is the transmission coefficient through the coupler and k is the mode coupling amplitude, or referred to as coupling coefficient. The coupled and returned wave through the round-trip are also linked by the circulation condition in the ring as a2=b2αejθ, where α is the inner circulation factor which equals to 1 for zero internal loss. The term θ = 2πLneff/λ0 is the round-trip phase shift, where L is the optical path length, neff is the effective refractive index of the ring and λ0 is the free-space wavelength. By substitution, the relationship between the power at input and output terminals can be easily obtained as:

| |

| |

| |

| |

| | | |

| |

| |

| |

(5) (6)

It can be seen that equation (5) reaches the minimum only when cosθ=1. That is, the phase shift is an integral number of wavelength periods, by which the resonance wavelength can be determined. If the ring is a circle with radius of R (i.e. L=2πR), for example, the resonance condition can then be deduced as:

(7) , where m is a integer. In addition, at resonance condition, equation (6) becomes:

| |

(

| |)

(

| |)

(8)

When α = |t|, the power transmitted to the output waveguide bus will be zero, which is also called “critical coupling” condition, and small changes in α for a given t, or vice versa, can control the ratio of power transmitted output waveguide bus from the input. This actually also provides a basis for switching techniques. An add-drop filter can be formed if adding one more waveguide bus at the other side of the ring (Figure 6 a)). When the incident signal passes the ring, the signal 150


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components whose frequency is close to the resonance condition (Equation (7)) will be attracted to the ring and then drop to the other waveguide bus (yellow path), while the off-resonance components will continue to go through the bus to the other terminal (green path). Similarly, if we input another signal at the resonant frequency at the “add” port, this signal can then couple to the waveguide bus above and is added to the transmitted signal there (red path). Two parameters, called free spectra range (FSR) and the extinction ratio (ER), are usually concerned when evaluating the performance of a filter. FSR is defined as the wavelength separation between resonances, and it can be expressed as:

(9)

where ng is the group index of the propagating modes. Larger FSR can provide more wavelength channels and higher aggregated data bandwidth. Therefore, smaller rings are usually preferred for multichannel optical telecommunication purpose. For example, rings with a free spectral range (FSR) larger than 30 nm would require a radius of 5 μm or less, but the size is also limited by the fabrication techniques. ER, which is also called resonance depth, is the ratio between the input and output power (i.e. the power at “through” port) through the filter in dB. High extinction ratio is always expected to avoid so-called coherent cross talk between drop and add data, which can be enhanced by cascade structure. Figure 6 b) shows a merit example which summarizes the explanation above by providing a microring add-drop filter made by MIT [100]. The filter has three stages, where each stage is formed by a three-order ring resonator. The component at the resonant frequency in the input signal will drop down through the rings (yellow path) at the first stage (the leftmost), while the signals far away from the resonant frequency will continue to transmit to the “through” port (green path). The filters at the second and the third stage have the similar but non-identical scheme with the first one, so that the residual signal components near the resonant 151


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frequency at the output of the first filtering stage can be removed. The purpose is to provide flattened passband and high extinction for the “through” ports even with substantial cavity losses. This is a crucial issue for add-drop filters, since the through-port extinction is a simple series-coupled cavity filter is particularly sensitive to errors, and a high extinction is very difficult to realize in practice. Extinction ratio exceeding 50 dB was achieved in the experiment, together with < 2 dB drop loss and a 20 nm FSR [100]. This topology enables flat passbands with lower drop loss and higher through-port extinction tolerance than either cascaded single cavities or a single series-coupled cavity filter of equivalent selectivity. a) Input

Through

Drop

Add

b)

Figure 6 The schemes of a) the basic add-drop filter and b) MIT‟s multistage add-drop filter reported in [100].

3- Modulators An optical modulator refers to a device that can vary the fundamental characteristics of a carrier light beam with respect to an information signal, and it is the workhorse of interconnect technology. A variety of mechanisms can be used for optical modulations in solid-state materials. The first one is Pockels effect. It was firstly studied by Friedrich Pockels in 1893, who found that the refractive index changes in proportion to the applied electric field, so the Pockels effect is sometimes also called linear electro-optic effect. Pockels effect occurs only in 152


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crystals that lack inversion symmetry, such as GaAs or LiNO3. Therefore, since pure silicon is centrosymmetric, Pockels effect is unfortunately absent in silicon. Kerr effect (or quadratic electro-optic effect), firstly discovered by John Kerr in 1875, is a change of refractive index in proportion to the square of the applied electric field, which provides another approach for optical modulation. However, Kerr effect is extremely small in silicon, which was quantified by Soref to be Î&#x201D;n=10-4 for an applied electric field of 106 V/cm at the wavelength of 1300 nm [12]. In addition to changing refractive index in crystals, optical modulation can also be realized by varying the absorption efficiency in materials at a specific wavelength band electrically, and this mechanism includes the Franz-Keldysh effect and the quantum confined stark effect (QCSE). Both of them use an electric field to modify the energy bands of a semiconductor in order to shift the absorption spectra of the material to longer wavelength. They are related with the Franz Keldysh effect being the limit of the QCSE as the quantum-well layers are increased in thickness. Franz-Keldysh effect appears very strong in some III-V group semiconductors, such as GaAs or InP, but it is very weak in Si because of indirect band structure. QCSE is discovered in quantum objects only, and it has been demonstrated to be relatively strong in Ge/SiGe quantum wells [101], hence providing an approach for the modulation in silicon photonics. However, the drawback of using quantum well structure is its complicated fabrication process and the low coupling efficiency between the wells and the standard waveguides. The thermo-optic effect is the thermal modulation of the optical properties of a material. The variation in temperature can cause a change in electron distribution as it alters the band structure and the electron-phonon interaction coefficients of the materials. This change in electron distribution translates into a change in refractive index, and it has been discovered in both crystalline and amorphous silicon [102, 103]. At room temperature, the thermo-optic coefficient of Si is as large as 2Ă&#x2014;10-4/K for telecommunication wavelengths. Despite of this, thermooptic effect is currently not a good modulation mechanism as it takes long time to make large and uniform temperature changes in a material. The speed of the 153


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device operation is therefore expected to be limited to about 1 MHz [104], which is too slow to be practical for modern telecommunications applications. As a result, this leaves plasma dispersion the only practical mechanism for optical modulation in silicon. Plasma dispersion refers to the phenomenon that both the real and imaginary parts of the refractive index Δn changes with the concentration of free carriers in semiconductors. Plasma dispersion in silicon was firstly evaluated by Soref and Bennett in 1987 experimentally for a wide range of electron and hole densities, over a wide range of wavelength [12]. They also quantified the changes in both the refractive index Δn and the absorption Δα corresponding to the variation of carrier density. At a wavelength of 1300 nm, the expression is: (

)

(10) (11)

At a wavelength of 1550 nm, the expression becomes: (

)

(12) (13)

where ΔNe and ΔNh are the changes in free-electron and free-hole concentration respectively. Two options can then be chosen to transfer this change in refractive index to intensity modulation. The first one is to use the refractive index change to shift the relative phase of two propagating waves so that two waves interfere either

constructively

or

destructively

with

each

other.

Mach-Zehnder

interferometer (MZI) is typically used to achieve this job (Figure 7 a)). The MZI modulator is usually realized by using a Y-shape waveguide, which splits the carrier wave into two arms. One silicon modulator is fabricated at one arm, and special design is carried out to shift the phase of the wave by 180 oC if the silicon modulator is on, or no phase shift if off, before it meets the reference wave in the 154


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other arm at the end of the path. The second method is to employ the refractive index change to modify the resonant condition in resonant structure. One example is the ring resonant modulator shown in Figure 7 b). The coupling can be switched between on-and off-resonance state for any given wavelength by applying electric signals to change the refractive index of the ring. MZI modulators usually occupy more area on chip than the ring resonant modulators because of their long optical path. However, the advantages of MZI modulators are their high extinction ration and stability, since the frequency selectivity of ring resonant modulators is very sensitive to the device geometry. a) Input

Input Phase shifter

b) Ring

n+ doped R p+ doped

VF

Input

Output Silicon waveguide

Figure 7 The schemes of a) Mach-Zehnder modulator and b) ring resonant modulator.

Electrical manipulation of the charge density interacting with the propagation of light can be achieved through three common mechanisms: carrier injection, accumulation and depletion, and all of them can be applied for optical modulation in silicon as shown in Figure 8. Carrier injection is the earliest technology to make carrier concentration variation in silicon, which was first proposed by Soref by using a p+-n-n+ diode in 1987 [105]. Later on, p-i-n structure was proved to have a 155


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better performance for modulation to avoid excessive optical loss, and it attracted intensive interests in 1900s. The major advantages of the injection-based modulators are their simple fabrication process and more importantly, the uniform index modulation across the entail waveguide, which gives complete overlapping between the optical mode and the modulation region. However, the key limitation of them is the slow operation speed because of the long minority carrier lifetime, as the injected holes and electrons will take a long time to recombine with each other. Although performance can be improved by reducing device size and optimizing the doping density, the modulation speed is recently limited to a few GHz [39]. In addition, the large power consumption is also a problem because of the forward-bias mechanism. This will not only be harmful to form powerefficient systems but also produce considerable heat when operating for a long time. This, however, will result in an inevitable temperature rise in modulator, which will weaken the refractive index changes because of the thermo-optical effect as introduced above [106].

a)

b)

Optical mode

Oxide barrier

Metal

Metal

Metal

Metal

p+

n+

p+

n+

c) Metal p+

p

n

Metal n+

Figure 8 The schemes of typical silicon optical modulator based on a) carrier injection, b) carrier accumulation and c) carrier depletion.

In accumulation-based modulators, a thin insulator layer is used to isolate two halves of the waveguide so that a capacitor can be formed. Carried density can be changed by charging or discharging the capacitor inside the modulators. Alternatively, metal-oxide-semiconductor structure can be employed for the same 156


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purposes. Compared with injection-based modulators, there are no slow carrier generation and/or recombination processes involved in the accumulation operation, so higher speed for modulation is achievable in this case. By using this mechanism, the first all silicon modulator with a modulation bandwidth in excess of 1 GHz was achieved by Liu from Intel in 2004 [34]. The performance was then improved to 10 GHz one year later [37]. In depletion-based modulators, the waveguide is slightly doped into p- and n-type to form a p-n diode. The diode is reverse biased so that the carriers originally located in the intrinsic region at equilibrium can be depleted, resulting in a change of carrier density. Similar to the accumulation mechanism, the operation of depletion-based modulators also relies on electric-field induced majority carrier dynamics, which makes them survive from the limitation from the minority carrier lifetime. The difference is that depletion-based modulators have much smaller device capacitance compared to that of the MOS capacitor used in accumulation type. In addition, the depletion-based modulators also have the superior in fabrication simplicity compared with the other two types, which make them currently the mainstream for silicon modulators. One milestone work was achieved by Liao from Intel by achieving a silicon depletion-based modulator with a record high data rate of 40 Gb/s in 2007 [45] (Figure 9). The modulator is based on an asymmetric MZI structure with a pn diode embedded in each of the two arms. The pn diode phase shifter comprises a p-type doped Si rib waveguide and an n-type doped Si formed using an epitaxial Si growth process. To ensure ohmic contact between Si and metal, the rib edge are heavily doped for the Ncontact and P-contacts. This device has 30 GHz bandwidth and can transmit data up to 40 Gbit/s. A few depletion-based modulators with similar performance have also been reported since then [107, 108]. Most recently, the performance has been improved to 50 Gb/s by Reedâ&#x20AC;&#x;s group [109].

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Traveling-wave electrodes signal

ground

n-Si p++

ground n++

P-Si oxide

p++

waveguide Si substrate Figure 9 The schematic cross-section of pn diode waveguide phase shifter in Intelâ&#x20AC;&#x;s 40 Gb/s modulator. The figure is taken from [45].

4- Detectors Photodetectors are a group of devices with the capability of converting light into electrical signals, and they act as the receivers of the whole optical interconnect loop. Semiconductor-based photodetectors employ a p-n junction structure. When photons are absorbed near the p-n junction, electron-hole pairs are produced, and electron-hole pairs are then separated, electrons drifting to the n-region and holes to the p-region. This results in a light induced electromotive force across the junction, by which the optical signals can be detected by measuring this potential. Alternatively, by reverse biasing the diode, optical current can be produced, which provides another mechanism for optical detection. An intrinsic area is usually added to form a p-i-n structure, so that the majority of applied voltage can drop across this intrinsic region for higher detection efficiency. Silicon has a bandgap of 1.12 eV, which ensures their high sensitivity in optical detection around the wavelength of 700 nm, and it is still useful for short-haul telecommunication wavelength of 850 nm. However, beyond 1100 nm (the absorption edge of silicon), silicon becomes transparent, which prevents its application for the direct optical to electrical conversion required in integrated photonic circuits operating at 1330 nm and 1550 nm. There are three main 158


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branches to overcome this problem: using Schottky structure, Ge detectors or IIIV heterogeneous detectors built on silicon substrates. Schottky photodiodes are based on the internal photoemission effect over the metal-semiconductor Schottky barrier. The detector consists of a metal layer built on a lightly doped Si waveguide, forming a Schottky contact at the material interface. The absorption wavelength in this case is determined by the Schottky barrier, and only incident photons with energy larger than the barrier can be absorbed by the junction. The barrier is usually 0.2-0.6 eV depending on different metals, and this offers a mean for the extension of the detection range for silicon. Schottky-based detectors have the advantages of high switching speed and simple fabrication process, but the drawback is that a small fraction of the incident photons actually causes photoemission because of the small interaction volume for photons and electrons in the metal. One latest example made by Levy is shown in [110]. The detector is fabricated using a self-aligned approach of localoxidation of silicon on silicon-on-insulator (SOI) substrate to define the nanoscale waveguide structure by oxide spacers, followed by a deposition of gold layer to form the Schottky contact. The detector has achieved the responsivity of 13.3 mA/W for incident optical wavelength of 1300 nm. However, further efforts are still required in Si Schottky detectors to improve the detection responsivity before applying them in practice. a)

b)

Figure 10 a) The SEM micrograph of the Schottky contact and b) Intensity mode profile of the plasmonic waveguide (Schottky contact) in [110]. The figures are taken from [110].

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With a smaller bandgap of 0.67 eV, germanium has strong absorption for the wavelength deeper into the infrared toward 1880 nm, which gives a motivation for the early work on photodetectors with an active layer of SiGe [111] or SiGeC alloys [112]. It is currently the most popular approach for photodetecting in silicon photonics because of the high compatibility of Ge with CMOS process. The problems of growing Ge on Si mainly result from the 4.2% lattice constant mismatch between Ge and Si, which introduces the crystalline defects such as dislocations. The dislocation defects are usually manifested in the form of large dark current, which is harmful for the sensitivity of the photodetectors. The first successful approach for growing high-quality epitaxial Ge layers on Si was reported Luryi in 1984, in which a graded SiGe buffer layer was deposited to reduce the dislocation density in the Ge layer [113]. The quality of Ge layers can be improved further by optimizing the SiGe graded buffer layers and growth temperature [114–116]. Ge films with dislocations of less than 2 ×106 cm-2 were successfully demonstrated [117]. At the end of 1990s, directly growing Ge on Si substrate was realized by using a two-step growth technique [118]. A thin epitaxial Ge layer of 30-60 nm is grown on Si first at low temperature around 320-360 ◦C to suppress the islanding of Ge. At this stage, Ge layer has been thick enough (> 30 nm) to eliminate the influence of lattice mismatch on further epitaxy. Then, in the second step, Ge is continuously grown on Si with a higher temperature of T>600 ◦C for higher growth rates and better crystal quality, followed by a post-annealing at T>750 ◦C to reduce the dislocation density by two orders of magnitude. To reduce the growth temperature, the thin Ge buffer layer can be altered by combining thin SiGe buffer layers with the two-step Ge growth approach [119]. Plasma-enhanced CVD has been report to replace traditional low-pressure CVD to achieve highquality Ge at a lower thermal budget [120]. Recently, Feng has reported a highperformance Ge detector produced by two-step growth introduced above (Figure 11). The detector is formed with a p-i-n diode integrated in a SOI waveguide and it has a large bandwidth of 32 GHz and a responsivity of 1.1 A/W at a wavelength 160


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of 1550 nm under a reverse bias of 1V [121]. Ge photodetectors with extremely high bandwidth 42 GHz, together with a high responsivity of 1 A/W and a low dark current density of 60 mA/cm2 under -4V bias, was also achieved by Vivien using similar methods [122]. In addition, Ge photodetectors have also been successfully integrated into CMOS circuits by Luxtera and MIT using standard 130 nm and 150 nm CMOS process respectively [123, 124] a) p-Ge

n-Ge

Metal

Metal i-Ge Si slab Box Si substrate

b)

Figure 11 a) Cross-section view and b) SEM image of the p-i-n region of the Ge photodetector reported by Feng. Figures are reproduced from [121].

Compared with silicon, III-V-based detectors usually have a much wider absorption bandwidth and lower dark current. Although III-V materials usually suffer from the compatibility problem with CMOS fabrication process, recently, some of them, like InGaAs detectors, have been successfully built on the SOI 161


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substrate using heterogeneous integration approach [125, 126]. The InGaAs p-i-n diode is built on the silicon waveguide via a benzocyclobutene bonding layer, and the light is coupled from the waveguide to the detector vertically. The device was reported to have a responsivity of 1.1 A/W and very low dark current of 10 pA at a wavelength of 1550 nm [126]. In addition, quantum confinement can be applied to III-V detectors to enhance their performance in long wavelength infrared range, which is useful for some commercial applications, like medical imaging, gas sensors, surveillance devices or high temperature detection [127, 128].

5- Light Sources The absence of efficient light sources is the greatest challenge in silicon photonics over the past two decades. Silicon is an indirect bandgap semiconductor, which means that the radiative recombination of electron-hole pairs in pure silicon is a phonon-assisted process for momentum conservation. Such second-order radiative recombination does not occur frequently, which results in a very long radiative lifetime (on the order of milliseconds) in silicon. This in turn makes electrons and holes more likely to be trapped in defect and release the energy into heat before they recombine with each other radiatively. As a result, the optical conversion efficiency of silicon is only around 10-6 at room temperature [28]. Although the non-radiative recombination can be minimized by using extremely high-purity samples, the conversion efficiency is still limited to only 10-4-10-3 [129]. In addition, there are two further mechanisms that also limit the population inversion in silicon. The first one is called Auger process, in which, instead of producing a photo, the energy of electron-hole recombination is absorbed by another free carrier to stimulate it to a higher energy state. This recombination mechanism is active as soon as more than one carrier is excited, and it becomes more serious for higher density of excited carriers and inversely proportional to the bandgap energy [130]. The second mechanism is the free-carrier absorption, in which the incident energy is directly absorbed by a free carrier in Si. Empirically, the absorption 162


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coefficient is related to the free-carrier density nf and the light wavelength λ as αn∼10-18nfλ2 at room temperature [131]. A schematic diagram is provided in Figure 12

as a summary of the transitions introduced above.

Auger

Energy

FCA

Momentum

Photon

Figure 12. Schematic diagram of three possible transitions in bulk Si, including radiative recombination (green), Auger recombination (red), and free-carrier absorption (blue).

Early research focused on using quantum confinement to squeeze light from pure Si. According to the Heisenberg uncertainty principle, when an electron is localized, its momentum becomes uncertain, which gives a chance to overcome the problem of indirect band in silicon. The first successful trial was made in 1990 by Canham who got light emitting by using porous silicon [132], but the application of this method is limited by the instability of porous silicon. Alternatively, people also tried to introduce silicon nanocrystals (Si-nc) into a SiO2 layer, or similar use natural silicon-rich oxide (SRO). Lasers based on this idea can be pumped optically or electrically using a tunnel diode, and the emission spectrum is around 800-900 nm [22, 133]. The emitting wavelength can be transferred to telecommunication range by doping trivalent erbium ions (Er+3) into SRO. Er+3 ions have an incomplete 4f electronic shell, permitting intra-4f transitions when pumped optically. Moreover, the first excited state is at energy of 0.8 eV, which permits the emission of photons at a wavelength of 1535 nm. 163


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Although Si is not a good host for Er+3 ions, Er-doping SRO instead provides a mean to realize emission at telecommunication wavelength. Photoluminescence and

electroluminescence

have

both

been

reported

[134].

Meanwhile,

photoluminescence can be enhanced by co-implantation with elements of carbon, hydrogen or silver [135, 136], while the improvement of electroluminescence can be realized by optimizing device geometries and doping distributions [137]. For example, multilayer Er-doped SRO/Si-nc structures has been proved to have higher emitting efficiency than single-layer structures [138]. Another approach to obtain luminescence in silicon was developed in 2002, when Jalali reported the observation of Raman scattering in silicon [139]. Two years later, the first silicon laser based on stimulated Raman scattering (SRS) was developed [33], followed by the demonstration of continuous-wave (CW) lasing in 2005 [41]. Raman scattering is an inelastic scattering of photon by an optical phonon. After incident photons are absorbed by atoms, the system energy is raised temporarily. Soon after that, the system quickly drops back to the original state, and release the energy in a form of photons in the same frequency (known as Rayleigh scattering). A small fraction of released photons, however, have the frequency different the incident frequency, which is called Raman scattering. Compared with the incident photons, the scattered photons can have a lower or higher frequency, which are called Stokes and anti-Stokes transitions respectively. With the presence of pump beam that stimulates atoms to a higher vibrational level, if one photon with a frequency resonant at Stokes transition is also injected to the crystal, it may trigger another triggers the generation of another Raman Stokes photon. Optical amplification is realized in this process, and this phenomenon is called stimulated Raman scattering. Although Raman effect is very strong in Si (104 times larger than in glass fiber), which allows for significant gain in centimeters, the amplification efficiency is still limited by the free carrier absorption (FCA) in Si. In order to suppress FCA, the energy of the pump should be smaller than the bandgap of Si. This, however, results in another optical loss mechanism called two-photon absorption (TPA), in which an electron 164


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simultaneously absorbs of two photons and is elevated to the conduction band. Because the carrier density is increases in this process, TPA will further introduce optical loss from FCA. One common solution is to employ a p-i-n diode structure in the waveguide, where a reverse bias is applied to sweep the free carriers away. Typical examples can be found in Intel‟s CW Raman lasers [41] (Figure 13). The laser is based on a S-shape p-i-n waveguide path, together with the reflective facets at both ends to form the cavity (Fabry-Pérot resonance structure). The p-i-n diode structure is reverse biased to sweep out the free carrier to reduce the TPAinduced FCA. Alternatively, the Raman conversion efficiency can also be enhanced by adding dopants, such as Nd or As [140, 141], or using nano-disk structure [142]. In addition, giant Raman gain was also discovered in silicon nanocrystals at the end of 2012, which is up to four orders of magnitude greater than in crystalline silicon, and this may provide a new chance to develop Si Raman lasers with high performance [143]. a)

b)

Figure 13 a) Scheme of the silicon laser cavity with the facets and a p-i-n structure along the waveguide; b) SEM cross-section image of a silicon rib waveguide with a p-i-n diode structure. Figures are taken from [41].

Despite the progress in silicon-based lasers introduced above, electrically pumped lasers with high efficiency are still expected. Compared with Si, practical lasing is easy to be obtained under electrical pumping in many III–V compounds, such as GaAs and InP, because of their direct bandgap structure. Therefore, currently combining III-V semiconductor active materials onto the silicon platform is more likely to be the solution, which can be realized either on-chip or off-chip. Off-chip lasers based on III-V compounds can supply the lasing at a high power level, and 165


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one distinct advantage is that the serious heat produced by the lasers will not be released to the Si platform, which offers a long-term stability of the whole system. Additional coupling process, however, is required to transfer the light from the lasers to the waveguides on chip in this chip. The topic of coupling has been discussed in part 0, so we mainly review the on-chip lasers here. The fundamental roadblock in monolithically integration of III-V lasers mainly focuses on a large mismatch of lattice constants and thermal expansion coefficients between III-V materials and silicon. For example, GaAs and InP have lattice mismatches of 4.1% and 8.1% and thermal expansion coefficient mismatches of 120.4% and 76.9%, respectively when comparing with Si. These will result in a threading or misfit dislocation density after the lasers operate for a relatively long time. There are two branches, namely epitaxial growth and hybrid, to overcome this challenge. The difference is that in epitaxial growth, lasers are directly built on the SOI substrate using metal-organic chemical vapor deposition (MOCVD), molecular-beam epitaxy (MBE) and so on, while hybrid lasers with a complex structure are usually first grown on some other substrates like Ge or SiO2, and then transferred onto the SOI substrate by wafer bonding. One common method to overcome the crystal mismatch in epitaxial growth is to employ a buffer layer between the III-V lasers and the Si substrate. For example, epitaxial techniques with SiGe or GaSb buffer layers have provide a feasible medium to build GaAs-based CW lasers on Si substrates at room-temperature, but the reliability is still a problem [144–146]. Alternatively, recent work on the
 epitaxial growth has involved
 the deposition of Ge on Si because of the compatibility with CMOS technology [147]. Despite the indirect-band structure of Ge, it has a direct bandgap at 0.8 eV compared with its indirect bandgap (0.66 eV), which makes electroluminescence possible in Ge. In 2010, MIT reported the first Ge on Si CW laser at 1520-1620 nm at room temperature. Band engineering was applied to this experiment, where the Ge layer was first grown at 650 °C and then cooled to room temperature, accumulated with a thermal induced tensile strain of 0.24%. Emitting efficiency can be increased in this process, as this 166


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tensile strain shrinks the direct bandgap of Ge to 0.76 eV, and the difference between the direct band and the indirect gap is reduced. Further compensation of this energy difference was realized by doping with phosphorous [63, 64]. Most recently, electrically pumped lasing in Ge-on-Si structure at room temperature has also been reported by the same group [148]. The Ge laser has a p-n-n heterojunction structure, and it is monolithically integrated into a CMOS process by doping phosphorous in Ge, the laser has an output power of 1 mW, together with a gain spectrum of nearly 200 nm. In addition to the progress above, attempts to „convert‟ silicon into
a platform for growing III-V lasers by a thin virtual substrate, such as InP, are in progress [149]. Wafer bonding technique provides a method to combine epitaxial films with low threading dislocation densities to the lattice-mismatched Si substrate. In this case, the III-V lasers can be made in a relatively complicated structure for better performance, since the fabrication is independent from other components on SOI substrate. The first electrically driven hybrid laser on silicon (Figure 14) was realized by Intel and UCSB in 2006 [43]. The laser consists of two a silicon waveguide etched with distributed-Bragg-grating mirrors for resonance and an InP chip for light source. The chip is bonded to the silicon at low temperature by oxygen plasma-assisted wafer bonding. An oxide layer is formed roughly at the interface in this process, which acts as the glue that fuses the two materials together [150]. A hybrid AlGaInAs-silicon laser with continuous wave at 1568 nm was reported in the same year as well using the same technique [151]. In addition, adhesive bonding technique can also be applied to bond III-V lasers to Si, and a III-V/Silicon Fabry-Pérot laser with 5.2 mW output power in continuous-wave regime at 1310 nm has been reported using this technique [152].

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Figure 14 Scheme of the hybrid laser structure developed by Intel and UCSB with the optical mode superimposed. Figure is taken from [43].

Finally, we review a latest technique of bonding III-V lasers on Si in large scale, called transfer printing (Figure 15) [153, 154]. An array of AlInGaAs lasers are first grown on a native substrate of GaAs. Instead of bonding entire III-V wafers to silicon, a microstructured elastomeric stamp is used to selectively release the lasers on GaAs and “print” them onto the Si substrate. The stamp are secured to the silicon surface by van der Waals forces, so heat or pressure are not required in this process. The laser has an optical output power of 15 mW per facet at room temperature. This method not only allows wafer-scale integration of III-V lasers on Si, but also provides an “economic” use of the III-V wafers.

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a)

b)

c)

d)

Figure 15 Scheme of the transfer-printing process. Figures are taken from [153].

6- Commercial Progress Currently, silicon photonics is proposed to have a wide applications, including broad-band Ethernet, on-chip interconnects for high-speed computers or the sensors for biomedicine, gas, evanescent or optical field, and so on [9], while the first two are perhaps the ultimate goals for this subject. Utilizing all the components in silicon photonics introduced above, it is now possible to build the on-chip photonic network for optical interconnect purpose. One ultimate goal of silicon photonics is to build the connecting loop between each component on central processing unit (CPU). Stacking structure in three dimension is now regarded as one promising platform to combine optical interconnects with CMOS processors, and the interconnect layer can be put at the bottom, in the middle or on the top of the system. Figure 16 gives an example of monolithic integration, where the cores of the processor and the optical interconnects are built at the top and the bottom of the stacks respectively, sandwiching several layers of local memory in the middle [155]. In inter-chip communication, the electrical signals 169


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from one core first goes down to the optics layer and is converted into the optical signals by the modulators. The optical signals transmit through the waveguides and go up to the objective core before being converted back to electrical form. Different from the copper-based interconnects, the communication signals between several cores can be transmitted simultaneously in the optical circuit by wavelength division multiplexing (WDM), permitting much larger bandwidth compared with metal interconnects. Off-chip communication can be realized using the same chip by adding the coupling system to the external photonic interfaces.

Figure 16 Scheme of the monolithically integrated silicon stack with multi-core processor, memory, and optical interconnects. Figure is taken from [155].

With the impetus from the potential of high-speed optical interconnects, the past three years is a fruitful period for the commercial application of silicon photonics. For example, at the end of 2012, IBM reported a pioneer work by integrating different photonic components side-by-side with electrical circuits on a single silicon chip, for the first time, in standard 90nm semiconductor fabrication process [156, 157]. The cross-section and angled view of the chip can be found in Figure 17. The system contains a Ge photodetector and a silicon modulator, which are integrated with the silicon transistors side-by side. The carrier light is generated by an off-chip laser and is input to the chip by a coupler. As shown in Figure 17 b), the optical signals and electrical signals are transmitted in waveguides and 170


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copper wires respectively on the same plane. The optical loop on chip forms a multi-channel WDM optical transceiver, and the performance is measured to be 25 Gb/s with extended temperature range up to 90 oC.

a)

b)

Figure 17 a) Scheme of IBMâ&#x20AC;&#x;s CMOS chip with a Ge photodetector (red feature on the left side of the cube) and a modulator (blue feature on the right side of the cube) fabricated side-by-side with silicon transistors (red sparks on the far right of the cube). Photonic circuits and silicon transistors are interconnected with nine levels of yellow metal wires; b)â&#x20AC;¨Angled view of a portion of IBMâ&#x20AC;&#x;s CMOS chip showing optical waveguides (blue) and copper wires (yellow). Figures are taken from [157].

Besides, huge progress was also made in optical communication in Ethernet as well. One distinct breakthrough was also reported at the end of 2011 by Luxtera, which delivered world first single-chip integrated optoelectronic transceiver at 100 Gb/s level, with the maximum data rate at 112 Gb/s [158]. This silicon-based optical device comprise four 28 Gb/s transmit and receive channels powered from a single laser and the on-chip component are fabricated with standard CMOS process. The chip was soon used by Molex for 100 Gigabit active optical cables at the beginning of 2012. This achievement in silicon photonics enables development of products for parallel 100 Gb/s InfiniBand and Ethernet, as well as the serial 32G Fiber Channel applications. In 2012, Luxtera also announced that it had shipped over one million Gigabit silicon CMOS photonic channel devices over the past ten years [159], and this is also an important milestone for the silicon photonic industry. In the same year, Kotura also reported its low-power 100 Gb/s 171


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optical network based on silicon photonics [160]. In the past two years, Oracle has achieved 10 Gb/s silicon transceiver circuits with power consumption of 530 fJ/b on a 40 nm CMOS SOI platform [161], and an 80 Gb/s arrayed CMOS silicon photonic transceivers using 130 nm CMOS SOI platform [162]. Therefore, silicon photonics has definitely provided crucial contribution to the next generation of optical interconnects at 100 Gigabit level.

Conclusions This chapter has reviewed the development of silicon photonics for optical interconnects, starting from a description of history. The progress in each component in silicon photonics, namely waveguides, filters, modulators, detectors, and lasers has been introduced in detail one by one. The commercial applications of silicon photonics for interconnect purposes have finally been introduced. It can be seen that silicon-based optical interconnects is now no longer a future for both on-chip and off-chip applications. The monolithic integration of optical and electronic devices on silicon platform is now on progress for the replacement of copper-based interconnects to extend the Mooreâ&#x20AC;&#x;s Law. In off-chip applications, a variety of evidence has been supplied for the capability of silicon photonics in the next generation of Ethernet at 100 Gigabit level. It is proposed that Si will play a leading role for the Ethernet and high-performance computing systems in the near future.

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