POLY(FERROCENYLDIMETHYLSILANES) AT THE ...IV Poly(ferrocenyldimethylsilanes) at the Interface of...

POLY(FERROCENYLDIMETHYLSILANES) AT THE INTERFACE OFCHEMISTRY AND MATERIALS SCIENCE:

SYNTHESIS, STRUCTURE-PROPERTIES AND THIN FILM APPLICATIONS

PROEFSCHRIFT

ter verkrijging van

de graad van doctor aan de Universiteit Twente,

op gezag van de rector magnificus,

Prof. dr. F. A. van Vught,

volgens besluit van het College van Promoties

in het openbaar te verdedigen

op vrijdag 9 juni 2000 te 15.00 uur.

door

Rob Gerhardus Hendrikus Lammertink

geboren op 3 april 1972

te Enter

II

Dit proefschrift is goedgekeurd door:

Promotor: Prof. dr. G. J. VancsoAssistent-promotor: dr. M. A. Hempenius

IV

Poly(ferrocenyldimethylsilanes) at the Interface of Chemistry and Materials Science:Synthesis, Structure-Properties and Thin Film Applications / Rob G. H. Lammertink

Thesis University of Twente, Enschede, The NetherlandsISBN 90 36514517

Cover illustration: AFM image of block copolymer thin film.Press: Print Partners Ipskamp, Enschede.

Contents

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1 NANOTECHNOLOGY AND INORGANIC POLYMERS 1

1.1 Science and Technology in Diminishing Dimensions 1

1.2 Polymers and the Periodic Table 21.2.1 The Birth of Metallocenes 31.2.2 Ferrocene-Containing Polymers 4

1.3 Ring-Opening Polymerization of Ferrocenophanes 41.3.1 Poly(ferrocenylsilanes) via Ring Opening Polymerization 5

1.4 Nanostructured Materials by Block Copolymer Self-assembly 6

1.5 Concept of this Thesis 7

1.6 References and Notes 8

2 HOMOPOLYMER AND BLOCK COPOLYMER ASSEMBLIES 13

2.1 Semi-Crystalline Homopolymers 132.1.1 Morphology 142.1.2 Crystal Melting Temperature 152.1.3 Crystallization Kinetics 15

2.2 Block Copolymers 162.2.1 Bulk Morphology 172.2.2 Identification of Block Copolymer Bulk Morphologies 192.2.3 Conformational Asymmetry 212.2.4 Homopolymer/Block Copolymer Blends 232.2.5 Block Copolymer Thin Films 24

2.3 Surface Treatment of Polymers 27


3 SYNTHESIS, CHARACTERIZATION AND THERMALPROPERTIES OF FERROCENYLDIMETHYLSILANE POLYMERS,OBTAINED BY ANIONIC RING-OPENING POLYMERIZATION 37

3.1 Introduction 37

3.2 Synthesis and Characterization of Poly(ferrocenylsilane) 38

VI

3.3 Thermal Properties of Poly(ferrocenyldimethylsilanes) 403.3.1 Glass Transition 403.3.2 Crystallization Kinetics 423.3.3 Morphology of Melt-Crystallized Poly(ferrocenylsilane) 443.3.4 Melting Behavior of Poly(ferrocenylsilanes) 47

3.4 Conclusions 51

3.5 Experimental 52


4 POLY(FERROCENYLDIMETHYLSILANES) FOR REACTIVE IONETCH BARRIER APPLICATIONS 55

4.1 Introduction 55

4.2 Results and Discussion 584.2.1 Oxygen Reactive Ion Etching 584.2.2 Tetrafluorocarbon Reactive Ion Etching 63

4.3 Conclusions and Outlook 65

4.4 Experimental 65


5 PERIODIC ORGANIC-ORGANOMETALLIC MICRODOMAINSTRUCTURES IN POLY(STYRENE-BLOCK -FERROCENYL-DIMETHYLSILANE) COPOLYMERS AND BLENDS WITHCORRESPONDING HOMOPOLYMERS 69

5.1 Introduction 69

5.2 Sequential Anionic Polymerization of Styrene and 1,1’-Dimethyl-silylferrocenophane 71

5.3 Bulk Morphology of Neat Diblocks 73

5.4 Temperature Induced Transitions 75

5.5 Diblock-Homopolymer Blends 79

5.6 Conclusions 84

5.7 Experimental 86


Contents

VII

6 MORPHOLOGY AND SURFACE RELIEF STRUCTURES OFASYMMETRIC POLY(STYRENE-BLOCK -FERROCENYLSILANE)THIN FILMS 91

6.1 Introduction 91

6.2 Results and Discussion 946.2.1 Thin Films on Silicon Substrates 946.2.2 Freestanding Diblock Copolymer Films 102

6.3 Conclusions 104

6.4 Experimental 105


7 MORPHOLOGY AND CRYSTALLIZATION OF THIN FILMS OFASYMMETRIC ORGANIC-ORGANOMETALLIC DIBLOCKCOPOLYMERS OF ISOPRENE AND FERROCENYLDIMETHYL-SILANE 109

7.1 Introduction 110

7.2 Results and Discussion 1117.2.1 Anionic Synthesis of Poly(isoprene-block-ferrocenylsilanes) 1117.2.2 Thin Film Morphology and Stability 112

7.3 Conclusions 121



8 NANOSTRUCTURED SURFACES BY A COMBINATION OF SELF-ASSEMBLY OF BLOCK COPOLYMERS AND REACTIVE IONETCHING 125

8.1 Introduction 125

8.2 Results and Discussion 1278.2.1 Patterns by Self-Assembly of Diblock Copolymer Thin Films 1278.2.2 Patterns by Self-Assembly on Different Length Scales 132

8.3 Conclusions and Outlook 134



VIII

SUMMARY 139

SAMENVATTING 141

PUBLICATIONS LIST 143

DANKWOORD 147

CURRICULUM VITAE 149

Nanotechnology and Inorganic Polymers

1

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1.1 Science and Technology in Diminishing Dimensions

A large amount of research and industrial interest is focussed on subjects that can behosted under a single umbrella called “Nanotechnology.” The demand is driven by a singlereason: Scale. For example, electronic systems are rapidly moving toward small, fast andhigh-density information storage devices. Already in 1959, Richard Feynman addressed thepossibility of manipulating matter in his visionary lecture, “There’s plenty of room at thebottom,” which has become a “classic” among micro- and nanoscientists nowadays.1 He was“…inspired by the biological phenomena in which chemical forces are used in a repetitiousfashion to produce all kinds of weird effects…” Nature already demonstrated that theinformation for the organization of a complex creature such as ourselves can be stored in verysmall cells. In his talk, Feynman foresaw the importance and technological applications ofmanipulating and controlling things on a small scale. Whether a structure is called small,depends on one’s viewpoint. Nanostructures, which are assemblies of molecules withdimensions in the range of 1 – 100 nanometers, will be considered small by electricalengineers and materials scientists, but are large to chemists.

As dimensions become smaller and smaller, new technologies and methodologiesneed to be developed in order to construct architectures at length scales characteristic ofatoms and molecules. Several strategies have so far been investigated to obtain and control(near) nano-sized structures. When using optical lithography in a top-down approach, thedecrease in feature sizes has been made possible by the application of lasers with shorterwavelengths.2 At present, 193 nm lithography using an ArF laser is a leading candidate forprinting devices with features of the size of 0.18 µm. Further decrease of wavelengths has ledto proximity x-ray,3 extreme UV,4 electron projection5 and ion projection lithography.6

Physical printing techniques like hot-embossing lithography7 and microcontact printing8 alsoappear to be good candidates to cheaply reproduce complex patterns on a sub-micron scale.

However, a variety of technological applications, including the fabrication of denselypacked magnetic domains9,10 or the use of silicon posts as model synthetic electrophoresisgels,11 require only a regular array of domains on a surface. Laser interference lithography(LIL) is a maskless technique that allows the preparation of sub 100-nm lines and dots.12

Using bottom-up self-assembly strategies, regular patterns can be generated on a variety of

Chapter 1

2

length scales. The application of self-assembly to form functional systems has been used inmany examples, for instance block copolymers,13,14,15 colloidal suspensions,16 surfactants,17

and proteins.18

Scanning probe microscopes, for example scanning tunneling microscopes (STMs),19

atomic force microscopes (AFMs),20 near-field scanning optical microscopes (NSOMs),21

and scanning electrochemical microscopes (SECMs),22 are able to image and modify surfaceswith atomic resolution. Although these microscopes were originally designed to provide highresolution images of surfaces, their lithographic capability was demonstrated in a set ofexperiments with an STM, just five years after the first STM images were recorded.23

Although these methods are serial techniques, and have limited writing speeds, they are“enabling technologies” for nanoscience.24

1.2 Polymers and the Periodic Table

The progress in organic polymer synthesis has been immense during the past 50 years.Polymers built from readily accessible organic units find numerous applications aselastomers, plastics and fibers in a variety of products.

In contrast to the versatility of organic polymer chemistry, the incorporation ofinorganic elements in macromolecules has been very limited. This can be understood if oneconsiders the main synthetic routes to organic polymers. Chain growth mechanisms canhardly be used since inorganic species that contain stable multiple bonds and that are suitablyreactive enough are difficult to prepare. The preparation of difunctional inorganic monomersfor polycondensation type reactions is also difficult, and the stoichiometry needed for thistype of reactions further limits the applicability. Polymerization via ring-opening seems to bevery promising, since inorganic ring chemistry is quite well-developed.25

Si

R

R

n

Si O

R

R

P

R

R

N

n n

Chart 1.1 Polysilane, polysiloxane and polyphosphazene are among the best-known examples of “inorganic”polymers that find many applications in present day life.

The few well-established and commercialized examples of “inorganic” polymersshow very diverse and interesting properties.26 Among them are polysilanes, polysiloxanes,and polyphosphazenes (Chart 1.1). Polysilanes can for instance be used as ceramicprecursors,27 photoresists28 or photoinitiators.29 Polysiloxanes, currently the commerciallymost important inorganic polymer systems, are used in a large number of applications,including mold-release agents and as biomedical materials.30 Polyphosphazenes consist of an


3

alternating phosphorus and nitrogen backbone31 and comprise the largest class of inorganicpolymers. These materials, often called “inorganic rubbers”, are of interest as biomedicalmaterial, or for example as polymeric electrolyte components in battery technology.30b

Organometallic based polymers can be divided in two major classes,32 namelypoly(metallocenes) and rigid-rod organometallic polymers like metal-acetylide compounds(Chart 1.2). Insoluble and uncharacterized polyacetylides with Cu or Hg were reportedalready in 1960.33 It took 15 more years before the first soluble Pt and Pd metal poly-yne wassynthesized.34 Since then much progress has been made on synthetic routes to similarstructures, incorporating group 8,35 9,36 and 1036,37 metals into metal acetylide polymers.These polymers display interesting non-linear optical, electrochemical and electricalproperties, due to the presence of low oxidation state transition metals and delocalized,extended π-electron systems.

M C C C C M C C C Cx

nn

Chart 1.2 Metal-acetylide polymers incorporate metal atoms, depicted by M , within an extended chain.

1.2.1 THE BIRTH OF METALLOCENES 38

Fe

Chart 1.3 Illustration of the ferrocene “sandwich” structure, and the electronic bonding.

In 1951, two groups of chemists independently prepared ferrocene by “accident”.39,40

An air-stable, sublimable, orange solid that melted at 173 °C and was soluble in organicsolvents was reported by both groups. Simple σ-bonding between the iron and the organicpart would be relatively unstable and ionic interaction could not explain the volatility of thesolid. The correct structure of ferrocene was reported shortly afterwards by Wilkinson andFischer,41,42 for which they shared the Nobel Prize for Chemistry in 1973. The fiveelectronically equivalent carbons in each ring had to participate with the bonding to iron in anequal way. In the resulting “sandwich” structure, the metal d orbital overlaps strongly with

Chapter 1

4

the carbon π-electrons in the p orbital (Chart 1.3).43 It was discovered that thecyclopentadienyl rings behaved chemically very similar to aromatic species. It was thisdiscovery that spurred the field of metallocene chemistry.

1.2.2 FERROCENE-CONTAINING POLYMERS

Poly(ferrocenylene), synthesized in a simple free-radical process, was first reported in1960.44 Low molar mass materials were obtained in low yields, but the quality and quantityof the materials gradually improved as better synthetic routes were developed.45,46

Copolymers that contain ferrocene in the main chain have been prepared in widevariety, featuring ester or amide linkages, and alkene, alkyne or aromatic repeating unitsalternating with ferrocene. The nature of the bridging unit in linked metallocene plays adecisive role in the degree of interaction between the iron centers.47 Difunctional ferroceneswere employed in controlled polycondensation reactions to produce heteroannular polymerchains, e.g. poly(arylene-siloxane-ferrocene),48 and ferrocene containing polyesters.49

Ferrocene copolymers have been of interest for a variety of reasons including thermalstability, electroactivity, and non-linear optical properties in conjugated systems.

Fe

n

Fe

R

n

Chart 1.4 Poly(ferrocenylene) and poly(ferrocenes) are heteroannular polymers that incorporate bothcyclopentadienyl rings within the polymer backbone.

1.3 Ring-Opening Polymerization of Ferrocenophanes

[n] Metallocenophanes are molecules feature linking of the cyclopentadienyl rings bythe introduction of a heteroannular bridge (with n being the number of bridging atoms). Theyhave very interesting chemical properties compared to the simple metallocenes. A shortbridge causes distortion, which affects the reactivity of the molecule or the bridge cansterically hinder the approach of attacking species.


5

M Xα

β

Chart 1.5 Structure of metallocenophanes, with inclination, α, of the cyclopentadienyl rings (ring-tilt) anddeviation, β, of the exocyclic bonds from the planes of the rings.

Metallocenophanes can be strained, depending on the length of the bridge, and as aconsequence they can undergo ring-opening polymerization as cyclic compounds. The typicalstructure of metallocenophanes is illustrated in Chart 1.5; data related to the ring strain inferrocenophanes (M = Fe) is given in Table 1.1. Poly(ferrocenes) with relatively shortheteroannular bridges can be prepared by ring-opening polymerization of the correspondingferrocenophanes.50,51 To date, the Sn(t-Bu)2 bridged [1] ferrocenophane, has the smallest ringtilt α (14.1°) to still undergo ring-opening polymerization,52 and [2] metallocenophanes withtilt angles up to 13 ° were resistant to ring-opening polymerization.

Table 1.1 Structural differences indicating ring strain of [1] and [2] ferrocenophanes.53

Bridging Group (X) Ring Tilt α (°) Angle β (°)SiMe2 20.8 37.0GeMe2 19.0 36.8Sn(t-Bu)2 14.1 36.2(CMe2)2 23.2 11.0(SiMe2)2 4.19 10.8(SnMe2)2 0.7 10.5

1.3.1 POLY (FERROCENYLSILANES ) VIA RING OPENING POLYMERIZATION

As previously mentioned, strained metallocenophanes can undergo ring-openingpolymerization as cyclic compounds. Recently, the ring-openings polymerization of such astrained monomer, 1,1’-silylferrocenophanes, (Chart 1.6) was reported by Manners and co-workers.54,55 The polymerization can take place via a surprisingly large variety ofpolymerization techniques. In the solid state, the monomer can be polymerized by thermaltreatment or by γ-irradiation56 or in solution by transition metal catalysts57 and anionicinitiators.58 The polymer consists of alternating ferrocene and substituted silane units in thepolymer main chain.59 Cyclic voltammetry experiments showed interesting redox propertiesthat indicate intermetallic coupling between the iron centers.59a,60 Pyrolysis of the metalcontaining polymer produced a shaped magnetic ceramic.61

Chapter 1

6

Fe

Si

CH3

CH3

SiFeCH3

CH3

n

Chart 1.6 Ring-opening of silicon bridged ferrocenophane.

The sequential anionic polymerization allows one to synthesize block copolymerswith narrow polydispersities and targeted compositions.62 These block copolymers open upthe way to nanostructured inorganic/organometallic materials with unconventional properties.

1.4 Nanostructured Materials by Block Copolymer Self-Assembly

For the fabrication of regular arrays of nanoscale domains, two approaches have beenpursued. Lithographic methods, in a top-down approach, are developed for directedpatterning of substrates. Electron beam,63 focussed ion beam,64 and scanning probelithography65 are techniques that are used for patterning on a sub 100-nm scale.66 However,these techniques are very expensive and time-consuming as they fabricate the nanodomainsin a point by point addressing method. Bottom-up self-organization processes can be used toprepare regular patterns over a large area in a quick and cheap fashion.

Molecular self-assembly is the spontaneous organization of molecules into stable,well-defined structures.67 Multiple weak, reversible interactions like hydrophobic andhydrophilic effects, hydrogen bonds, Coulombic interactions, and van der Waals forcesaccount for the assembly of molecules into stable aggregates. Self-assembly is a strategy forfabrication at thermodynamic equilibrium, so defects are rejected and a high degree of ordercan be achieved.

Well-controlled architectures can be obtained in block copolymers and theequilibrium structures of interest can be controlled by designing the macromolecules.68 Theself-assembly is driven by the immiscibility of the two different phases. By forming asegregated phase, the number of contacts between the dissimilar phases is minimized, therebylowering the free energy of the system, which offsets the corresponding loss of entropy. Thedegree of chemical incompatibility, χ, between the different blocks, the degree ofpolymerization, N, of the polymer, and the composition, ƒ, can be adjusted to manufactureself-assembled structures with corresponding dimensions on a length scale of about ten to ahundred nanometers. Furthermore, the copolymers benefit from excellent processability,which makes them extremely interesting for nanotechnological applications.

Polymer technology can benefit from self-assembling interactions to formnanostructured materials combined with inorganic polymer chemistry to induce well-defined


7

patterns with topologically or chemically distinct regions that are unparalleled in organicpolymers. For example, the creation of structures with dimensions below 100 nm onsemiconductor substrates can be achieved. Controlled nanodomain formation and surfacesegregation of etch resistant layers, built from inorganic polymers, may prove to be usefultools.

1.5 Concept of this Thesis

The research described in this thesis comprises the synthesis, characterization andapplication of organometallic-inorganic homopolymers and block copolymers. Blockcopolymer phase separation was employed as a self-assembling strategy to obtainnanostructured materials.

Chapter 2 serves as a general introduction to the topics that are relevant for the workdescribed in this thesis. The crystallization and melting behavior of semi-crystallinehomopolymers is briefly discussed. Self-assembly in block copolymer systems is reviewed,and special attention is given to conformationally asymmetric block copolymers, blends ofblock copolymers and corresponding homopolymers, and block copolymer thin films.

In Chapter 3, the thermal behavior and morphology of ferrocenyldimethylsilanehomopolymers is described. A series of well-defined polymers was prepared by anionic ring-opening polymerization. The isothermal crystallization kinetics were analyzed by means ofthe Avrami equation. The polymers displayed a melting-recrystallization behavior uponheating after isothermal crystallization.

Chapter 4 covers the surface modification of poly(ferrocenylsilane) by means ofoxygen reactive ion etching (O2-RIE) treatment. It was found that the organometallic polymerwas highly resistant towards O2-RIE treatment. The formation of a thin iron-silicon-oxidelayer, of approximately 10 nm thickness, protected the underlying polymer towardsoxidation. The etching rate is then determined by a competition between the oxide formationand oxide removal by ion sputtering. Due to the presence of iron in the oxide layer thetransfer of an initial pattern into the underlying silicon or siliconnitride substrate by means ofCF4/O2-RIE treatment is possible.

Block copolymers of styrene and ferrocenylsilane (PS-b-PFS) are described inChapter 5. The bulk morphology of these organic-organometallic diblock copolymers wasstudied by transmission electron microscopy (TEM) and small angle x-ray scattering (SAXS).Blends of block copolymer with corresponding homopolymer were prepared to vary thecomposition and thereby control the morphology. Besides the three classic morphologies(spheres, cylinders, and lamellae) more exotic structures like the double gyroid and theperforated lamellae were observed. Information about order-order and order-disordertransitions was obtained from rheological experiments.

In Chapter 6, the thin film morphologies of PS-b-PFS copolymers, studied by atomicforce microscopy (AFM), transmission electron microscopy (TEM), secondary ion mass

Chapter 1

8

spectroscopy (SIMS), and x-ray reflectivity measurements are described. Thin films wereinvestigated both on solid silicon substrates as well as freestanding films. Upon annealing thefilms form islands or holes if the thickness is not commensurate with the domain spacing ofthe copolymer. However, these islands disappear upon longer annealing due to arearrangement of the phase separated nanodomains. The nanodomains could be visualized byAFM after selective removal of the organic PS phase using O2-RIE. In freestanding films,both parallel and perpendicular orientations for the cylinders were observed.

Chapter 7 describes thin films of asymmetric block copolymers of isoprene andferrocenylsilane (PI-b-PFS). Thin films were prepared and studied by AFM. Severalmorphologies were observed, depending on the block copolymer composition and filmthickness. The films show dewetting behavior if the initial film thickness is notcommensurate with the interdomain spacing. On the longer term, these films are unstable dueto crystallization of the poly(ferrocenyldimethylsilane) block.

The use of self-assembled organic-organometallic block copolymer films fornanolithographic applications is discussed in Chapter 8. Small Fe/Si/O domains can easily bedeposited on substrates like silicon or siliconnitride. The inorganic domains are also resistanttowards removal by a CF4/O2 plasma. This makes it possible to transfer the structure into theunderlying substrate by conventional reactive etching techniques and to increase the aspectratio of the nanodomains.

1.6 References and Notes

1 Annual meeting of the American Physical Society on December 29, 1959.2 M. Rothschild, J. A. Burns, S. G. Cann, A. R. Forte, C. L. Keast, R. R. Kunz, S. C.

Palmateer, J. H. C. Sedlacek, R. Uttaro, A. Grenville, D. Corliss, J. Vac. Sci. Technol.B 1996, 14, 4157.

3 J. P Silverman, J. Vac. Sci. Technol. B 1998, 16, 3137.4 C. W. Gwyn, R. Stulen, D. Sweeney, D. Attwood, J. Vac. Sci. Technol. B 1998, 16,

3142.5 S. T. Stanton, J. A. Liddle, W. K. Waskiewicz, A. E. Novembre, J. Vac. Sci. Technol.

B 1998, 16, 3197.6 G. Gross, R. Kaesmaier, H. Löschner, G. Stengl, J. Vac. Sci. Technol. B 1998, 16,

3150.7 In hot-embossing lithography, nanosized holes are physically printed in a thin

polymer layer with a mold: (a) S. Y. Chou, P. R. Krauss, L. Kong, J. Appl. Phys.1996, 79, 6101. (b) L. Kong, L. Zhuang, M. Li, B. Cui, S. Y. Chou, Jpn. J. Appl.Phys. 1998, 37, 5973.

8 (a) A. Kumar, H. A. Biebuyck, G. M. Whitesides, Langmuir 1994, 10, 1498. (b) J. L.Wilbure, A. Kumar, E. Kim, G. M. Whitesides, Adv. Mater. 1994, 6, 600. (c) A.Kumar, G. M. Whitesides, Science 1994, 263, 60. (d) J. Jackman, J. L. Wilbur, G. M.Whitesides, Science 1995, 269, 664.

9 M. H. Kryder, Thin Solid Films 1992, 216, 174.


9

10 S. Y. Chou, P. R. Krauss, L. Kong, J. Appl. Phys. 1996, 79, 6101.11 W. D. Volkmuth, T.. Duke, M. C. Wu, R. H. Austin, A. Szabo, Phys. Rev. Lett. 1994,

72, 2117.12 (a) L. F. Johnson, G. W. Kammlott, K. A. Ingersoll, Appl. Opt. 1978, 17, 1165. (b) E.

H. Anderson, C. M. Horwitz, H. I. Smith, Appl. Phys. Lett. 1983, 43, 874.13 M. Muthukumar, C. K. Ober, E. L. Thomas, Science 1997, 277, 1225.14 G. H. Frederickson, F. S. Bates, Annu. Rev. Mater. Sci. 1996, 26, 501.15 J. Ruokalainen, R. Makinen, M. Torkkeli, T. Makela, R. Serimaa, G. ten Brinke, O.

Ikkala, Science 1998, 280, 557.16 A. Van Blaaderen, R. Ruel, P. Wiltzius, Nature 1997, 385, 321.17 Q. Huo, D. I. Margolese, U. Ciesla, P. Feng, T. E. Gier, P. Sieger, R. Leon, P. M.

Petroff, F. Schüth, G. D. Stucky, Nature 1994, 368, 317.18 M. R. Ghadiri, J. R. Cranja, L. K. Buehler, Nature 1997, 369, 301.19 N. Kramer, H. Birk, J. Jorritsma, C. Schonenberg, Appl. Phys. Lett. 1995, 66, 1325.20 (a) G. Binning, C. F. Quate, Ch. Gerber, Phys. Rev. Lett. 1986, 56, 930. (b) J. A.

Dagata, Science 1995, 270, 1625.21 E. Betzig, K. Trautman, Science 1992, 257, 189.22 A. J. Bard, G. Denault, C. Lee, D. Mandler, D. O. Wipf, Acc. Chem. Res. 1990, 23,

357.23 R. S. Becker, J. A. Golovchenko, B. S. Swartzentruber, Nature 1987, 325, 419.24 To increase the writing speed, parallel array probe microscopes have been developed:

S. C. Minne, Ph. Fluekiger, H. T. Soh, C. F. Quate, J. Vac. Sci. Technol. B 1995, 13,1380.

25 “The Chemistry of Inorganic Homo- and Heterocycles”, I. Haiduc, D. B. Sowerby,Academic Press, Toronto, 1987.

26 The definition “inorganic polymers” is commonly quite confusing. Generally, ifpolymers contain metallic elements in their repeat unit they are regarded as inorganicpolymers.

27 M. Kumada, K. Tamao, Adv. Organometal. Chem. 1968, 6, 19. For a more completedescription: R. Baney, G. Chandra, In “Encyclopedia of Polymer Science andEngineering”, 2nd ed., John Wiley, New York, 1988, vol. 13, 312.

28 S. Gauthier, D. J. Worsfold, Macromolecules 1989, 22, 2214.29 A. R. Wolff, R. West, Appl. Organomet. Chem. 1987, 1, 7.30 See for Instance: (a) “Siloxane Polymers”, S. J. Clarson, J. A. Semlyen; Prentice Hall,

Englewood Cliffs, NJ, 1991. (b) “Inorganic Polymers”, J. E. Mark, H. R. Allcock, R.West; Prentice Hall, Englewood Cliffs, NJ, 1992.

31 H. R. Allcock, R. L. Kugel, J. Am. Chem. Soc. 1965, 87, 4216.32 P. Nguyen, P. Gomez-Elipe, I. Manners, Chem. Rev. 1999, 99, 1515.33 A. S. Hay, J. Org. Chem. 1960, 25, 1275.34 (a) Y. Fujikura, K. Sonogashiri, N. Hagihara, Chem Lett. 1975, 1067. (b) K.

Sonogashira, S. Takahashi, N. Hagihara, Macromolecules 1977, 10, 879. (c) N.Hagihara, K. Sonogashiri, S. Takahashi, Adv. Polym. Sci. 1981, 41, 149.

35 (a) S. J. Davies, B. F. G. Johnson, J. Lewis, P. Raithby, J. Organomet. Chem. 1991,414, C51. (b) B. F. G. Johnson, A. K. Kakkar, M. S. Khan, J. Lewis, P. Raithby, J.Organomet. Chem. 1991, 409, C12. (c) M. S. Khan, A. K. Kakkar, S. L. Ingham, P. R.Raithby, J. Lewis, B. Spencer, F. Wittman, R. H. Friend, J. Organomet. Chem. 1994,

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472, 247. (d) C. W. Faulkner, S. L. Ingham, M. S. Khan, J. Lewis, N. J. Long, P. R.Raithby, J. Organomet. Chem. 1994, 482, 139.

36 (a) S. J. Davies, B. F. G. Johnson, M. S. Khan, J. Lewis, J. Chem. Soc., Chem.Commun. 1991, 187. (b) M. S. Khan, S. J. Davies, A. K. Kakkar, D. Schwartz, B. Lin,B. F. G. Johnson, J. Lewis, J. Organomet. Chem. 1992, 424, 87.

37 (a) B. F. G. Johnson, A. K. Kakkar, M. S. Khan, J. Lewis, A. E. Dray, R. H. Friend, F.Wittman, J. Mater. Chem. 1991, 1, 485. (b) J. Lewis, M. S. Khan, A. K. Kakkar, B. F.G. Johnson, T. B. Marder, H. B. Fyfe, F. Wittman, R. H. Friend, A. E. Dray, J.Organomet. Chem. 1992, 425, 165.

38 Recently, a short essay on the discovery of ferrocene was published: P. Laszlo, R.Hoffmann, Angew. Chem. Int. Ed. 2000, 39, 123.

39 T. J. Kealey, P. L. Pauson, Nature 1951, 168, 1039.40 S. A. Miller, J. A. Tebboth, J. F. Tremaine, J. Chem. Soc. 1952, 632.41 (a) G. Wilkinson, M. Rosenblum, M. C. Whiting, R. B. Woodward, J. Am. Chem. Soc.

1952, 74, 2125. (b) G. Wilkinson, J. Am. Chem. Soc. 1952, 74, 6146.42 E. O. Fischer, W. Pfab, Z. Naturforsch.B 1952, 7, 378.43 “Metallocenes, An Introduction To Sandwich Complexes”, N. J. Long, Blackwell

Science, Oxford, 1998.44 (a) V. V. Korshak, S. L. Sosin, V. P. Alekseeva, Akad. Nauk SSSR 1960, 132, 360. (b)

V. V. Korshak, S. L. Sosin, V. P. Alekseeva, Vysokomol. Soedin. 1961, 3, 1332. (c) A.N. Nesmeyanov, V. V. Korshak, N. S. Voevodskii, S. L. Kochetkova, S. L. Sosin, R.B. Materikova, T. N. Bolotnikova, N. M. Bazhin, Dokl. Akad. Nauk SSSR 1961, 137,1370.

45 The most impressive early results (85% yield, molar mass ranging from 103 – 104

g/mol) on poly(ferrocenylenes) were obtained in 1979: see reference 46a.46 (a) E. W. Neuse, L. Bednarik, Macromolecules 1979, 12, 187. (b) T. Yamamoto, K.

Sanechika, A. Yamamoto, M. Katada, I. Motoyama, H. Sano, Inorg. Chim. Acta.1983, 73, 75.

47 S. Barlow, D. O’Hare, Chem. Rev. 1997, 97, 637.48 W. J. Patterson, S. MacManus, C. U. Pittman, J. Polym. Sci. Part A-1 1974, 12, 837.49 K. Gonsalves, L. Zhanru, M. D. Rausch, J. Am. Chem. Soc. 1984, 106, 3862.50 The most reported and studied system can be prepared by the ring-opening

polymerization of silicon bridged ferrocenophanes. See for example: (a) D. A.Foucher, B. Z. Tang, I. Manners, J. Am. Chem. Soc. 1992, 114, 6246. For a recentreview on poly(ferrocenysilanes): (b) I. Manners, Chem. Commun. 1999, 857.

51 Several poly(ferrocenes) with bridging groups, or atoms, other than silicon have beenprepared, as well. Phosphorus-bridge: (a) C. H. Honeyman, D. A. Foucher, F. Y.Dahmen, R. Rulkens, A. J. Lough, I. Manners, Organometallics 1995, 14, 5503. (b)T. J. Peckham, J. A. Massey, C. H. Honeyman, I. Manners, Macromolecules 1999, 32,2830. Sulfur-, selenium- and boron-bridges: (c) R. Rulkens, D. P. Gates, D. Balaishis,J. K. Pudelski, D. F. McIntosh, A. J. Lough, I. Manners, J. Am. Chem. Soc. 1997, 119,10976. (d) D. P. Gates, R. Rulkens, R. Dirk, P. Nguyen, J. K. Pudelski, R. Resendes,H. Braunschweig, I. Manners, Phosphor Sulfur Silicon Relat. Elem. 1997, 125, 561.Carbothiaferrocenophane: (e) R. Resendes, P. Nguyen, A. J. Lough, I. Manners,Chem. Commun. 1998, 1001. Hydrocarbon bridge: (f) J. M. Nelson, H. Rengel, I.Manners, J. Am. Chem. Soc. 1993, 115, 7035. (g) J. M. Nelson, P. Nguyen, R.


11

Petersen, H. Rengel, P. M. Macdonald, A. J. Lough, I. Manners, N. P. Raju, J. E.Greedan, S. Barlow, D. O’Hare, Chem.-Eur. J. 1997, 3, 573. Vinylene bridge: (h) M.A. Buretea, T. D. Tilley, Organometallics 1997, 16, 1507. Tin bridge: (i) R. Rulkens,A. J. Lough, I. Manners, Angew. Chem.-Int. Edit. Engl. 1996, 35, 1805. Germaniumbridge: (j) D. A. Foucher, M. Edwards, R. A. Burrow, A. J. Lough, I. Manners,Organometallics 1994, 13, 4959. For a review on ring opening of strainedmetallocenophanes see: (k) I. Manners, Can. J. Chem. 1998, 76, 371.

52 (a) R. Rulkens, A. J. Lough, I. Manners, Angew. Chem.-Int. Edit. Engl. 1996, 35,1805. (b) F. Jakle, R. Rulkens, G. Zech, D. A. Foucher, A. J. Lough, I. Manners,Chem.-Eur. J. 1998, 4, 2117.

53 “Metallocenes: an introduction to sandwich complexes”, N. J. Long, BlackwellScience, Oxford, 1998.

54 The synthesis of 1,1’-dimethylsilylferrocenophane was first reported in 1979: A. B.Fischer, J. B. Kinney, R. H. Staley, M. S. Wrighton, J. Am. Chem. Soc. 1979, 101,6501.

55 The first report on the ring-opening polymerization of 1,1’-dimethylsilyl-ferrocenophane was made 13 years after the monomer synthesis: D. A. Foucher, B. Z.Tang, I. Manners, J. Am. Chem. Soc. 1992, 114, 6246.

56 J. Rasburn, R. Petersen, T. Jahr, R. Rulkens, I. Manners, G. J. Vancso, Chem. Mater.1995, 7, 871.

57 (a) Y. Z. Ni, R. Rulkens, J. K. Pudelski, I. Manners, Macromol. Rappid Commun.1995, 16, 637.(b) P. Gomez-Elipe, R. Resendes, P. M. Macdonald, I. Manners, J. Am.Chem. Soc. 1998, 120, 8348.

58 (a) R. Rulkens, A. J. Lough, I. Manners, J. Am. Chem. Soc. 1994, 116, 797. (b) R.Rulkens, Y. Z. Ni, I. Manners, J. Am. Chem. Soc. 1994, 116, 12121. (c) Y. Z. Ni, R.Rulkens, I. Manners, J. Am. Chem. Soc. 1996, 118, 4102. (d) R. Rulkens, A. J. Lough,I. Manners, S. R. Lovelace, C. Grant, W. E. Geiger, J. Am. Chem. Soc. 1996, 118,12683.

59 Since the first the ring-opening polymerization of 1,1’-dimethylsilylferrocenophane in1992, numerous ferrocenylsilane polymers have been synthesized with differentsubstituents on the silicon atom and on the cyclopentadienyl-ring. For example, arylor longer alkyl chains on the silicon atom: (a) M. T. Nguyen, A. F. Diaz, V. V.Dementev, K. H. Pannell, Chem. Mater. 1993, 5, 1389. (b) D. A. Foucher, R.Ziembinski, B. Z. Tang, P. M. Macdonald, J. Massey, C. R. Jaeger, G. J. Vancso, I.Manners, Macromolecules 1993, 26, 2878. J. Rasburn, D. A. Foucher, W. F.Reynolds, I. Manners, G. J. Vancso, Chem. Commun. 1998, 843. Dihydrosilanebridged poly(ferrocene): J. K. Pudelski, R. Rulkens, D. A. Foucher, A. J. Lough, P.M. Macdonald, I. Manners, Macromolecules 1995, 28, 7301. Silicon substituted withalkoxy, aryloxy, and amino: (c) P. Nguyen, A. J. Lough, I. Manners, Macromol.Rappid Commun. 1997, 18, 953. (d) P. Nguyen, G. Stojcevic, K. Kubalka, M. J.Maclachlan, X. H. Liu, A. J. Lough, I. Manners, Macromolecules 1998, 31, 5977.Silicon substituted with chlorine: (e) D. L. Zechel, K. C. Hultzsch, R. Rulkens, D.Balaishis, Y. Z. Ni, J. K. Pudelski, A. J. Lough, I. Manners, Organometallics 1996,15, 1972. Methylated cyclopentadienyl ring: (f) J. K. Pudelski, D. A. Foucher, C. H.Honeyman, A. J. Lough, I. Manners, S. Barlow, D. O’Hare, Organometallics 1995,

Chapter 1

12

14, 2470. (g) J. K. Pudelski, D. A. Foucher, C. H. Honeyman, P. M. Macdonald, I.Manners, S. Barlow, D. O’Hare, Macromolecules 1996, 29, 1894.

60 R. Rulkens, A. J. Lough, I. Manners, S. R. Lovelace, C. Grant, W. E. Geiger, J. Am.Chem. Soc. 1996, 118, 12683.

61 M. J. MacLachlan, M. Ginzburg, N. Coombs, T. W. Coyle, N. P. Raju, J. E. Greedan,G. A. Ozin, I. Manners, Science 2000, 287, 1460.

62 (a) R. Rulkens, Y. Z. Ni, I. Manners, J. Am. Chem. Soc. 1994, 116, 12121. (b) Y. Z.Ni, R. Rulkens, I. Manners, J. Am. Chem. Soc. 1996, 118, 4102. (c) J. Massey, K. N.Power, I. Manners, M. A. Winnik, J. Am. Chem. Soc. 1998, 120, 9533. (d) J. Massey,K. N. Power, M. A. Winnik, I. Manners, Adv. Mater. 1998, 10, 1559.

63 (a) R. M. H. New, R. F. W. Pease, R. L. White, J. Vac. Sci. Technol. B 1994, 12,3196. (b) P. R. Krauss, S. Y. Chou, J. Vac. Sci. Technol. B 1998, 13, 2850. (c) R.O’Barr, S. Y. Yamamoto, S. Schultz, W. Xu, A. Scherer, J. Appl. Phys. 1997, 81,4730.

64 S. Matsui, Y. Ochia, Nanotechnology 1996, 7, 247.65 (a) K. Wilder, H. T. Soh, A. Atalar, C. F. Quate, J. Vac. Sci. Technol. B 1997, 15,

1811. (b) H. Sugimura, N. Nakagiri, Nanotechnology 1997, 8, A15. (c) M. Ishibashi,S. Heike, H. Kajiyama, Y. Wade, T. Hashizume, Appl. Phys. Lett. 1998, 72, 1581.

66 For a comparison between electron beam and scanning probe lithography, see: K.Wilder, C. F. Quate, B. Singh, D. F. Kyser, J. Vac. Sci. Technol. B 1998, 16, 3864.

67 G. M. Whitesides, J. P. Mathias, C. T. Seto, Science 1991, 254, 1312.68 (a) F. S. Bates, G. H. Frederickson, Annu. Rev. Phys. Chem. 1990, 41, 525. (b) F. S.

Bates, J. H. Rosedale, G. H. Frederickson, J. Chem. Phys. 1990, 92, 6255.

Homopolymer and Block Copolymer Assemblies

13

&KDSWHU��


This chapter serves as an introduction to topics that are relevant for the workdescribed in the remainder of this thesis. The crystallization and meltingbehavior of semi-crystalline homopolymers is briefly discussed. Self-assembly in block copolymer systems is reviewed, and special attention isgiven to conformationally asymmetric block copolymers, blends of blockcopolymers and homopolymers, and block copolymer thin films.

2.1 Semi-Crystalline Homopolymers

Figure 2.1 Three proposed models for the lamellar polymer crystals. (a) regular, adjacent reentry folds similarto those postulated as present in pyramidal crystals that have been grown from solution. (b) irregular, adjacentreentry folds in which the extent or thickness of the irregular layer is suggested to be proportional to thetemperature. (c) switchboard, or non-adjacent reentry, model in which an even more non-ordered amorphouslayer is present on both sides of the lamellae than in the regular model.1

Chapter 2

14

Since the discovery of stereospecific polymerization by Ziegler and Natta, thecrystallization of linear polymers has been studied extensively. In 1954 Keller, Fischer, andTill found independently that linear polyethylene crystallizes from dilute solutions into thinplatelet crystals, the so-called lamellae. These crystalline lamellae consist of arrays of foldedchains (Figure 2.1). The reentry of the folded chains can be either adjacent or nonadjacent.The thickness of lamellae is typically 10 to 20 nanometer (depending on the crystallizationtemperature), but can extend laterally over micrometers.

2.1.1 MORPHOLOGY

Semi-crystalline polymers exhibit a structural hierarchy depending on the lengthscale. The monomers are packed in crystal unit cells (nanometer level order). Lamellaeconsist of regularly packed unit cells whereas one chain belongs to a large number of cells,depending on the chain length. The lamellae in most polymers, crystallized either from themelt or from concentrated solutions, are organized in larger spherical-structures, calledspherulites. Starting from a nucleus, the lamellae grow in a direction perpendicular to thechain axis. When the lamellae start curving and splaying apart, so-called axialites2 orhedrites3 are formed (Figure 2.2), which eventually mature into spherulites.4 The lamellaegrow in the radial direction from the spherulite until individual spherulites will impinge uponone another and the primary crystallization is terminated.

Figure 2.2 AFM image (deflection) of an edge-on view of a melt-crystallized hedrite of poly(ferrocenylsilane).5


15

2.1.2 CRYSTAL MELTING TEMPERATURE

The equilibrium melting temperature is perhaps the most important macroscopiccharacteristic for a given macromolecular crystal. In general, the observed meltingtemperature Tm is always lower than the equilibrium melting temperature Tm

0. Experimentallyequilibrium is difficult to reach since crystals with large dimensions are difficult to grow. Thefinite crystal size, which depends also on the crystallization conditions, as well as onimpurities, contributes to melting point depression. The Thomson-Gibbs equation can beemployed to extrapolate the melting temperature of finite crystals Tm to the equilibriummelting temperature Tm

0:

∆⋅

−=f

mm HlTT

γ210

where Tm0 is the melting temperature of the infinitely large crystal, γ is the specific surface

free energy of the fold surface, l is the lamellar thickness, and ∆Hf corresponds to the bulkheat of fusion. As the lamellar thickness l decreases (i.e. the crystals get thinner), the meltingpoint decreases. Assuming that during isothermal crystallization the fold length is fixed at avalue proportional to the reciprocal of the supercooling ∆T, one can derive that:6

( )β

β2

120cm

m

TTT

+−=

where Tm is the observed melting temperature, Tm0 is the extrapolated equilibrium melting

temperature, Tc is the crystallization temperature, and β indicates the fold length. A plot ofthe experimental melting temperatures versus the crystallization temperatures (the so-calledHoffman-Weeks plot) can be extrapolated to a point where:7

0mcm TTT ==

2.1.3 CRYSTALLIZATION KINETICS

For a given polymer, the rate of crystallization depends on the degree ofundercooling. The crystallization rate exhibits a maximum as a function of temperature. Atthe melting point the crystallization rate is zero. As the undercooling and thus the drivingforce for crystallization increases, the crystallization rate increases. As the glass temperatureof the polymer is approached, the crystallization rate becomes smaller again due to theabsence of large scale chain mobility. The rate of crystallization can be determined at a giventemperature, and the process may be analyzed by the Avrami equation:8

Chapter 2

16

( )nt ttKX )(exp1 0−−−=

where Xt is the relative conversion of the crystallization process at time t, K is a temperaturedependent growth-rate parameter, t0 is the point in time at which crystallization begins, and nis a temperature-independent nucleation index, the so-called Avrami exponent. The Avramiexponent characterizes the type of growing structure and nucleation mechanism.

2.2 Block Copolymers

Block copolymers are macromolecular architectures that consist of two or morechemically different polymer chains, linked to one another covalently. They find applicationsfor instance as compatibilizers,9,10,11 impact modifiers or as thermoplastic elastomers.12 Insolution their properties are exploited for applications in foams, solubilizers, thickeners andas dispersion agents.13

Figure 2.3 Typical block copolymer architectures.

The scientific interest in block copolymers has been tremendous during the last fewdecades and is still a field of major research efforts. Structural and dynamical properties havebeen considered for melts, solids, dilute solutions and concentrated solutions. In the melt,block copolymers self-assemble into a variety of structures via the process of microphaseseparation. The important energies in the formation of phase separated structures are theentropic chain-conformation energy and the enthalpic interaction energy (characterized by theFlory-Huggins interaction parameter χ). The chemical connectivity of the demixing blocksprevents macrophase separation and therefore domains of both components are created. Thedomain morphology is prescribed by the area-minimization of the surface separating thesedomains. Consequently, these so-called inter material dividing surfaces (IMDS) are said tohave constant mean curvature.


17

The self-assembly of block copolymers provides a versatile means to createnanostructures with potential applications in biomaterials, optics, and microelectronics.14 Inbiomineralization, for instance, complex morphologies on different length scales are usuallyobtained through cooperative self-assembly of organic and inorganic species.15

Well-defined block copolymers can be prepared by controlled/living polymerizationvia several methods: sequential monomer addition, coupling reaction of living polymerchains, and mechanism transformation.16 The sequential monomer addition is restricted byrelative monomer reactivity.17 First a block is polymerized and after completion, the livingpolymer chain functions as the initiator for the next monomer. If initiation is fast compared tothe propagation, well-defined block copolymers with relatively narrow molar massdistributions can be obtained.

2.2.1 BULK MORPHOLOGY

In a diblock copolymer phase diagram the order-disorder transitions (ODT) and theorder-order transitions (OOT) between different morphologies are plotted in Figure 2.6.Typically, a phase diagram describes the morphology as a function of temperature andcomposition. It has been established that the temperature, the composition, the chemicalnature of the phases, and the molar mass influence the morphology.18 These quantities can bedescribed by the Flory-Huggins interaction parameter χ, the number of statistical segments inthe block copolymer N (which is proportional to its molar mass), and the copolymercomposition f. The χ parameter accounts for the temperature and the chemical nature of thecopolymer. In a diblock copolymer phase diagram, the product χN is plotted as a function ofcomposition f, since theories predict that such phase diagrams would be universal.19,20,21 Theproduct χN can be regarded as a measure for the degree of segregation. The strongsegregation limit (SSL) extents to very high values of χN, whereas the weak segregation limit(WSL) is reached near the order-disorder transitions. Typically, the domain interfaces aresharp in the strong segregation limit and more diffuse in the weak segregation limit.

Figure 2.4 Three classical morphologies in block copolymer systems. With increasing volume fraction of thedarker phase, spheres on a bcc lattice, hexagonally packed cylinders, alternating lamellae and the inverse phasescan be observed.

Chapter 2

18

The morphology of microdomains formed by pure and simple diblock copolymers hasbeen an intensively researched and is by now well understood area.22 In neat diblocks, three“classic” ordered microphases are usually distinguished (Figure 2.4). These includealternating lamellae (LAM), hexagonally packed cylinders (HEX) and body-centered-cubicpacked spheres (BCC). In addition, some other more complex microstructures, like thebicontinuous gyroid (see Figure 2.5)23,24,25, the hexagonally modulated layer,26 or theperforated lamellae,26,27,28 may appear especially near the order-disorder transition.29,30

Figure 2.5 Illustration of the double gyroid structure. A triply periodic interface with constant mean curvature.31

In the strong segregation limit (SSL), i.e. at large values for χN, the volume fractionsfor the transitions between the different ordered phases are almost independent of χN. Forexample, for PS-PI diblocks, the transitions were situated at the following PS volumefractions: BCC < 0.17 < HEX < 0.28 < DG < LAM < 0.62 < DG < 0.66 < HEX <0.77 <BCC.22a

For relatively high values of χN, the classical phases remain stable as opposed to themore exotic ones. Theoretical extrapolations to χN = ∞ provide evidence for the order-ordertransition from lamellar to cylindrical at a relative volume fraction of 0.310 and from cylinderto sphere at a volume fraction of 0.105.32 In the strong segregation limit (SSL) the phaseboundaries are almost independent of χN.

Triblock copolymers composed of three different phases (ABC) form even morecomplex structures,33 such as helical strands surrounding cylinders embedded in a continuousmatrix,34 or the so-called “knitting pattern”.35 Architecturally more complex nonlinear blockcopolymers, such as star blocks and graft copolymers, also display a very diversemorphological behavior.36


19

2.2.2 IDENTIFICATION OF BLOCK COPOLYMER BULK MORPHOLOGIES

A few techniques have been used to study the bulk morphology of block copolymers.Transmission electron microscopy (TEM) is a widely used tool to visualize the phaseseparated nanodomains.37 An ultrathin sample slice (approximately 50 – 70 nm) ismicrotomed and projected in an electron microscope. The contrast in TEM imaging ofdiblock copolymers is due to differences in electron scattering power of the phases. Mostpolymers have low atomic number atoms (organic polymers) and therefore the scatteringcontrast is usually weak. In general, selective staining is necessary to obtain suitable contrastin the TEM.38,39 Care must be taken concerning the identification of the morphology by TEMexperiments alone, since it is based only on the 2D projection of a 3D structure of a smallregion of the sample.40

Figure 2.6 Phase diagram for conformationally symmetric diblock melts. Phases are labeled L (lamellar), H

(hexagonal cylinders), Q (cubic, bicontinuous gyroid Ia3d and bcc spheres Im3m), CPS (close-packedspheres), and DIS (disordered). Figure was adopted from ref. 21.

X-ray diffraction identifies the lattice type and lattice parameters by their specificBragg-reflections. Diffraction is governed by the Bragg law:

θλ sin2 hkld=

where λ is the wavelength of the x-rays, dhkl is the spacing between the allowed reflectinghkl-planes, where hkl are the Miller indices, and θ is half of the scattering angle 2θ.

Chapter 2

20

Typically, the x-ray wavelength is about 1.5 Å and the first reflecting hkl-planes are spacedby approximately 30 nanometers. This implies that the scattering angles must be detectedfrom 2θ ~ 0.3º upwards, by so-called small-angle x-ray scattering (SAXS). SAXS data arerepresented by plots of scattered intensity as a function of the scattering vector qhkl:

λθππ sin42 ==

hklhkl d

q

Each lattice type has a characteristic sequence of distances between the scattering planes.41

For a lamellar structure, the interdomain distance alam is equal to the d100 spacing, and thefollowing reflecting dh00 spacings are given by (h = 1, 2, 3, 4,...):

h

ad lam

h =00

This results in peak positions, relative to the first bragg-peak q100, at:

etc. 4 3, 2, 1,100

00 =q

qh

For a morphology that consists of cylinders that are packed in a hexagonal lattice, theinterdomain distance ahex, and the reflecting dhk0 spacings are connected by:

( )22

0

3

4khkh

ad hex

hk

++=

In this case the hk0-planes diffract for hk = 10, 11, 20, 21, 30,..., so the corresponding relativepeak positions will be at:

etc. ,9 ,7 ,4 ,3 1,100

0 =q

qhk

The simple cubic and body-centered cubic lattices require seven well-developed peaks to bedistinguished. A cubic lattice type has interdomain distance acubic and dhkl spacings:

( )222 lkh

ad cubic

hkl++

=

A bcc lattice type has allowed reflections for hkl = 110, 200, 211, 220, 310, 222, 321, 400,...and will therefore display diffraction peaks at relative positions of:


21

7,6,5,4,3,2,1100

=q

qhkl

For a simple cubic lattice, the allowed indices are hkl are 100, 110, 111, 200, 210, 211, 220,300,... Therefore, the √7 reflection allows one to distinguish between the body-centered cubicand the simple cubic structure, since for the simple cubic structure, the √7 reflection is absent.

For the more exotic double gyroid structure, the space group was determined to be Ia3d (Int.tables no. 230).41 The first three allowed reflections are 211, 220, and 321, which result inrelative peak positions at √3, √4, and √10.23

From the specific peak reflections it is possible to accurately measure the dimensionsof the structure. Combining the structure and dimensions that are known from SAXSmeasurements, together with the block copolymer composition, an accurate determination ofthe individual domain sizes can be calculated. Due to the form factor of a single domain, theintensity will display minimal values at corresponding positions (so-called particle scatteringprofile).42 More specifically, the scattering intensity displays a minimum at qR equals 5.76for spheres and 3.83 for cylinders respectively, where q is the scattering vector and R is theradius of the sphere or cylinder.43

Finally, rheological measurements are often employed to observe order-order andorder-disorder transitions.44 The flow properties of block copolymer melts depend on thestate of order in the system. The order-disorder transition temperature (TODT) can beidentified from rheological measurements during a heating ramp. Going through TODT, thedynamic elastic modulus displays a sharp decrease. Also, the frequency dependence of theelastic and storage moduli gives valuable information concerning transitions. In the phase-separated state, i.e. below TODT, the elastic modulus displays a somewhat limited (non-Newtonian) behavior (G’~ω0.5). Rheological experiments are also very useful to locate order-order transitions, although they do not identify phases unambiguously. Some reports in theliterature, however, succeeded to assign a certain rheological “fingerprint” to definedphases.28,29,45

2.2.3 CONFORMATIONAL ASYMMETRY

Most theories on microphase separated morphologies of diblock copolymers deal withconformationally symmetric diblocks. Conformational asymmetry can arise from eitherdifferences in densities, or in Kuhn lengths, or both. The asymmetry ratio ε, whichcharacterizes the conformational asymmetry of a system, can be defined by:46

2AA0

2BB0

2B,

2A,

BA

b

b

RR

ff

gg ρρ

ε ==

Chapter 2

22

In this equation f is the volume fraction, Rg is the radius of gyration, ρ0 is the densityof the pure component and b denotes the Kuhn segment length. Vavasour and Whitmore,47

and later Matsen and Schick,48 calculated the effect of conformational asymmetry on themicrophase diagram of diblock copolymers (see Figure 2.7). They found for aconformationally asymmetric diblock that the order-disorder transition was virtuallyunchanged with respect to a conformationally symmetric diblock, but all the order-orderboundaries were shifted toward the higher A content, in which A is the block with the largerρ0b

2.

Figure 2.7 Phase diagram of a conformationally asymmetric diblock copolymer. Phases are labeled fordisordered (D), lamellar (L), gyroid (G), hexagonally packed cylinders (H), and spheres on a cubic lattice (C).The figure was adopted from ref. 48.

The most investigated diblock system, poly(styrene-block-isoprene), displays a certainasymmetry of the phase diagram with respect to the composition. Hasegawa et al.49 reportedthe order-order transitions of PS-PI diblock copolymers for BCC - HEX - LAM - Tetrapod -HEX - BCC to be located at PS volume fractions of approximately 0.18, 0.31, 0.61, 0.66 and0.76, respectively. The order-order transitions were found to be asymmetric with respect tothe PS volume fraction. The shift of the order-order transitions was explained byconformational asymmetry of the PS-PI diblocks. The authors considered that the tetrapod-network structure and the ordered bicontinuous double diamond (OBDD, space group Pn3m)were essentially the same type of morphology. It must be noted that some structuresidentified as OBDD may in fact have been double gyroid (DG) morphologies (space groupIa 3d).23,25

Pochan et al.50 investigated a series of conformationally asymmetric poly(isoprene-b-tert-butylmethacrylate) (PI-PtBMA) diblocks (ε ~ 0.75) in the strong segregation limit and


23

compared their results with the theory. They found that the shift of phase boundaries tohigher volume fractions of PtBMA is consistent with the theory. However, they observedPtBMA spheres up to a relatively large volume fraction of PtBMA compared to the theory.

2.2.4 HOMOPOLYMER /BLOCK COPOLYMER BLENDS

Another approach to explore the composition dependence of the morphology diagramis to alter the composition by mixing homopolymer (hA) into the diblock copolymer (AB).51

Two different situations may occur when a homopolymer is blended into a diblockcopolymer. The components either can mix with preservation of the microdomains, or theycan demix leading to macrophase separation. If they mix, the addition of homopolymereffectively swells the matching block of the copolymer and this may induce a phasetransition. The solubilization of the homopolymer in a block copolymer depends on therelative molar mass α of the homopolymer to the molar mass of the corresponding block inthe copolymer (α = MA,homo / MA,diblock). Both theory52 and experiment53 have shown that thisratio α is a crucial parameter for determining the topology of the phase diagram. If 1 > α >0.5 the homopolymers swell the corresponding domains and are weakly segregated towardsthe center of these domains. This case corresponds to the “dry brush” situation, where a freechain cannot diffuse into domains of shorter chains grafted to a surface (here the interfacebetween the blocks), since the loss of conformational entropy is larger than the gain intranslational or mixing entropy.54 However, if α < 0.5 the homopolymers are much shorterthan their matching block and they distribute themselves more uniformly throughout thedomains (“wet brush” situation). Consequently, the homopolymer solubility for low α ismuch higher compared to high α.

Figure 2.8 Schematic illustration of homopolymer distribution within a phase separated block copolymer forrelatively high (but still smaller than approximately 1) and low values of α (<0.5). Both situations lead toselective swelling of the corresponding microdomain.

Extensive work on binary blends of polystyrene homopolymers with styrene-isoprenediblocks or styrene-butadiene diblocks demonstrated the importance of the relative

Chapter 2

24

homopolymer molar mass and overall blend composition. The overall composition of theblend mainly determines the final morphology and approximately matches the morphology inneat diblocks at the same composition.53f-i

An important aspect about diblock/homopolymer blends found both in theory andexperiment is that the addition of homopolymer can stabilize new morphologies especiallybetween the lamellar and cylindrical phases. This is attributed to the relief of packingfrustration by the homopolymer. In neat diblocks, the junctions are placed on the interfaceand pull the monomers near these junctions, thereby stretching them. The addedhomopolymer will preferentially occupy regions in the domains, which would otherwise beoccupied by stretched chains of the corresponding block.

2.2.5 BLOCK COPOLYMER THIN FILMS

In thin films, the presence of a substrate and surface can induce orientation of thestructure,55 and can result in changes in domain dimensions or in phase transitions56,57 due topreferential segregation of one of the blocks at the substrate or the surface.

If one block has a higher affinity for the substrate, it will exhibit preferential wetting,resulting in an orientation of the phase separated domains parallel to the substrate.58 Similarwetting will occur at the free surface, i.e. the block with the lower surface free energy willenrich the surface.59 For symmetric block copolymers in unconfined films, the lamellae orientparallel to the surface and substrate and a mismatch between the film thickness and lamellarperiod results in the formation of holes or islands in thin films (the term “unconfined” wasused to indicate situations when the film has at least one free surface).60 This allows thepreferred blocks to be present at both the substrate and free surface, without changing theequilibrium period of the lamellae (Figure 2.9).

Figure 2.9 Schematic illustration of a symmetric block copolymer thin film. The same block covers both thesubstrate and the surface. By the formation of islands or holes, the preferred block can still be at the substrateand surface although the thickness is not commensurate with the equilibrium spacing d.


25

If the same block is present at both the substrate and the surface, the film thicknessmust match nL0, and correspondingly, if opposite blocks are present at the substrate andsurface, the film thickness is quantified as (n+½)L0 (where n is an integer and L0 is the bulkequilibrium period).61 This is for example the case for thin PS-b-PMMA copolymer films onsilicon substrates. PS will segregate at the film surface, due to its lower surface energy,whereas PMMA has a higher affinity for the hydrophilic silicon substrate. If the initialthickness of the film differs from the quantified thickness, islands or holes will form, asillustrated in Figure 2.9. The lateral size and number of the islands depend on the initial filmthickness and annealing time. Such surface relief structures have been studied in detail.62

By confinement, holes and islands can not be formed. The lamellar period in confinedthin block copolymer films was found to deviate from the equilibrium bulk period as afunction of film thickness.63,64 There is in some cases, however, an extent to which the periodcan increase or decrease to accommodate the thickness constraint. The frustration generatedwithin the film can lead to perpendicular orientation of lamellae.65,66 If the film thickness isbelow the bulk equilibrium period, lateral structures are sometimes found even in unconfinedfilms.56,67,68 Perpendicular orientation of lamellae can also be obtained by balancing allinterfacial interactions, i.e. by making the substrate neutral to both blocks.69 The substrate canbe made neutral to both blocks by surface grafting of a random copolymer with the samecomposition as the block copolymer. The domains in films that were confined between twonon-preferential surfaces were found to orient perpendicular to the substrate throughout theentire film.70 In such structures, both blocks are present at the polymer-substrate interface.Theoretical predictions71,72,73 and simulations74 have confirmed such observations.

The ordering in thin films of asymmetric block copolymers, with spherical orcylindrical morphologies, has been investigated to a much smaller extent compared tosymmetric block copolymers. In a theoretical study, the surface induced ordering of acylindrical phase was treated.75 A lamellar region (covering layer) at the surface existed if thesurface tension difference between the two phases and the surface is above a certain criticalvalue. Suh et al.76 found theoretically that cylinders in triblock and diblock copolymers couldorient parallel or perpendicular in confined geometries, depending on the film thickness andsurface tensions. In films of asymmetric poly(styrene-block-butadiene) copolymers (PS-b-PB) annealed without the presence of a substrate, a transition from the bulk cylindricalmorphology was observed as the film thickness decreased.77 A block copolymer that formedPS cylinders transformed into an interlayer, penetrated by PB channels as the film thicknessdecreased. A block with a PB cylindrical bulk structure formed PB spheres and thin PBsurface wetting layers in very thin films. This was induced by the surface segregation of thelower surface tension PB component. For different PS-b-PB diblock compositions, themorphology in thin film droplets of non-uniform thickness on carbon substrates was studiedby Henkee et al.78 Although there was no clear affinity for one of the two blocks to wet thesubstrate, a parallel orientation was observed for spherical and cylindrical domains in a singledomain layer. For a spherical morphology, the domains oriented in a two dimensionalhexagonal packing. Adjacent layers consisted of hexagonally packed arrays of spheres werestacked onto each other in a fashion so that the spheres in the upper layer are located over the

Chapter 2

26

interstitial sites of the lower layer. Cylindrical domains arranged themselves also parallel tothe surface, with transversely shifted adjacent layers. The morphology of spherical andcylindrical PS-b-PB and PS-b-PI copolymer thin films on silicon wafers was investigated byHarrison et al.79 The authors combined reactive ion etching (RIE) and high resolutionscanning electron microscopy (HRSEM) to visualize thin block copolymer films at variousdepths. In a single domain layer, the cylinders were oriented parallel to the substrate, whichwas wetted by a layer of PB. The wetting of the silicon by the PB was ascribed to the lowerinterfacial tension between unsaturated hydrocarbons and a polar substrate, compared toaromatic hydrocarbons and a polar substrate.80 In a monolayer of spheres, the domains werearranged in a hexagonal lattice, also parallel to the substrate. Both the spherical andcylindrical films formed islands and holes, if the initial film thickness was not compatiblewith the domain spacing. Along the edges of islands and holes in a cylindrical film theauthors observed hexagonally packed spheres. In these regions, that are slightly thinner thanthe monolayer thickness, the PB wetting layers were depleting the PB in the interior of thefilm and this drove the morphology to an effectively lower PB concentration, producing PBspheres. Liu et al.81 studied the morphology of both cylindrical di- and triblock copolymers,consisting of deuterated poly(styrene) (dPS) and poly(vinylpyridine) (PVP). Uniform filmswere prepared by spin-casting from toluene solution onto native oxide covered silicon wafers.Due to the strong interaction between the substrate and the PVP block, the substrate waspreferentially wetted by the PVP block, even though it was present in the minority phase. Theauthors concluded that also asymmetric diblock copolymers may orient parallel to a surface.The formation of islands was not observed for the diblock copolymer films even afterannealing for five days in contrast to the triblock copolymer films. Yokoyama et al.82 studiedthin films of asymmetric PS-b-PVP copolymers and found that the minority PVP phasepreferentially absorbed at the silicon substrate. This block copolymer displayed a sphericalmorphology in the bulk and the (110) planes were oriented parallel to the substrate in thinfilms, up to 5 layers from the substrate. Island and hole formation was also observed in theirfilms, indicating a high degree of ordering in the direction perpendicular to the substrate. Theordering in asymmetric diblock copolymer films, consisting of poly(ethylenepropylene-block-ethylethylene) (PEP-b-PEE) was investigated by using neutron reflectivity and phasecontrast microscopy.83 The PEE minority phase segregated at both the substrate and thesurface. Island formation was not observed for any thickness. The lattice distortions in thebulk of these films and the formation of a wetting surface layer did not favor the creation ofislands and holes.

Thin block copolymer films have been studied using several techniques. Visualizationtechniques include atomic force microscopy (AFM), scanning electron microscopy (SEM),84

and transmission electron microscopy (TEM). Atomic force microscopy is particularly usefulto study block copolymer thin films since only minimal sample preparation is required.85 Thereduction of lateral forces can be accomplished by using tapping mode, so even soft polymerfilms can be imaged. In tapping mode, the cantilever is resonated near its resonancefrequency. When the tip is scanning the surface, the amplitude can be kept constant. Thephase shift of the cantilever’s frequency is used to obtain a so-called phase image.86 The


27

contrast in phase images is due to mechanical interaction by the different phases and forvisualization of the phase separation in thin copolymer films, phase images often give goodcontrast.87,88 The disadvantage of performing atomic force microscopy on block copolymersamples is that only information about the surface structure can be obtained.

Secondary Ion Mass Spectrometry (SIMS) is a surface characterization technique thatgives in-depth information about the atomic composition in thin films. While an ion gun, e.g.Ar+ ion gun, is sputtering and ionizing particles of the sample upon surface bombardment, theenergy of the incident ions is dissipated in a series of collisions and secondary ions from thesample are emitted. A mass spectrometer detects the ions that have been removed.

Thin film layer systems are often investigated by x-ray or neutron specularreflectivity. The reflected intensity is determined by the thickness dj, the interfacial roughnessσj and the density ρj of the layer denoted by j.89 For x-rays the ρj densities correspond to theelectron densities and for neutrons to the scattering length density of the nuclei, respectively.An arbitrary density profile ρ(z) perpendicular to the sample surface will result in a reflectedintensity I(qz) given by the Born approximation:90

( ) ( ) ( )2

4exp

1dzziq

dz

zd

qqI z

z

z −≈ ∫ρ

where qz = (4π sinθ)/λ is the wave vector transfer, λ is the wavelength of the radiation and θcorresponds to the incident angle with respect to the surface. The calculated scattering spectracan be compared with the experimental data, and after optimization a density depth profilecan be obtained. Very recently, a new Fourier method has been developed for the analysis ofx-ray reflectivity data from low contrast polymer systems.91 This technique was used toprofile the system of interest in this thesis as will be described later.

2.3 Surface Treatment of Polymers

Plasma processes have become important industrial processes in modifying polymersurfaces. Phenomena such as wetting, adhesion, and thin film stability are important factorsfor their applications. Therefore, surface treatment of polymers and substrates is ofsubstantial interest. Generally, a plasma can be defined as a gas containing a variety ofspecies including: electrons, positive ions, negative ions, radicals, atoms, and othermolecules. To produce a plasma, an energy source is required to produce ionization of thefeed gas.

Plasma etching can proceed by physical sputtering or chemical reaction and ion-assisted mechanisms (Figure 2.10).92 Chemical etching refers to the chemical reactions thattake place between the active plasma species and the surface of interest. After absorption ofthe reactive species, the reaction takes place and the product desorbs from the surface. Purely

Chapter 2

28

chemical etching has therefore no preferential direction, which leads to isotropic etching. Ionenhanced plasma etching, on the other hand, produces an anisotropic profile. An electricalfield makes the ions in the plasma strike the surface almost at normal incidence. The ionbombardment accelerates the reaction in the direction of incidence so that etching occursanisotropic (Ion-enhanced energetic). A second class of ion enhanced plasma etching requiresan etchant and an inhibitor (Ion-enhanced protective). The inhibitor forms a thin protectivelayer on the surface. Chemical etching can then only occur at places that have been disruptedby the bombarding ions. Since vertical surfaces see little or no ion bombardment the ion-enhanced etching results in anisotropic profiles. Physical etching occurs due to thebombardment of positive ions. The positive ions are accelerated towards the surface, andbreak bonds upon impact, which results in physical etching. Since this sputter etching ispurely physical it is least selective.

Figure 2.10 Basic mechanisms of plasma etching: 1. Sputtering - ion bombardment physically removesmaterial. 2. Chemical – neutral radicals react with the surface and a volatile reaction product is formed. 3. Ion-enhanced energetic – impinging ions increase the reactivity of the surface so chemical reactions can take placeto form volatile products. 4. Ion-enhanced protective – a film forms on the surface that acts as a barrier to theetchant. The ion bombardment disrupts the protective film and allows the chemical etching. Figure was obtainedfrom ref. 92.

Which plasma etching mechanisms take place depends on a large number ofvariables, such as plasma gas, pressure, reactor design, and substrate.93 The most importantparameters in etching are etch rate, selectivity, and anisotropy.94 Selectivity is defined as theetch rate ratio between the material which is to be etched and the material which is not to beremoved (e.g. the masking layer) and anisotropy is related to the directionality of etching.


29

Polymer resists in the lithographic industry are being developed for smaller andsmaller feature sizes. Organic polymers are known to oxidize rapidly while silicon containingpolymers form an etch resistant oxide layer when exposed to an oxygen plasma.95 Thecompetition between the oxide formation and the oxide removal by ion sputtering results inrelatively low etching rates compared to organic polymers. Selective silylation of exposedareas in polymer resist films will therefore be developable in oxygen plasmas.96 Also otherinorganic components, such as tin,97,98 germanium,99 or titanium,100 have been employed toenhance the etch resistance in oxygen plasmas. Another approach combines etch resistance ofinorganic or organometallic polymers with the sensitivity of organic polymers in so-calledmultilevel resists. 101,102 Organometallic materials have been employed for direct writingmethods in lithographic applications, as well. Direct writing of metallic features is possibleby irradiating thin films of organometallic materials with highly focussed charged particlebeams.103

X-ray Photoelectron Spectroscopy (XPS) is a powerful technique for studying thechemical composition of surfaces. A sample surface is irradiated with x-ray photons thatinteract with an inner shell electron of an atom. The energy from the photon is thencompletely transferred to the electron. It the transferred energy is larger than the bindingenergy of the electron, the electron will be ejected from the atom. Only electrons that are notreabsorbed in the material and escape the sample are detected (typically less than 10 nmdepth, depending on the angle of take off). The kinetic energy of the photoelectrons (Ek) thatare emitted from the sample surface is determined by the difference between the x-ray photonenergy (hν) and the binding energy of the emitted inner shell electron (Eb). The bindingenergy of a particular shell of an atom is unique for each element.

kb EhE −= ν

Electrons from all orbitals with a binding energy less than the x-ray photon energycan be excited. In a so-called survey spectrum, the photoelectron binding energies can readilybe identified. A high resolution scan even allows one to detect small shifts in these bindingenergies and can therefore give information about the chemical environment of the atom.


1 P. H. Geil, Chem. Eng. News 1965, August 16, 72.2 D. C. Bassett, A. Keller, S. Mitsuhashi, J. Polym. Sci. 1963, A1, 763.3 P. H. Geil, in ”Growth and Perfection of Crystals”, R. H. Doremus, B. W. Roberts,

D. Turnbull, Wiley, New York, 1958, 579.4 H. D. Keith, J. Polym. Sci. 1964, A2, 4339.5 See also Chapter 3 of this Thesis. The figure was taken from: R. G. H. Lammertink,

M. A. Hempenius, I. Manners, G. J. Vancso, Macromolecules 1998, 31, 795.

Chapter 2

30

6 Assuming that the crystal thickness is proportional to the fold length of the primary,homogenous nucleus, the crystal thickness is: l = β [4γTm

0/(∆Hf∆T)].7 J. D. Hoffman, J. J. Weeks, J. Res. Natl. Bur. Stand. (U.S.) 1962, 66A, 13.8 (a) M. Avrami, J. Chem. Phys. 1939, 7, 1103. (b) M. Avrami, J. Chem. Phys. 1940, 8,

212. (c) M. Avrami, J. Chem. Phys. 1941, 9, 177.9 “Mulitcomponent polymer systems”, I. S. Miles, S. Rostami, Polymer Science and

Technology Series, Longman Group UK Ltd., Great Britain, 1992.10 D. R. Paul, J. W. Barlow, J. Macromol. Sci.-Rev. Macromol. Chem. 1980, C18, 109.11 “Compatibility” , D. W. Fox, R. B. Allen in “Encyclopedia of Polymer Science and

Engineering”, H. F. Mark, N. Bikales, C. G. Overberger, G. Menges, WileyInterscience, New York, 1988, 2nd ed., vol. 3, 758.

12 Useful reviews on applications of block copolymers in the solid state are: (a) I.Goodman, “Developments in block copolymers”, Vol. 1, Applied Science, London,1982. (b) I. Goodman, “Developments in block copolymers”, Vol. 2, ElsevierScience, London, 1985. (c) “Block copolymers”, G. Riess, G. Hurtrez, P. Bahadur in“Encyclopedia of Polymer Science and Engineering”, H. F. Mark, J. I. Kroschwitz,Wiley, New York, Vol. 2, 1985.

13 For applications of block copolymers in solutions see: (a) I. R. Schmolka, “Polymersfor controlled drug delivery”, CRC Press, Boston, 1991. (b) V. M. Nace, “Nonionicsurfactants. Polyoxyalkylene block copolymers”, Surfactant Science Series, MarcelDekker, New York, 1996.

14 See for instance: (a) G. M. Whitesides, J. P. Mathias, C. T. Seto, Science 1991, 254,1312. (b) T. L. Morkved, M. Lu, A. M. Urbas, E. E. Ehrichs, H. M. Jaeger, P.Mansky, T. P. Russell, Science 1996, 273, 931. (c) S. I. Stupp, V. LeBonheur, K.Walker, L. S. Li, K. E. Huggins, M. Keser, A. Amstutz, Science 1997, 276, 384. (d)M. Templin, F. Achim, A. Du Chesne, H. Leist, Y. Zhang, R. Ulrich, V. Schädler, U.Wiesner, Science 1997, 278, 1795. (e) D. Zhao, J. Feng, Q. Huo, N. Melosh, G. H.Fredrickson, B. F. Chmelka, G. D. Stucky, Science 1998, 279, 548. (f) S. A. Jenekhe,X. L. Chen, Science 1999, 283, 372. (g) Y. Fink, A. M. Urbas, M. G. Bawendi, J. D.Joannopoulos, E. L. Thomas, J. Lightwave Technol. 1999, 17, 1963.

15 S. Mann, G. A. Ozin, Nature 1996, 382, 313.16 F. Schue, in “Comprehensive Polymer Science”, G. Allen, J. C. Bevington,

Pergamon, Oxford, 1989.17 O. W. Webster, Science 1991, 251, 887.18 F. S. Bates, Science 1991, 251, 898.19 L. Leibler, Macromolecules 1980, 13, 1602.20 G. H. Frederickson, E. Helfand, J. Chem. Phys. 1987, 87, 697.21 M. W. Matsen, F. S. Bates, Macromolecules 1996, 29, 1091.22 (a) F. S. Bates, G. H. Frederickson, Annu. Rev. Phys. Chem. 1990, 41, 525. (b) F. S.

Bates, J. H. Rosedale, G. H. Frederickson, J. Chem. Phys. 1990, 92, 6255.23 D. A. Hajduk, P. E. Harper, S. M. Gruner, C. C. Honeker, G. Kim, E. L. Thomas, L. J.

Fetters, Macromolecules 1994, 27, 4063.24 M. F. Schulz, F. S. Bates, K. Almdal, K. Mortensen, Phys. Rev. Lett. 1994, 73, 86.25 The ordered bicontinuous double-diamond (OBDD) morphology is a cubic structure

with Pn3m symmetry (E. L. Thomas, D. B. Alward, D. J. Kinning, D. C. Martin, D.L. Handlin, L. J. Fetters, Macromolecules 1986, 19, 2197). Later a growing


31

skepticism in regard to this phase, motivated a reexamination of this phase and

showed that the structure possessed a cubic structure with Ia3d symmetry: D. A.Hajduk, P. E. Harper, S. M. Gruner, C. C. Honeker, G. Kim, E. L. Thomas, L. J.Fetters, Macromolecules 1995, 28, 2570.

26 I. W. Hamley, K. A. Koppi, J. H. Rosedale, F. S. Bates, K. Almdal, K. Mortensen,Macromolecules 1993, 26, 5959.

27 E. L. Thomas, D. M. Anderson, C. S. Henkee, D. Hoffman, Nature 1988, 334, 598.28 S. Förster, A. K. Khandpur, J. Zhao, F. S. Bates, I. W. Hamley, A. J. Ryan, W. Bras,

Macromolecules 1994, 27, 6922.29 D. A. Hajduk, H. Takenouchi, M. A. Hillmyer, F. S. Bates, M. E. Vigild, K. Almdal,

Macromolecules 1997, 30, 3788.30 Later it was found that both the hexagonally modulated and perforated layers were

long-lived nonequilibrium states that convert to the bicontinuous gyroid uponisothermal annealing. See for instance ref. 29.

31 K. Grosse-Brauckmann, J. Colloid Int. Sci. 1997, 187, 418.32 M. W. Matsen, M. D. Whitmore, J. Chem. Phys. 1996, 105, 9698.33 Much work on the morphological behavior of especially SBM (styrene-butadiene-

methyl methacrylate) triblock copolymers has been reported by Stadler’s group: (a) C.Auschra, R. Stadler, Macromolecules 1993, 26, 2171. (b) R. Stadler, C. Auschra, J.Beckmann, U. Krappe, I. Voigt-Martin, L. Leibler, Macromolecules 1995, 28, 3080.(c) U. Breiner, U. Krappe, V. Abetz, R. Stadler, Macromol. Chem. Phys. 1997, 198,1051. (d) U. Breiner, U. Krappe, T. Jakob, V. Abetz, R. Stadler, Polym. Bull. (Berlin)1998, 40, 219. (e) S. Brinkmann, R. Stadler, E. L. Thomas, Macromolecules 1998, 31,6566.

34 U. Krappe, R. Stadler, I. Voigt-Martin, Macromolecules 1995, 28, 4558.35 (a) U. Breiner, U. Krappe, R. Stadler, Macromol. Rapid Commun. 1996, 17, 567. (b)

U. Breiner, U. Krappe, E. L. Thomas, R. Stadler, Macromolecules 1998, 31, 135.36 (a) N. Hadjichristidis, H. Iatrou, S. K. Behal, J. J. Chludzinski, M. M. Disko, R. T.

Garner, K. S. Liang, D. J. Lohse, S. T. Milner, Macromolecules 1993, 26, 5812. (b)D. J. Pochan, S. P. Gido, S. Pispas, J. W. Mays, A. J. Ryan, J. P. A. Fairclough, I. W.Hamley, N. Terrill, Macromolecules 1996, 29, 5091. (c) Y. Tselikas, H. Iatrou, N.Hadjichristidis, K. S. Liang, K. Mohanty, D. J. Lohse, J. Chem. Phys. 1996, 105,2456. (d) S. P. Gido, Z.-G. Wang, Macromolecules 1997, 30, 6771. (e) C. M. Turner,N. B. Sheller, M. D. Foster, B. Lee, S. Corona-Galvin, R. P. Quirk, B. Annis, J.-S.Lin, Macromolecules 1998, 31, 4372. (f) M. Pitsikalis, S. Pispas, J. W. Mays, N.Hadjichristidis, Adv. Polym. Sci. 1998, 135, 1. (g) F. L. Beyer, S. P. Gido, D. Uhrig, J.W. Mays, N. Beck Tan, S. F. Trevino, J. Polym. Sci., Part B Polym. Phys. 1999, 37,3392.

37 A comprehensive review of electron microscopy from polymers is provided by: E. L.Thomas, “Electron microscopy”, in “Encyclopedia of Polymer Science andEngineering”, H. F. Mark, N. M. Bikales, C. G. Overberger, G. Menges, vol. 5,Wiley, New York, 1989.

38 (a) K. Kato, J. Electron Microscopy 1965, 14, 220. (b) K. Kato, Polym. Eng. Sci.1967, 7, 38.

39 D. W. Schwark, D. L. Vezie, J. R. Reffner, E. L. Thomas, B. K. Annis, J. Mater. Sci.Lett. 1992, 11, 352.

Chapter 2

32

40 D. C. Martin, E. L. Thomas, Polymer 1995, 36, 1743.41 The space group determines which reflections are allowed. The ratios of the allowed

reflections are given for each space group in: “International Tables for X-rayCrystallography”, vol. I, The Kynoch Press, Birmingham, 1952.

42 Small angle scattering analysis has been presented for lamellae: (a) T. Hashimoto, K.Nagatoshi, A. Todo, H. Hasegawa, H. Kawai, Macromolecules 1974, 7, 364. (b) M.Shibayama, T. Hashimoto, Macromolecules 1986, 19, 740. For hexagonally packedcylinders: (c) T. Hashimoto, T. Kawamura, M. Harada, H. Tanaka, Macromolecules1994, 27, 3063. For spheres stacked with cubic symmetries: (d) H. Matsuoka, H.Tanaka, H. Hashimoto, N. Ise, Phys. Rev. B 1987, 36, 1754. (e) H. Matsuoka, H.Tanaka, N. Iizuka, T. Hashimoto, N. Ise, Phys. Rev. B 1990, 41, 3854.

43 “ International Tables for X-ray Crystallography”, vol. III, The Kynoch Press,Birmingham, 1962.

44 (a) C. I. Chung, J. C. Gale, J. Polym. Sci., Polym. Phys. Ed. 1976, 14, 1149. (b) E.Gouinlock, R. Porter, Polym. Eng. Sci. 1977, 17, 534. (c) C. I. Chung, M. I. Lin, J.Polym. Sci., Polym. Phys. Ed. 1978, 16, 545. (d) F. S. Bates, Macromolecules 1984,17, 2607.

45 (a) K. Almdal, K. A. Koppi, F. S. Bates, K. Mortensen, Macromolecules 1992, 25,1743. (b) A. K. Khandpur, S. Förster, F. S. Bates, I. W. Hamley, A. J. Ryan, W. Bras,K. Almdal, K. Mortensen, Macromolecules 1995, 28, 8796. (c) M. F. Schulz, A. K.Khandpur, F. S. Bates, K. Almdal, K. Mortensen, D. A. Hajduk, S. M. Gruner,Macromolecules 1996, 29, 2857.

46 E. Helfand, Z. R. Wasserman, In “Developments in Block Copolymers”, I. Goodman,Applied Science, New York, 1982, vol. 1.

47 J. D. Vavasour, M. D. Withmore, Macromolecules 1993, 26, 7070.48 M. W. Matsen, M. Schick, Macromolecules 1994, 27, 4014.49 H. Hasegawa, H. Tanaka, K. Yamasaki, T. Hashimoto, Macromolecules 1987, 20,

1651.50 D. J. Pochan, S. P. Gido, J. Zhou, J. W. Mays, M. Withmore, A. J. Ryan, J. Polym.

Sci. Part B, Polym. Phys. 1997, 35, 2629.51 For a recent review on blending of block copolymers see: V. Abetz, T. Goldacker,

Macromol. Rapid Commun. 2000, 21, 16.52 (a) M. D. Withmore, J. Noolandi, Macromolecules 1985, 18, 2486. (b) M. W. Matsen,

Macromolecules 1995, 28, 5765. (c) H. Xi, S. T. Milner, Macromolecules 1996, 29,2404. (d) A. E. Likhtman, A. N. Semenov, Macromolecules 1997, 30, 7273. (e) P. K.Janert, M. Schick, Macromolecules 1998, 31, 1109.

53 (a) R. J. Roe, W. C. Zin, Macromolecules 1984, 17, 189. (b) D. J. Kinning, K. I.Winey, E. L. Thomas, Macromolecules 1988, 21, 3502. (c) T. Hashimoto, H. Tanaka,H. Hasegawa, Macromolecules 1990, 23, 43787. (d) H. Tanaka, H. Hasegawa, T.Hashimoto, Macromolecules 1991, 24, 240. (e) H. Tanaka, T. Hashimoto,Macromolecules 1991, 24, 5398. (f) K. I. Winey, E. L. Thomas, L. J. Fetters, J. Chem.Phys. 1991, 95, 9367. (g) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules1991, 24, 6182. (h) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules 1992,25, 422. (i) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules 1992, 25, 2645.(j) R. J. Spontak, S. D. Smith, A. Ashraf, Macromolecules 1993, 26, 956. (k) M. M.


33

Disko, K. S. Liang, S. K. Behal, R. J. Roe, K. J. Jeon, Macromolecules 1993, 26,2983. (l) K. J. Jeon, R. J. Roe, Macromolecules 1994, 27, 2439.

54 T. A. Witten, L. Leibler, P. A. Pincus, Macromolecules 1990, 23, 824.55 S. H. Anastasiadis, T. P. Russell, S. K. Satija, C. F. Majkrzak, Phys. Rev. Lett. 1989,

62, 1852.56 J. P. Spatz, S. Sheiko, M. Möller, Adv. Mater. 1996, 8, 513.57 J. P. Spatz, M. Möller, M. Noeske, R. J. Behm, M. Pietralla, Macromolecules 1997,

30, 3874.58 T. P. Russell, G. Coulon, V. R. Deline, D. C. Miller, Macromolecules 1989, 22, 4600.59 D. T. Clark, J. Peeling, J. M. O’Malley, J. Polym. Sci., Polym. Chem. 1976, 14, 543.60 G. Coulon, D. Auserre, T. P. Russell, J. Phys. (France) 1990, 51, 777.61 G. Coulon, T. P. Russell, V. R. Deline, P. F. Green, Macromolecules 1989, 22, 2581.62 (a) G. Coulon, B. Collin, D. Ausserre, D. Chatenay, T. P. Russell, J. Phys. (France)

1990, 51, 2801. (b) M. Maaloum, D. Ausserre, D. Chatenay, G. Coulon, Y. Gallot,Phys. Rev. Lett. 1992, 68, 1575. (c) B. Collin, D. Chatenay, G. Coulon, D. Ausserre,Y. Gallot, Macromolecules 1992, 25, 1621. (d) Z. Cai, K. Huang, P. A. Montano, T.P. Russell, J. M. Bai, G. W. Zajac, J. Chem. Phys. 1993, 98, 2376. (e) R. Mutter, B.Stühn, Macromolecules 1995, 28, 5022. (f) P. C. M. Grim, I. A. Nyrkova, A. N.Semenov, G. ten Brinke, G. Hadziioannou, Macromolecules 1995, 28, 7501. (g) G.Vignaud, A. Gibaud, G. Grübel, S. Joly, D. Ausserre, J. F. Legrand, Y. Gallot, Phys.B 1998, 248, 250. (h) J. Heier, E. Sivaniah, E. J. Kramer, Macromolecules 1999, 32,9007.

63 P. Lambooy, T. P. Russell, G. J. Kellog, A. M. Mayes, P. D. Gallagher, S. K. Satija,Phys. Rev. Lett. 1994, 72, 2899.

64 N. Koneripalli, N. Singh, R. Levicky, F. S. Bates, P. D. Gallagher, S. K. Satija,Macromolecules 1995, 28, 2897.

65 G. J. Kellogg, D. G. Walton, A. M. Mayes, P. Lambooy, T. P. Russell, P. D.Gallagher, S. K. Satija, Phys. Rev. Lett. 1996, 76, 2503.

66 N. Koneripalli, R. Levicky, F. S. Bates, J. Ankner, H. Kaiser, S. Satija, Langmuir1996, 12, 6681.

67 M. J. Fasolka, D. J. Harris, A. M. Mayes, M. Yoon, S. G. J. Mochrie, Phys. Rev. Lett.1997, 79, 3018.

68 T. L. Morkved, H. M. Jaeger, Europhys. Lett. 1997, 40, 643.69 G. T. Pickett, T. A. Witten, S. R. Nagel, Macromolecules 1993, 26, 3194.70 E. Huang, T. P. Russell, C. Harrison, P. M. Chaikin, R. A. Register, C. J. Hawker, J.

Mays, Macromolecules 1998, 31, 7641.71 D. G. Walton, G. J. Kellogg, A. M. Mayes, P. Lambooy, T. P. Russell,

Macromolecules 1994, 27, 6225.72 G. T. Pickett, A. C. Balazs, Macromolecules 1997, 30, 3097.73 M. W. Matsen, J. Chem. Phys. 1997, 106, 7781.74 (a) M. Kikuchi, K. Binder, Europhys. Lett. 1993, 21, 427. (b) M. Kikuchi, K. Binder,

J. Chem. Phys. 1994, 101, 3367.75 M. S. Turner, M. Rubinstein, C. M. Marques, Macromolecules 1994, 27, 4986.76 K. Y. Suh, Y. S. Kim, H. H. Lee, J. Chem. Phys. 1998, 108, 1253.77 L. H. Radzilowski, B. L. Carvalho, E. L. Thomas, J. Polym. Sci., Polym. Phys. 1996,

34, 3081.

Chapter 2

34

78 C. S. Henkee, E. L. Thomas, L. J. Fetters, J. Mater. Sci. 1988, 23, 1685.79 (a) C. Harrison, M. Park, P. M. Chaikin, R. A. Register, D. H. Adamson, N. Yao,

Polymer 1998, 30, 2733. (b) C. Harrison, M. Park, P. M. Chaikin, R. A. Register, D.H. Adamson, N. Yao, Macromolecules 1998, 31, 2185.

80 J. Israelachvili, “Intermolecular and Surface Forces”, 2nd ed.; Academic Press:London, 1991; p. 315.

81 Y. Liu, W. Zhao, X. Zheng, A. King, A. Singh, M. H. Rafailovich, J. Sokolov, K. H.Dai, E. J. Kramer, S. A. Schwarz, O. Gebizlioglu, S. K. Sinha, Macromolecules 1994,27, 4000.

82 (a) H. Yokoyama, E. J. Kramer, M. H. Rafailovich, J. Sokolov, S. A. Schwarz,Macromolecules 1998, 31, 8826. (b) H. Yokoyama, T. E. Mates, E. J. Kramer,Macromolecules 2000, 33, 1888.

83 A. Karim, N. Singh, M. Sikka, F. S. Bates, W. D. Dozier, G. Felcher, J. Chem. Phys.1994, 100, 1620.

84 (a) D. W. Schwark, D. L. Vezie, J. R. Reffner, E. L. Thomas, B. K. Annis, J. Mater.Sci. Lett. 1992, 11, 352. (b) D. L. Vezie, E. L. Thomas, W. W. Adams, Polymer 1995,36, 1761.

85 B. K. Annis, D. W. Schwark, J. R. Reffner, E. L. Thomas, B. Wunderlich, Makromol.Chem., Macromol. Chem. Phys. 1992, 193, 2589.

86 M. Binggeli, R. Christoph, H. E. Hinterman, J. Colchero, O. Marti, Nanotechnology1993, 4, 59.

87 Ph. Leclère, R. Lazzaroni, J. L. Brédas, J. M. Yu, Ph. Dubois, R. Jérôme, Langmuir1996, 12, 4317.

88 However, care should be taken in identifying the phases by their relative contrast inheight and phase mode images, since the contrast is not unique to a particular phasebut is sensitive to changes in the operating conditions of the microscope: J. P.Pickering, G. J. Vancso, Polym. Bull. 1998, 40, 549.

89 L. G. Parratt, Phys. Rev. 1954, 95, 359.90 J. Als-Nielsen, “Topics in Current Physics. Structure and Dynamics of Surfaces”,

Springer-Verlag, 1986, Berlin.91 O. H. Seeck, I. D. Kaendler, M. Tolan, K. Shin, M. H. Rafailovich, J. Sokolov, R.

Kolb, Appl. Phys. Lett. 2000, in press.92 D. L. Flamm, G. K. Herb in “Plasma Etching, An Introduction”, D. M. Manos, D. L.

Flamm, Academic Press: San Diego, 1988.93 A. J. van Roosmalen, J. A. G. Baggerman, S. J. H. Brader, “Dry Etching for VSLI”,

Plenum Press, New York, 1991.94 S. J. Fonash, Solid Stat. Technol. 1985, 28, 150.95 N. J. Chou, C. H. Tang, J. Parszczak, E. Babich, Appl. Phys. Lett. 1985, 46, 31.96 (a) S. A. MacDonald, H. Schlosser, H. Ito, N. J. Clecak, C. G. Willson, Chem. Mater.

1991, 3, 435. (b) S. A. MacDonald, H. Schlosser, N. J. Clecak, C. G. Willson, J. M. J.Fréchet, Chem. Mater. 1992, 4, 1364.

97 T. M. Wolf, G. N. Taylor, J. Electrochem. Soc. 1984, 131, 1664.98 J. W. Labadie, S. A. MacDonald, C. G. Willson, J. Imag. Sci. 1986, 30, 169.99 (a) H. Fujioka, H. Nakajima, S. Kishimura, H. Nagata, Proc. SPIE 1990, 1262, 554.

(b) Y. Yoshida, H. Fujioka, H. Nakajima, S. Kishimura, H. Nagata, J. Photopolym.


35

Sci. Technol. 1991, 4, 497. (c) Y. Yoshida, S. Kubota, H. Koezuka, H. Fujioka, J.Vac. Sci. Technol. B 1994, 12, 1402.

100 O. Nalamasu, F. A. Baiocchi, G. N. Taylor, ACS Symp. Ser. 1989, 412, 189.101 (a) M. Suzuki, K. Saigo, H. Gokan, Y. Ohnishi, J. Electrochem. Soc. 1983, 30, 1962.

(b) M. Morita, A. Tanaka, S. Immamura, T. Tamamura, O. Kogure, Jpn. J. Appl.Phys. 1983, 22, 659. (c) E. Reichmanis, G. Smolinski, C. G. Wilkins Jr., Solid StateTechnol. 1985, 28, 130.

102 M. A. Hartney, A. E. Novembre, F. S. Bates, J. Vac. Sci. Technol. B 1985, 3, 1346.103 (a) H. G. Graighead, L. M. Schiavone, Appl. Phys. Lett. 1986, 48, 1748. (b) Y.

Ohmura, T. Shiokawa, K. Toyoda, S. Namba, Appl. Phys. Lett. 1987, 51, 1500. (c) M.E. Gross, W. L. Brown, L. R. Harriott, K. D. Cummings, J. Linnros, H. Funsten, J.Appl. Phys. 1989, 66, 1403. (d) P. Hoffmann, H. van dan Bergh, J. Flicstein, G. BenAssayas, J. Gierak, J.-F. Bresse, J. Vac. Sci. Technol. B 1991, 9, 3483. (e) A. A. Avey,R. H. Hill, J. Am. Chem. Soc. 1996, 118, 237. (f) M. T. Reetz, M. Winter, J. Am.Chem. Soc. 1997, 119, 4539.

Chapter 2

36

Synthesis, Characterization and Thermal Properties

37

&KDSWHU��

Synthesis, Characterization and Thermal Propertiesof Ferrocenyldimethylsilane Polymers, Obtained by

Anionic Ring-Opening Polymerization #

Well-defined ferrocenyldimethylsilane homopolymers were obtained viaanionic ring-opening polymerization. The thermal behavior and morphologyof these materials was studied using differential scanning calorimetry (DSC),temperature dependent wide- and small-angle x-ray scattering (WAXS andSAXS), and atomic force microscopy (AFM). The isothermal crystallizationkinetics were analyzed in terms of the Avrami equation. The value of theAvrami exponent was approximately 3 for the majority of samples andcrystallization temperatures. Hedrites (immature spherulites) were formed atlow undercooling whereas at high levels of undercooling mature, welldeveloped spherulites were observed with small hedritic features in theircenter. The melting enthalpy for a 100% crystalline polymer was estimated tobe 35 J/g from DSC and x-ray diffraction data obtained by using specimenswith different degrees of crystallinity. All polymers investigated showedmultiple melting endotherms upon heating after isothermal crystallization,which was ascribed to a melting-recrystallization process. The semi-crystalline polymers showed a melting behavior as described by the Hoffman-Weeks equation, leading to an equilibrium melting temperature of 143 oC.

3.1 Introduction

Since the 1970’s, poly(ferrocenylsilanes) were made accessible via a condensationpolymerization of dilithioferrocene and dichlorodialkylsilane.1 However, the recent discoveryof the ring-opening polymerization (ROP) of silicon bridged ferrocenophane2 has increased

# Parts of the work described in this Chapter was published: (a) R. G. H. Lammertink, M. A. Hempenius, I.Manners, G. J. Vancso, Macromolecules 1998, 31, 795. (b) R. G. H. Lammertink, M. A. Hempenius, G. J.Vancso, Macromol. Chem. Phys. 1998, 199, 2141. (c) M. A. Hempenius, R. G. H. Lammertink, G. J. Vancso,Macromol. Symp. 1998, 127, 161.

Chapter 3

38

the opportunities for the tailoring of poly(ferrocenylsilanes). For example, thermal ROP andγ-irradiation3 in the solid state produces high molar mass poly(ferrocenylsilanes). Transition-metal catalyzed4 ROP allows for the synthesis of regioregular poly(ferrocenylsilanes).Anionic polymerization leads to poly(ferrocenylsilanes) with controlled molar masses andlow polydispersities.5 Using these methods, a range of novel polymers have been synthesized,featuring the following variations: substituents on silicon,6 number and type of bridgingatoms,7 and substituents on the cyclopentadienyl rings.8 This chapter describes the thermalbehavior and morphology of a series of well-defined poly(ferrocenyldimethylsilanes),prepared via anionic ring-opening polymerization.

3.2 Synthesis and Characterization of Poly(ferrocenylsilane)

1,1’-Dimethylsilylferrocenophane was prepared as described earlier9 and purified bysequential dissolution in heptane followed by vacuum sublimation. The monomer wasobtained in approximately 70% yield.

Polymerizations of 1,1’-dimethylsilylferrocenophane were carried out in THF usingn-butyllithium as the initiator. The living polymerization was terminated by the addition of afew drops of methanol (Chart 3.1). The polymers were precipitated in methanol and driedunder vacuum.

Fe

Si

CH3

CH3

n-Bu

H

SiFeCH3

CH3

1. n-BuLi2. MeOH

Fe Fe

Li

Lin-BuLi Cl2SiMe2

SiFeCH3

CH3

Chart 3.1 Preparation and anionic ring-opening polymerization of dimethylsilyl bridged ferrocenophane.


39

Figure 3.1 1H-NMR spectra of 1,1’-dimethylsilylferrocenophane (indicated as monomer in the spectrum) and ofpoly(ferrocenyldimethylsilane) (indicated as polymer). The peaks are labeled according to the labels in thestructures. In the spectrum of the polymer, the arrow indicates the unsubstituted terminal cyclopentadienyl ring(end-Cp).

1H-NMR spectra of the monomer and the polymer in deuterated chloroform areshown in Figure 3.1. In order to characterize the molar mass distribution function, GPCexperiments were performed on polymers dissolved in THF. The GPC traces showed narrowmolar mass distributions as expected for anionic polymerization. The characteristics of thepolymers are shown in Table 3.1. The molar mass of the polymer with the lowest degree ofpolymerization was also estimated by end group analysis using 1H-NMR. In the NMR spectrain Figure 3.1 the unsubstituted cyclopentadienyl of the terminal ferrocene unit is indicated byan arrow. The polymers are indicated as PFS, followed by their approximate degree ofpolymerization. A reasonable agreement exists between the molar mass according to GPC(using a polystyrene calibration curve) and 1H-NMR for specimen PFS-25. Also, the molarmasses estimated from the ratio of monomer to initiator agree within 10% from the valuesobtained by GPC.

Chapter 3

40

Table 3.1 Characteristics of the poly(ferrocenyldimethylsilanes) used in this study.a

Mn

[g/mol]Mw

[g/mol]Mw/Mn

PFS-25 6500, 5800b 7100 1.09PFS-50 11400 12500 1.09PFS-100 24200 26200 1.09PFS-200 44900 49500 1.10PFS-350 76700 87100 1.14a GPC in THF using polystyrene calibration curves.b Determined by end group analysis using 1H-NMR.

3.3 Thermal Properties of Poly(ferrocenyldimethylsilanes)

3.3.1 GLASS TRANSITION

The glass transition temperatures (Tg), determined with DSC on quenched samples at10 K/min, heat capacity values (∆Cp) at Tg, and the crystallization enthalpies (∆Hc) of thepoly(ferrocenyldimethylsilanes) with different molar mass are given in Table 3.2.

Table 3.2 Thermal data of the poly(ferrocenylsilanes).

Tg

[°C]∆Cp

[J/gK]∆Hc

[J/g]PFS-25 22.4 0.15 14.7PFS-50 25.7 0.16 14.2PFS-100 28.0 0.16 12.6PFS-200 31.0 0.18 10.0PFS-350 31.3 0.17 7.7

Due to the excess free volume associated with the chain ends, the value of Tg dependson the number average molar mass Mn. According to Flory,10 the value of Tg is proportionalto 1/Mn. For many polymer systems a better fit can be obtained using O’Driscoll’s equation.According to O’Driscoll, the Tg can be described as:11


41

32/n

g,gM

KTT −= ∞

In this equation, Tg,∞ is the glass transition temperature of the polymer with infinite molarmass, and the polymer dependent constant K is related to the excess free volume at the chainends. For poly(ferrocenyldimethylsilanes), this results in an extrapolated Tg,∞ value of 33.5 °C(see Figure 3.2). This value agrees well with the glass transition temperature of 33 °Cobserved for high molar mass poly(ferrocenyldimethylsilane), prepared by thermalpolymerization.6 A fit of our data to O’Driscoll’s equation resulted in a regression coefficientof 0.991, as opposed to the regression coefficient of 0.981 for the 1/Mn fit.

0.0 0.5 1.0 1.5 2.0 2.5 3.020

22

24

26

28

30

32

34

Tg

[o C]

103/M n

2/3

Figure 3.2 Tg as a function of Mn-2/3 for poly(ferrocenyldimethylsilane). The line represents a fit according to

the equation of O’Driscoll.11

The values of the crystallization enthalpies ∆Hc, determined from the crystallizationexotherm during slow cooling are given in Table 3.2. Cooling scans showed a singleexotherm from which accurate ∆Hc values could be obtained. On the other hand, it was notpossible to obtain meaningful ∆Hm values from heating scans as these displayed multipleendotherms (see later in this chapter). The magnitude of the crystallization enthalpiesdecreased with increasing molar masses.

Chapter 3

42

3.3.2 CRYSTALLIZATION KINETICS

The isothermal crystallization process was monitored in time at differentcrystallization temperatures. Figure 3.3 shows the exothermic heat flow signal of the threesamples crystallized at 90 °C as measured by DSC (PFS-200 and PFS-350 could not bemonitored accurately due to slow crystallization kinetics). The magnitude of the exothermicheat flow was approximately constant for each polymer at different temperatures.

0 5 10 15 20 25 30

PFS-100

PFS-50

PFS-25

< E

xo

Time [minutes]

Figure 3.3 Isothermal crystallization exotherms at 90 °C of PFS homopolymers.

It is obvious from Figure 3.3 that the lower molar mass results in faster completion ofthe crystallization process. Values of the exothermic area at time t and the total exothermicarea can be obtained from the DSC heat flow signal. The relative completion of thecrystallization, Xt, can be determined as a function of time as the ratio between theexothermic area at time t and the heat flow of the complete process. The results obtainedwere analyzed by using the Avrami equation:12

( )X K t ttn= − − −1 0exp ( )

where Xt is the relative conversion of the crystallization process at time t, K is a constantwhich accounts for the growth rate, nucleation rate, and temperature terms, t0 is the point intime at which crystallization begins, and n is the Avrami exponent which characterizes thetype of growing structure and nucleation mechanism. For every sample and temperature, t0was taken as the time at which the onset of the crystallization exotherm occurred. The


43

linearized form of the Avrami equation, log(-ln(1-Xt)) vs. log(t-t0), was plotted for all samplesat different crystallization temperatures. Figure 3.4 displays Avrami plots for all threesamples at 90 °C.

-0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6-3.5

-3.0

-2.5

-2.0

-1.5

-1.0

-0.5

0.0

0.5

1.0 PFS-25

PFS-50

PFS-100

log(

-ln(1

-Xt))

log(t-t0)

Figure 3.4 Avrami plots for PFS homopolymers.

Table 3.3 Results on the Avrami analysis of the isothermal crystallization of poly(ferrocenyldimethylsilanes).

TC

[°C]n average n K (10-3)

[min-n]t½

[min]PFS-25 80 3.1 2.25 6.47

85 2.9 5.21 5.5390 3.0 3.83 5.5395 3.4 3.1 ± 0.2 1.14 6.70

PFS-50 85 3.1 1.28 7.4790 2.9 3.75 6.0995 3.0 4.47 5.50100 2.9 3.0 ± 0.1 2.31 7.05

PFS-100 90 3.0 0.56 11.095 3.1 0.83 8.79100 3.7 0.10 10.6105 3.8 3.4 ± 0.4 0.08 10.7

The results of the Avrami analysis, including the crystallization temperatures, thecorresponding (and the average) Avrami exponents n, the value of the constant K in theAvrami equation, and the half-time of crystallization t½ are given in Table 3.3. An average

Chapter 3

44

Avrami exponent of approximately 3 suggests simultaneous nucleation followed by three-dimensional growth.13

A significantly higher value for the Avrami exponent was obtained only for PFS-100at 100 and 105 °C compared with the rest of the samples and temperatures studied. Thismight indicate that mixed nucleation modes are operative. The evaluation of crystallizationkinetics can also be influenced by the accuracy of the determination of the heat ofcrystallization and the choice of the starting point, t0, of the phase transition.14 Theaforementioned deviations of the Avrami exponent found for PFS-100 seem to be beyond theexperimental error.

The rate of crystallization can be expressed by the inverse of the half-time ofcrystallization, t½, which is the time needed for development of 50% of the finalcrystallinity:13

tK

n

12

12=

ln

A distinct crystallization temperature exists for each polymer at which the growth ratehas a maximum and thus the t½ value exhibits a minimum. For PFS-25 this temperature isbetween 85 and 90 °C; and for PFS-50 and PFS-100 the corresponding value is locatedaround 95 °C. A slightly lower optimal crystallization temperature for PFS-25 is probablycaused by the lower melting temperature due to dilution by end-groups.

3.3.3 MORPHOLOGY OF MELT -CRYSTALLIZED POLY (FERROCENYLSILANE )

The morphology of PFS-100 was studied by means of AFM on melt crystallizedsamples. According to the Avrami exponent (approximately 3) we can expect a spheruliticsuperstructure for this polymer when crystallized at 90 or 95 °C. Figure 3.5 A displays anAFM image of PFS-100 which was crystallized isothermally at 90 °C. Mature spheruliteswith small hedritic cores in the centers can be observed. Note the straight lines of intersectionbetween neighboring spherulites on impingement (Figure 3.5 A), which suggest simultaneousnucleation. This is in agreement with the proposed mechanism indicated by the Avramiexponents.

Figure 3.5 B displays an AFM image of PFS-100, crystallized at 105 °C. Non-maturespherulites exhibiting hedritic features are observed. In general, hedrites transform intospherulitic morphologies at more mature stages of crystallization.15

The connection between the appearance of spherulites or hedrites and crystallizationtemperature has been observed for other polymers, as well.16 At lower crystallizationtemperatures, the smaller “hedritic” cores and higher branching density allow for the


45

formation of a mature spherulite. At higher crystallization temperatures these hedritesimpinge before they develop into spherulites.

Figure 3.5 AFM images (deflection) of PFS-100, isothermally crystallized at 90 °C (A) and at 105 °C (B).According to the Avrami analysis, the exponent is higher at 105 °C (Table 3.3).

If DSC is used to determine the percent crystallinity of a polymer, then the heat offusion of the 100% crystalline material, ∆Hm,100%, must be known. The value can of ∆Hm,100%

can e.g. be determined by measuring the heat of fusion, ∆Hm, of a series of samples withdifferent degrees of crystallinity. These ∆Hm values can be correlated with the weight fractionof crystallinity for the given sample, determined from wide-angle x-ray scattering (WAXS)by comparing the integrated crystalline and amorphous diffracted x-ray intensities.17 Wefollowed this approach and determined ∆Hm for poly(ferrocenyldimethylsilane) assuming avalue of 1 for the WAXS correction factor which takes into account differences in thescattering power for the amorphous and crystalline phases for a given angular range ofdiffraction. In our analysis we also assumed equality for the values of the heat ofcrystallization and the heat of fusion for a given specimen.

Figure 3.6 shows the WAXS patterns for PFS-25 after isothermal crystallization at 90°C for different times in the DSC apparatus. One strong reflection evolves at 6.06 Å (2 theta= 14.6°) and a weaker one at 6.75 Å (2 theta = 13.1°).

The degree of crystallinity can be determined from the x-ray diffraction spectra usingthe following equation:

%100Cryst.% ⋅−

=SC

ASC

A

AA

where ASC is the area under the semi-crystalline WAXS pattern (Figure 3.6 b-e) and AA is thearea under the amorphous WAXS pattern (Figure 3.6 a). The different crystallization timescorrespond to different heats of crystallization, as can be determined form the isothermalcrystallization exotherms (see Figure 3.3).

Chapter 3

46

5 10 15 20 25 30 35 40

ed

c

ba

2θ

Figure 3.6 WAXS pattern of PFS-25 after isothermal crystallization at 90 °C for (a) 0 min, (b) 6 min, (c) 8 min,(d) 12 min, and (e) 15 min.

Figure 3.7 displays the fusion enthalpy (taken to be equal to the heat ofcrystallization) as a function of the degree of crystallinity as determined by WAXSexperiments for PFS-25. From this plot the fusion enthalpy of a 100 % crystalline polymercan be extrapolated to approximately 35 J/g.

0 10 20 30 40 50 60 70 80 90 1000

5

10

15

20

25

30

35

40

Ent

halp

y [J

/g]

Crystallinity [%]

Figure 3.7 Enthalpy of fusion vs. crystallinity as estimated by WAXS measurements. The line represents alinear fit that was extrapolated to 100% crystallinity.


47

3.3.4 MELTING BEHAVIOR OF POLY (FERROCENYLSILANES )

90 95 100 105 110 115 120 125 130 135 140 145 150 155

e

II

I 90

oC

95 oC

100 oC

105 oC

End

o >

Temperature [oC]

Figure 3.8 DSC scans (heat flow versus temperature) obtained at a scan rate of 10 K/min. Samples werecrystallized isothermally at temperatures indicated on the figure: (a) PFS-25, (b) PFS-50, (c) PFS-100, (d) PFS-200, (e) PFS-350. Sample characteristics are shown in Table 3.1.

90 95 100 105 110 115 120 125 130 135 140 145 150 155

a

II

I 80 oC

85 oC

90 oC

95 oC

End

o >

Temperature [oC]

90 95 100 105 110 115 120 125 130 135 140 145 150 155

b

II

I 85 oC

90 oC

95 oC

100 oC

End

o >

Temperature [oC]

90 95 100 105 110 115 120 125 130 135 140 145 150 155

c

II

I 90 oC

95 oC

100 oC

105 oC

End

o >

Temperature [oC]

90 95 100 105 110 115 120 125 130 135 140 145 150 155

d

II

I

90 oC

95 oC

100 oC

105 oC

End

o >

Temperature [oC]

Chapter 3

48

Figure 3.8 shows the DSC heating traces for all polymers after isothermalcrystallization at different temperatures. Before isothermal crystallization, the samples werekept in the melt for 5 minutes in order to erase any thermal history. After heating, thespecimens were quenched (300 K/min) to the crystallization temperature indicated in thefigures. After completion of the crystallization process, the corresponding heating trace wasrecorded starting directly from the crystallization temperature.

The heating scans display multiple melting transitions. Although such transitions arefrequently observed in DSC experiments, their origins are not always well understood.Several models have been proposed in the literature. These include dual lamellar thicknessmodels,18 melting recrystallization models,19and physical aging models.20

In the heating traces of all polymers, two types of transitions indicated by I and II inthe figures can be distinguished. The isothermal crystallization temperature does not affectthe position of peak II. In contrast, peaks of type I do depend on the crystallizationtemperature and shift to higher temperatures when a higher crystallization temperature wasapplied. A crystallization exotherm following transition I can be observed for samples PFS-25, PFS-50 and PFS-100 (see Figure 3.8 a, b, and c).

100 110 120 130 140 150

c

b

a

II

I

End

o >

Temperature [oC]

Figure 3.9 Irreversible melting illustrated for PFS-25: (a) melting after isothermal crystallization at 90 °C; (b)melting as in part (a) but scanned to 125 °C; (c) crystallization at 90 °C after scan (b).

A striking feature on the DSC traces is that peak I is a result of an irreversibletransition, as shown in Figure 3.9 for PFS-25. Scan (a) was taken up to 150 °C after thesample was crystallized at 90 °C. Scan (b) was recorded following the same thermaltreatment, but was only obtained by scanning to a temperature of 125 °C. The sample was


49

then cooled to 90 °C, was allowed to crystallize, and then reheated (scan (c)). Scan (c) showsthat peak I completely disappeared due to the irreversible transition. Peak I can only beformed again after complete melting and crystallization.

From these results it can be concluded that a structural reorganization process occursduring the DSC scan. This was also confirmed by the observation that the enthalpy of peak IIbecomes smaller when higher scanning rates are applied, thus not allowing the sample tosufficiently reorganize. Similar reorganization processes are known for polymers includingPET,21 PEEK22 and PE.23 For these polymers it was found that peak I represents the meltingpoint of the original crystals which are present prior to the first heating scan. Peak II wasascribed to the melting of reorganized crystals formed from the original crystals duringheating.

The position of peak I (i.e. the temperature at which the reorganization occurs)depends on the crystallization temperature. For example, in Figure 3.8 a, a sample of PFS-25is crystallized at 80 °C, then partially melts and recrystallizes at approximately 112 °C (peakI). This recrystallization temperature is still low enough to allow for the formation of astructure which melts at a temperature lower than the final melting temperature (peak II),giving rise to an additional peak, visible as a shoulder, at the lower temperature side of peakII. If on the other hand, the sample is crystallized at 95 °C, recrystallization occurs at 120 °C(peak I), and the final melting transition occurs as a single peak II. At 120 °C, thesupercooling is small enough to allow formation of the most stable structure.

The same dual melting behavior is observed for PFS-50 (Figure 3.8 b) and PFS-100(Figure 3.8 c), but the transitions occur at slightly higher temperatures. (This may be due tothe less pronounced diluting effect of the n-butyl end-groups). The shoulder in peak II is notvisible, but the endotherm becomes broader at lower recrystallization temperatures. Thehigher molar mass polymer, PFS-200 (Figure 3.8 d), displays an analogous behavior. Its DSCtrace exhibits multiple transitions. The position of the lower temperature transitions dependson the crystallization temperature (type I) whereas a final melting endotherm remains at aconstant temperature (type II). The members with lower molar mass, PFS-25, PFS-50 andPFS-100, display one melting transition followed by recrystallization. The higher polymer,PFS-200, showed additional melting endotherms. PFS-350 (Figure 3.8 e) exhibits a meltingbehavior similar to PFS-200. Its type I transitions are crystallization temperature dependentas expected. The final peak II, only visible as a shoulder, is positioned at a constanttemperature.

The position of peak I depends linearly on the crystallization temperature for allpolymers. Its position may be influenced by the presence of the recrystallization exotherm.24

It appears, however, that the position of peak I is not affected to a great extent by therecrystallization exotherm since the enthalpy of peak I hardly changes when higher scanningrates are applied. The position of peak I, Tm, can be described by the Hoffman-Weeksequation:25

( )β

β2

120cm

m

TTT

+−=

Chapter 3

50

in which Tm is the observed melting temperature, Tc the isothermal crystallizationtemperature, Tm

0 the extrapolated equilibrium melting temperature, and β is the fold length interms of primary, homogeneous nuclei. Figure 3.10 shows the Hoffman-Weeks plots for allpolymers studied in this work using the position of the first peak I transition upon heating asTm. This peak was chosen because the transition belongs to the melting of the originalcrystals formed during the isothermal crystallization.

80 90 100 110 120 130 140 150110

115

120

125

130

135

140

145

150

PFS-25 PFS-50 PFS-100 PFS-200 PFS-350

Tm [

0 C]

Tc [oC]

Figure 3.10 Hoffman-Weeks25 plot of melting temperature vs. crystallization temperature forpoly(ferrocenyldimethylsilanes) with different chain lengths.

From the intersect Tm = Tc = Tm0, the equilibrium melting temperature could be

determined. For all polymers, except for PFS-25 (Tm0 = 139 °C), an equilibrium melting

temperature of 143 °C was found.Previously, Manners et al.26 reported on the melting points of oligomers of ferrocenyl-

dimethylsilanes that have 2 to 9 repeat units. These authors found that the melting pointdisplayed an “odd-even” effect and converged to 145 °C. This is in close agreement with theequilibrium melting temperature found for the polymers described here, but is significantlyhigher than the previously reported melting temperature of 122 °C of high molar masspoly(ferrocenyldimethylsilane).6a

Figure 3.11 shows the long-period L (thickness of a crystalline lamella and anamorphous layer) of PFS-25 as a function of temperature after isothermal crystallization at 90°C, deduced from the maximum of the SAXS intensity peak using Bragg’s law. In Figure3.11 the temperatures of the transition I and II, respectively, as determined on the DSCthermogram (see Figure 3.8 a), are also shown. The average long-period increases from about13 nm to 15.5 nm in the temperature range of our experiments.


51

The observed increase in the long-period (SAXS), in association with findings fromDSC, WAXS and AFM suggests that it is a lamellar thickness transformation which accountsfor the multiple endotherms observed in the DSC experiments. This is due to thin lamellaeformed at high degrees of supercooling resulting in the first melting endotherm (type I) uponheating. Subsequently, recrystallization occurs when these thin lamellae melt. Since a largenumber of nuclei are still present up to the equilibrium melting temperature, therecrystallization occurs very rapidly, leading to thicker lamellae. These thicker lamellae meltat a higher temperature resulting in the second endotherm (type II).

80 90 100 110 120 130 140 15012

13

14

15

16

III

Long

Per

iod

[nm

]

Temperature [oC]

Figure 3.11 Long-period as a function of temperature for PFS-25 after isothermal crystallization at 90 °C.

3.4 Conclusions

The melting and crystallization behavior and morphology of a series of anionicallyprepared poly(ferrocenyldimethylsilanes) is described. Isothermal crystallization kineticswere analyzed in terms of the Avrami equation and the Avrami exponents were found to beapproximately three for all studied temperatures. Some deviations were found for PFS-100 atrelatively high crystallization temperatures. Furthermore, it was found that at relatively lowundercooling, hedrites were formed rather than spherulites. The melting enthalpy of a 100%crystalline poly(ferrocenyldimethylsilane) was estimated to be 35 J/g, as was obtained fromx-ray scattering and DSC measurements on a series of polymers with different degrees ofcrystallinity. The multiple melting endotherms that were observed in DSC heating scans wereascribed to a lamellar reorganization process that occurs during the experiment. The observedmelting temperatures of crystals formed during isothermal crystallization could be described

Chapter 3

52

by the Hoffman-Weeks equation. The equilibrium melting temperature of poly(ferrocenyl-dimethylsilane) at 143 °C was obtained by extrapolation. During small angle X-ray scatteringexperiments there was an increase observed in the long period upon heating through the typeI endotherm. From these findings it was concluded that thinner lamellae, which are formedduring isothermal crystallization, melt upon heating (peak I) and recrystallize into thickerlamellae. The thicker lamellae then melt at a higher temperature, giving rise to a secondendotherm (peak II).

3.5 Experimental

N,N,N’,N’-Tetramethylethylenediamine (TMEDA), ferrocene, dichlorodimethylsilaneand n-butyllithium were purchased from Aldrich. TMEDA was distilled from sodium, andferrocene was Soxhlet extracted with cyclohexane prior to use. Dimethylsilylferrocenophanewas prepared as described earlier.9 Polymerization was carried out in THF in a gloveboxpurged with prepurified nitrogen. n-Butyllithium was used as initiator, and the reaction wasterminated after 2 hours by adding a few drops of methanol. The polymers were precipitatedtwice in methanol and dried in vacuo.

1H NMR spectra were recorded in deuterated chloroform using a Bruker AC 250spectrometer that was operated at 250.1 MHz.

GPC measurements were carried out in THF using microstyragel columns with poresizes of 105, 104, 103 and 106 Å (Waters) equipped with a dual detection system consisting ofa differential refractometer (Waters model 410) and a differential viscometer (Viscotekmodel H502). The molar masses were referenced to a polystyrene calibration curve.

A Perkin-Elmer DSC-7 equipped with a PE-7700 computer using TAS-7 softwarewas used for thermal analysis. In DSC experiments, unless otherwise reported, a scan rate of10 K/min was employed. Tg values quoted in this study correspond to the inflection point atthe heat capacity change as determined by the software of the equipment used. Meltingtemperatures were taken from the peak maximum. Isothermal crystallization experimentswere performed by quenching (300 K/min) the samples from the melt to a selected holdingtemperature. When crystallization was complete, the thermograms were recorded startingfrom the holding temperature.

Wide angle X-ray scattering (WAXS) patterns were recorded by a Siemens D500diffractometer equipped with a water cooled hot stage. A CuKα x-ray source (λ = 1.542 Å)was used. The samples were melted on an aluminum plate, isothermally crystallized, andexamined in reflection mode. The angle of diffraction, 2-theta, ranged from 5 to 30 °. A stepsize of 0.1 ° with a collection time of 12 s for each step was used. For crystallinitydetermination a Philips X’Pert-1 set-up using a Cu Kα x-ray source (λ = 1.542 Å) was used.2 Theta step sizes of 0.04 ° were taken. The specimens for WAXS measurements wererapidly cooled from the melt and isothermally crystallized at 90 °C for different times in theDSC. After a certain crystallization time, the sample was quenched (300 K/min) to room


53

temperature. The value of the heat of crystallization, ∆Hm, was taken from the isothermalcrystallization exotherm. Samples were removed from DSC cups and pressed to pellets forWAXS measurements.

Small angle X-ray scattering (SAXS) experiments were performed using a Kratkycamera (Paar, Graz) in vacuo and a CuKα source (λ = 1.542 Å). A water cooled hot stagecontrolled temperatures up to 300 °C with an accuracy of 0.4 °C. Thin samples of about 150µm were pressed under 3 kN at 50 °C and slowly heated under vacuum inside the SAXScamera. From the melt the sample was cooled to the desired crystallization temperature, andthe SAXS experiments were started. Measurements were performed at temperatures rangingfrom 90 to 160 °C.

Atomic force microscopy (AFM) images were taken in air using a NanoScope III(Digital Instruments) equipped with a J scanner. The scans were carried out in contact modeutilizing a NanoProbe 100 µm triangular Si3N4 cantilever with a quoted spring constant of0.38 N/m. Samples for AFM investigations were crystallized with a free surface in a MettlerFP82 hotstage.


1 H. Rosenberg, U.S. Pat. 1969, 3,426,053.2 D. A. Foucher, B. Z. Tang, I. Manners, J. Am. Chem. Soc. 1992, 114, 6246.3 J. Rasburn, R. Petersen, T. Jahr, R. Rulkens, I. Manners, G. J. Vancso, Chem. Mater.

1995, 7, 871.4 (a) Y. Ni, R. Rulkens, J. K. Pudelski, I. Manners, Macromol. Rapid Commun. 1995,

16, 637. (b) P. Gómez-Elipe, P. M. Macdonald, I. Manners, Angew. Chem. Int. Ed.Engl. 1997, 36, 762.

5 (a) R. Rulkens, Y. Ni, I. Manners, J. Am. Chem. Soc. 1994, 116, 12121. (b) R.Rulkens, A. J. Lough, I. Manners, J. Am. Chem. Soc. 1994, 116, 797.

6 (a) M. T. Nguyen, A. F. Diaz, V. Dement’ev, K. H. Pannell, Chem. Mater. 1993, 5,1389. (b) D. A. Foucher, R. Ziembinski, B. Z. Tang, P. M. Macdonald, J. Massey, C.R. Jaeger, G. J. Vancso, I. Manners, Macromolecules 1993, 26, 2878. (c) J. K.Pudelski, R. Rulkens, D. A. Foucher, A. J. Lough, P. M. Macdonald, I. Manners,Macromolecules 1995, 28, 7301.

7 (a) J. M. Nelson, H. Rengel, I. Manners, J. Am. Chem. Soc. 1993, 115, 7035. (b) J. K.Pudelski, D. P. Gates, R. Rulkens, A. J. Lough, I. Manners, Angew. Chem. Int. Ed.Engl. 1995, 34, 1506. (c) C. Angelakos, D. B. Zamble, D. A. Foucher, A. J. Lough, I.Manners, Inorganic Chemistry 1994, 33, 1709. (d) D. A. Foucher, M. Edwards, R. A.Burrow, A. J. Lough, I. Manners, Organometallics 1994, 13, 4959.

8 J. M. Nelson, A. J. Lough, I. Manners, Angew. Chem. Int. Ed. Engl. 1994, 33, 989.9 A. B. Fischer, J. B. Kinney, R. H. Staley, M. S. Wrighton, J. Am. Chem. Soc. 1979,

101, 6501.10 T. G. Fox, P. J. Flory, J. Appl. Phys. 1950, 21, 581.

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11 K. O’Driscoll, R. Amin Sanayei, Macromolecules 1991, 24, 4479.12 (a) M. Avrami, J. Chem. Phys. 1939, 7, 1103. (b) M. Avrami, J. Chem. Phys. 1940, 8,

212. (c) M. Avrami, J. Chem. Phys. 1941, 9, 177.13 B. Wunderlich, “Macromolecular Physics”, vol. 2, Academic Press, New York.14 (a) J. Tomka, Eur. Polym. J. 1968, 4, 237. (b) D. Grenier, R. E. Prud’homme, J.

Polym. Sci., Part B Polym. Phys. 1980, 18, 1655.15 D. C. Basssett, Phil. Trans. R.. Soc. Lond. 1994, 348A, 29.16 (a) P. H. Geil “Polymer single Crystals”, InterScience, New York, 1963, p. 201. (b) Z.

Pelzbauer, A. Galeski, J. Polym. Sci., Polym. Symp. 1972, 38, 23. (c) R. C. Allen, L.Mandelkern, J. Polym. Sci., Part B Polym. Phys. 1982, 20, 1465. (d) M. Mihailov, E.Nedkov, I. Goshev, J. Macromol. Sci. Phys. 1987, B15, 313.

17 B. Wunderlich, “Macromolecular Physics”, vol. 1, Academic Press, New York,p.387.

18 (a) P. Cebe, S. D. Hong, Polymer 1986, 27, 1183. (b) D. C. Bassett, R. H. Olley, I. A.M. Al Raheil, Polymer 1988, 29, 1745.

19 (a) D. J. Blundell, Polymer 1987, 28, 2248. (b) S. Z. D. Cheng, M. Y. Cao, B.Wunderlich, Macromolecules 1986, 19, 1868.

20 V. Velikov, H. Marand, Bull. Am. Phys. Soc. 1994, 39, 569.21 (a) P. J. Holdsworth, A. Turner-Jones, Polymer 1971, 12, 195. (b) G. C. Alfonso, E.

Pedemonte, L. Ponzetti, Polymer 1979, 20, 104.22 (a) Y. Lee, R. S. Porter, Macromolecules 1987, 20, 1336. (b) B. S. Hsiao, K. H.

Gardner, D. Q. Wu, B. Chu, Polymer 1993, 34, 3996. (c) A. M. Jonas, T. P. Russell,D. Y. Yoon, Macromolecules 1995, 28, 8491. (d) F. J. Medellin-Rodriguez, P. J.Philips, J. S. Lin, Macromolecules 1996, 29, 7491.

23 D. P. Pope, J. Polym. Sci., Part B Polym. Phys. 1976, 14, 811.24 Y. Lee, R. S. Porter, J. S. Lin, Macromolecules 1989, 22, 1756.25 J. D. Hoffman, J. J. Weeks, J. Res. Natl. Bur. Stand.(U.S.) 1962, 66A, 13.26 R. Rulkens, A. J. Lough, I. Manners, S. R. Lovelace, C. Grant, W. E. Geiger, J. Am.

Chem. Soc. 1996, 118, 12683.

Reactive Ion Etching of Poly(ferrocenylsilanes)

55

&KDSWHU��

Poly(ferrocenyldimethylsilanes) for Reactive IonEtch Barrier Applications #

The behavior of poly(ferrocenyldimethylsilane) (PFS) under reactive ionetching conditions was investigated. Due to the presence of iron and silicon,the polymer was found to be relatively stable towards oxygen plasma etchingcompared to common organic polymers. Depending on the conditionsemployed, the etching rate ratio for poly(isoprene) vs. poly(ferrocenylsilane)ranged from 20:1 to 50:1. During etching a thin oxide layer is formed at thesurface of the organometallic polymer that contains iron and silicon. Theoxide layer was characterized by x-ray photoelectron spectroscopy (XPS) andAuger electron spectroscopy (AES) using depth profiling. Iron was found tobe more resistant towards removal by oxygen plasma compared to silicon.Due to the presence of iron, the polymer is also resistive to CF4/O2 etching. Aregular pattern consisting of parallel lines of PFS on a solid substrate wasobtained by micromolding in capillaries (MIMIC) (a microcontact printingtechnique) which could be transferred into the underlying substrate. Under theconditions employed, silicon, and silicon nitride substrates were etched atapproximately 20 nm/min whereas no thickness decrease was observed for thepoly(ferrocenylsilane) film, even after etching times of 10 minutes.

4.1 Introduction

Production of integrated circuits is often performed using low-pressure radiofrequency plasmas to etch thin films. The patterns of interest are formed by lithographicprocesses. Such processes consist of two steps, namely delineation of the pattern in a thinpolymer film and transfer of the pattern into the underlying substrate using an appropriateetching technique. In order to obtain smaller and smaller feature sizes, thinner resist layershave proven to be necessary to avoid pattern collapse at small dimensions. The sensitivity for

# The work described in this Chapter has been submitted for publication in Chemistry of Materials.

Chapter 4

56

the employed irradiation (UV, X-rays, etc.) and the etch resistance of the photoresist aresignificant factors that determine the quality of the photoresist.

Since the mid-1970s the same basic chemistry associated with a typical conventionalpositive photoresist has been used. The system consists of a diazonaphthoquinone sensitizerand a novolac resin. Upon irradiation, an indene carboxylic acid is generated which allowsthe novolac resin to dissolve in an appropriate solvent.1 However, as the dimensions of thefeatures desired to be fabricated become progressively smaller, the sensitivity and absorptionproperties of this resist system has reached its limitations.

The next generation of resist systems employs the photogeneration of species that cancatalyze chemical reactions in a subsequent process step. Such systems are called chemicallyamplified resists.2 The most widely studied system employs photoacid generation combinedwith an acid labile group such as the tert-butyloxycarbonyl (“t-BOC”). The acid catalyzes thedecomposition of the t-BOC groups, and thereby modifies the chemical composition ( Chart4.1). Several different trajectories can be followed from here leading to both positive andnegative tone resists.

O O

O

t-Bu OH

OSiMe3

H+

in exposed areas,acid is generated

Me2N-SiMe3

aqueous base soluble

O2 plasma developable

A

B

Chart 4.1 Schematic representation of the chemical amplification process.

The so-called plasma-developable resist systems generally involve a hydrocarbonpolymer film containing areas which are selectively loaded with an inorganic component.3

When placed in an oxygen plasma environment, the organic regions etch quickly while theareas containing the inorganic component (such as silicon,4 tin,5,6 germanium,7 or titanium8)are converted into a nonvolatile oxide. For example, the selective silylation of areas is anapproach to obtain etch resistance in an oxygen plasma.9 The selective incorporation ofinorganic components can be accomplished by irradiating the film and in a subsequent steptreating the film with an inorganic reagent, either from solution or from the gas phase. Mostof the published work that follows this general process deals with phenolic hydroxyl-containing films that react with a silylating agent to yield the corresponding silylether (seeroute B in Chart 4.1).

Multilevel resist processes have been developed to separate the imaging layer fromthe process etch mask. Simply stated, multilevel lithography combines the properties of theimaging layer with those of an etching barrier. Generally either random copolymers10 orblock copolymers11 are used, which possess a monomer unit that is incorporated to provide


57

sensitivity while a monomer containing an inorganic component is added to enhance the etchselectivity. Such systems are essential in the drive towards smaller dimensions. Muchattention has been devoted to polymers containing inorganic components that are useful asetch barrier in multilevel applications. In particular organosilicon polymers are goodcandidates for multilevel lithography. Organic polymers are converted into volatile productsupon exposure to oxygen plasmas, whereas silicon-containing polymers form a thin, non-volatile silicon oxide layer at the surface.12,13 The high etch rate ratio between organic andorganosilicon polymers has been frequently used to obtain high aspect ratio features in anoxygen plasma.

Organometallic materials have been employed for direct writing methods inlithographic applications as well. Direct writing of metallic features was possible byirradiating thin films of organometallic materials, a metal-organic cluster coordinationcompound, with highly focussed charged particle beams.14 Irradiation of the organometallicspecies resulted in the physical removal of the organic stabilizers with concomitant formationof insoluble metallic features.

Oxygen and oxygen containing plasmas are most commonly employed to modifypolymer surfaces. Reactive ion etching can be divided into two etching processes, namelychemical and physical.15 Chemical etching refers to the chemical reactions that take placebetween the active plasma species and the surface of interest. After absorption of the reactivespecies, the reaction takes place and the product desorbs from the surface. Physical etchingoccurs due to the bombardment of positive ions. The positive ions are accelerated towards thesurface, and break bonds upon impact, which results in physical etching. The balancebetween the two etching mechanisms depends on many variables, such as gas pressure andcomposition, reactor design, and temperature.16,17

Pattern transfer into the substrate is frequently accomplished using fluorocarbonreactive ion etching. Fluorocarbon plasmas have been extensively used to etch SiO2 andalthough the process of etching with such plasmas is very complex and diverse, some generalobservations can be highlighted. The C/F atomic ratio of the feed gas is a crucial parameterfor the nature of the plasma.18 CF and CF2 radicals [CFx] correlate with fluorocarbon filmformation at the substrate, whereas fluorine atoms [F] present in the plasma account for theetching of the substrate. When the C/F atomic ratio increases, relatively more CF and CF2

radicals will be present at the expense of atomic fluorine, and the plasma becomespolymerizing rather than etching.19 Even a very thin fluorocarbon film deposited on thesubstrate results in significant reduction of the SiO2 etch rate. To suppress this so-called etchstop, the fluorocarbon gases have been diluted with other gases like oxygen.20,21 Oxygenatoms act as scavengers for carbon and consequently a relative higher [F] (the etchingcomponent) rather than [CFx] (the polymerizing component) is obtained.

This chapter describes the potential of poly(ferrocenyldimethylsilane) as an etchbarrier. As a proof of principle, a regular pattern of the polymer was generated by a softlithographic method,22 similar to microcontact printing (µCP),23 and etched into theunderlying substrate.

Chapter 4

58

4.2 Results and Discussion

4.2.1 OXYGEN REACTIVE ION ETCHING

300 298 296 294 292 290 288 286 284 282 280 278 276

0.0

5.0x104

1.0x105

1.5x105

2.0x105

2.5x105

3.0x105

3.5x105

C1s

A

PFS PFS O

2 Etched

Cou

nts

Binding Energy (eV)

540 538 536 534 532 530 528 526 524 522 520 518 5160.0

5.0x104

1.0x105

1.5x105

2.0x105

2.5x105

3.0x105

O1s

B

PFS PFS O

2 Etched

Co

unts

Binding Energy (eV)

Figure 4.1 XPS spectral details of poly(ferrocenyldimethylsilane) (PFS) thin films for carbon (A, C1s) andoxygen (B, O1s). Upon oxygen etching, the carbon concentration decreased while the oxygen concentrationincreased.


59

Poly(ferrocenylsilane) films were exposed to oxygen reactive ion etching (O2-RIE)conditions. During such treatments, oxygen ions are formed and accelerated towards thesubstrate by means of a voltage (bias voltage). Upon impact of the ionized species with theorganometallic film, surface reactions and etching processes take place. Since organometallicspecies are known to act as an etching barrier in oxygen plasma, the influence of O2-RIE onpoly(ferrocenylsilane) surfaces was investigated.

The XPS spectra of poly(ferrocenyldimethylsilane) give information regarding theatoms present at the surface. Furthermore, the shifts in binding energies also indicate changesof the chemical environment. A survey scan of a thin poly(ferrocenyldimethylsilane) filmdetects the elements that are present within approximately the top ten nm.24 More detailedscans can then give additional information about the chemical environment of the differentelements.

During O2-RIE, a power of 75 W was applied which resulted in a bias voltage ofapproximately 500 V. The pressure inside the etching chamber was kept constant at 10mTorr, while an oxygen gas flow of 20 cm3/min was used. From Figure 4.1 it is clear that thecarbon concentration was reduced at the surface upon oxygen etching, while the oxygenconcentration increased. The aromaticity has been destroyed during the oxidation, as inferredfrom the disappearance of the small shakeup peak for the C1s signal at about 291 eV. This isconsistent with the formation of a thin oxide layer, from which carbon atoms are largelyexpelled.

The XPS spectra for iron are shown in Figure 4.2. The iron signals at 708.8 eV and721.5 eV have not decreased in magnitude as significantly as the carbon signal has. Therelative atomic concentration of iron has even increased (see Table 4.1). Furthermore, uponO2-RIE the iron signals have shifted to higher binding energies of 712.6 eV and 726.0 eV,and the peaks have become broader compared to the iron signals of the untreated film.Obviously, the chemical environment has been changed. The shift to higher binding energiesfor the Fe-electrons is indicative for the formation of an oxide.

The XPS data for silicon show comparable trends as were discussed for iron. Therelative concentration of silicon increased upon oxygen plasma treatment (see Table 4.1).However, the relative increase is not as pronounced as observed for iron. In the XPS spectrafor silicon (Figure 4.3), a shift to higher binding energy can be observed as a result of theoxygen plasma treatment (from 101.1 eV to 102.7 eV), which is indicative of the formationof an oxide. The peak position of silicon after the O2-RIE treatment at 102.7 eV is aboutmidway between the expected positions for siloxane (101.8 eV) and SiO2 (103.6 eV).

Chapter 4

60

735 730 725 720 715 710 705 7005.0x104

1.0x105

1.5x105

2.0x105

2.5x105

3.0x105

Fe2p PFS PFS O2 Etched

Cou

nts

Binding Energy (eV)

Figure 4.2 XPS spectra of Fe2p of the as cast film (solid line) and the O2 RIE treated film (dashed line).

110 108 106 104 102 100 98 96 94 92 901.0x104

1.5x104

2.0x104

2.5x104

3.0x104

3.5x104

4.0x104

Si2p PFS PFS O

2 Etched

Cou

nts

Binding Energy (eV)

Figure 4.3 XPS spectra of Si2p of the as cast film (solid line) and the O2 RIE treated film (dashed line).

The atomic concentrations according to the XPS measurements, prior to and after theO2-RIE treatment, are given in Table 4.1. For the PFS homopolymer, atomic concentrationsfor C : Si : Fe of 12 : 0.86 : 0.97 are found, which is near the theoretical relative amount of 12: 1 : 1 (C12H14FeSi for a repeat unit). Following O2-RIE the oxygen concentration is increased


61

significantly, while the carbon concentration decreased. Interesting phenomena can be seen,if one considers the relative amounts of the different elements. Both the Si/C and the Fe/Cratio increase upon oxygen etching, indicating that Si and Fe are more stable towards removalusing an oxygen plasma. Furthermore, the Fe/Si ratio increases, suggesting that Fe is morestable than Si towards RIE treatment using oxygen plasma. Although silicon-containingpolymers were mainly employed to enhance stability towards oxygen plasma treatments,these results indicated that the incorporation of other inorganic elements, such as iron, cansignificantly improve the etch resistance in oxygen plasmas.

From the XPS results it is clear that an oxide layer is formed at the surface ofpoly(ferrocenyldimethylsilane) upon oxygen plasma treatment. The simultaneous formationand removal due to ion sputtering of the oxide results in relatively low etching rates in O2

RIE compared to common organic polymers. The etching rates for organic polymers andpoly(ferrocenylsilane) depend strongly on the operational settings. As mentioned before, theinfluence of etching can be divided into two parts, namely chemical and physical etching.Physical etching will become more pronounced when the bombardment of ions take place athigher velocity. Therefore, the physical etching will be stronger at higher power settings, andcorrespondingly the chemical etching effect will be more pronounced at relative low powersettings.

Table 4.1 Atomic concentrations of PFS homopolymer films before and after O2-RIE.

PFS homopolymer PFS homopolymerO2 etched

element atom % atom %C1s 85.1 33.3O1s 1.9 41.8Si2p 6.1 8.0Fe2p 6.9 16.9

The etching rates were determined for poly(styrene), poly(isoprene), and poly(ferro-cenyldimethylsilane) at a fixed pressure of 10 mTorr and an oxygen gas flow of 20 cm3/min,while varying the applied power which influences the acceleration voltage of the ionstowards the surface. At this point, it is important to notice that the etching rates for theorganic polymers were measured over a much shorter time period (typically 15 to 20seconds) compared to poly(ferrocenyldimethylsilane) (up to 2 minutes). Steady-state etchingconditions are in general not achieved after very short times. However, a significant influenceof the applied power could be observed. The etching rate ratios increased from approximately20 : 1 to almost 50 : 1 for organic polymers compared with poly(ferrocenyldimethylsilane) asthe power was decreased from 75 W to 20 W.

From Auger electron spectroscopy (AES), using argon ion-sputtering depth profiling,the thickness of the surface oxide layer can be estimated. The Auger electrons of the elementsof interest can be measured as an argon ion-beam is slowly digging (~ 1 nm/cycle) through

Chapter 4

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the film. A 106 nm thin poly(ferrocenyldimethylsilane) film was exposed to an O2-RIEtreatment for 15 seconds and the film composition was measured as a function of sputteringcycles. If the sputtering rate by the argon ion gun is constant through the film, the sputteringtime can be correlated with the z-distance in the film if the thickness is known.

Figure 4.4 Auger electron spectroscopy of an oxygen etched film of poly(ferrocenylsilane) of 106 nm. Cyclesstart at the surface and end at the substrate.

The AES depth profile confirms the observations made by XPS. The film thickness ofthe etched polymer layer was approximately 106 nm. In the AES spectra, the first sputteringcycle corresponds to the sample surface, and the increase of the silicon signal around cycle 90indicates the underlying silicon substrate. A thin oxide-rich layer of approximately 10 nm ispresent at the surface (assuming a constant sputtering rate). It is interesting to compare thecomposition at the surface (sputtering cycle 0) with the internal composition of the film.From this, it can be seen that relatively less carbon and a significant amount of oxygen ispresent at the surface oxide layer compared to the film interior. Furthermore, the amount ofsilicon at the surface seems to be approximately equal compared to the amount below theoxide layer of the film. As it was also shown in Table 4.1, more silicon was removed fromthe surface compared to iron after oxygen plasma treatment. This is also very clear fromFigure 4.4, where the surface layer shows a higher intensity of iron compared to the bulk ofthe film.


63

4.2.2 TETRAFLUOROCARBON REACTIVE ION ETCHING

Radio-frequency discharges in low-pressure fluorocarbon gases are often used foretching silicon, silicon oxide, and silicon nitride. The gas mixtures are composed of CF4,CHF3, C2F6, C4F8, etc. By adjusting the composition of the gas, the character of the plasmacan be dramatically changed.19

Tetrafluorocarbon (CF4) shows the highest relative etching characteristics for amaterial reactive with fluorine atoms. Its decomposition in the plasma is characterized by thehighest concentration of fluorine atoms (F) and the lowest concentration of CF and CF2

radicals. If the F/C atomic ratio of the feed-in gas decreases, the concentration of CF and CF2

radicals will be higher at the expense of F, and the plasma becomes “polymerizing” andhence protective rather than “etching”. The deposition of polymer film on silicon25 protectsthe substrate from chemical etchants (F), thus stopping the etching. In order to reduce thedeposition of fluorocarbon polymer during the reactive ion etching treatment, the CF4 gaswas diluted with oxygen (20 vol% CF4 in the gas mixture).18 The addition of O2 to afluorocarbon etching plasma usually results in a significant increase in the etch rate of Si, aswell as reduced fluorocarbon polymer deposition.20

A pattern was generated by a microcontact printing technique known asmicromolding in capillaries (MIMIC).26 To briefly describe this process; in MIMIC apolydimethylsiloxane (PDMS) stamp is placed on a clean silicon substrate. When a drop ofthe polymer solution in THF is placed next to the PDMS stamp, the capillary force transportsthe polymer solution into the pattern. After several minutes the solvent is evaporated ordiffused into the stamp, and the stamp can be removed, leaving the polymer on the substrate(Figure 4.5).

Figure 4.5 Schematic of the procedure used in MIMIC. A PDMS stamp is placed onto a substrate (A and B). Adrop of the polymer solution is placed in contact with the stamp and substrate, and the channels are filled bycapillary action (C). After evaporation of the solvent, the stamp is removed and the polymer remains on thesubstrate (D).

The MIMIC process can be used to transfer a polymer pattern on a solid substrateconsisting of parallel lines. The AFM image in Figure 4.6 shows this homopolymer filmpattern, both before and after CF4/O2 RIE treatment. The employed plasma gas consisted of20 vol% CF4 and 80 vol% O2. The pressure inside the etching chamber was kept constant at 2mTorr, with a gas flow of 5 cm3/min. Applying a power of 20 W resulted in a bias voltage ofapproximately 160 V.

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64

Figure 4.6 AFM images (left) and single line scans (right) of a ferrocenyldimethylsilane homopolymer pattern,before and after CF4/O2 RIE treatment. The top half of the AFM image corresponds to an unetched area and thebottom half of the image displays the etched area. The sample was etched for 10 minutes, which led to anincrease of the domain height of approximately 200 nm.

The pattern transfer clearly demonstrates the potential of the ferrocenyldimethylsilanepolymer as an etch resist. The features that are present in the original line pattern can still beobserved after the CF4/O2 plasma treatment. The shapes of the polymer domains have notchanged during etching. Therefore, the section analysis from the AFM images can besuperimposed onto each other, which results in a clear illustration of the amount of siliconthat has been removed by the reactive ion etching treatment. Figure 4.7 displays the linescans from the unetched and etched samples. The gray area in this figure corresponds to theremoved silicon. Approximately 200 nm of silicon was removed in 10 minutes, resulting inan etch rate of 20 nm/min which is about equal to the etch rate of silicon nitride under similarconditions.

Figure 4.7 AFM section analysis, illustrating the etched substrate as a result of a CF4/O2 RIE treatment. During10 minutes of etching, approximately 200 nm of silicon has been removed, indicated by the gray area.


65

The AFM images and line scans indicate that poly(ferrocenyldimethylsilane) exhibitsgood dimensional stability during the plasma treatment. Although no effort was taken tooptimize and alter the etching conditions, it has been clearly shown thatpoly(ferrocenylsilane) has promising properties to function as an etch barrier.

4.3 Conclusions and Outlook

The potential of poly(ferrocenyldimethylsilane) as etch resistant polymer has beendemonstrated under reactive ion etch conditions. Upon exposure to an oxygen plasma, anapproximately 10 nm thin oxide layer was formed at the surface. The oxide resulted inrelatively low etching rates in oxygen plasmas compared to organic polymers. Due to thepresence of iron in the polymer and its oxide layer, low etching rates were also obtainedwhen fluorocarbon plasmas were applied. It was possible to increase the aspect ratio of thefeatures by etching into the substrate while the masking layer was not affected (i.e. very highselectivity towards silicon). Masking layers with such high etching barrier are potentiallyuseful for very thin resist layer applications or for high aspect ratio structures. Suchproperties are preferable for high-density imaging, and can not be expected from siliconcontaining etch barriers. The etching resistance has been demonstrated, so a challenge is leftto combine the etching properties with developable species. For example, the preparation ofblock copolymers consisting of poly(ferrocenylsilane) combined with a radiation sensitiveblock might result in a bilevel resist.11

4.4 Experimental

Thin films (approximately 60-100 nm) were prepared by spin-coating of a polymersolution in toluene onto a silicon wafer, containing its native oxide, at 4000 rpm for 30 s. Thewafers were cleaned prior to spin-coating. The cleaning involved treatment for 10 minutes in100% nitric acid, rinsing subsequently with deionized water, followed by additional treatmentfor 15 minutes in 70% nitric acid at 95 °C, and finally rinsing with deionized water and dryspinning. Another cleaning procedure involved cleaning the wafers with piranha solution (70vol% sulfuric acid, 30 vol% hydrogen peroxide (60% solution)).

Film thicknesses were varied by spin-casting from different concentrations of polymersolutions and were determined by ellipsometry. The thickness was determined on 25 spots ona wafer, and the standard deviation was always less than 1%.

Microcontact printed patterns were prepared using PDMS stamps. After the stampwas placed on the wafer, a drop of polymer solution in THF was placed next to the stamp.Due to capillary action, the solution diffused into the voids between the stamp and the wafersubstrate. After a few minutes the stamp could be removed.

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Reactive ion etching (O2-RIE) experiments were carried out in an Elektrotech PF 340apparatus. During oxygen RIE the pressure inside the etching chamber was 10 mTorr, thesubstrate temperature was set at 10 °C, and an oxygen flow rate of 20 cm3/min wasmaintained. The power was varied between 20 and 75 W. To determine the etch rate of theorganic polymers, films were etched under the condition mentioned above for 10-20 s, andthe decrease in film thickness was determined by ellipsometry. The etch rate ofpoly(ferrocenyldimethylsilane) could be determined over longer etching times, typically 1-2minutes. CF4/O2 RIE (20% CF4 and 80% O2) was performed at a pressure of 2 mTorr, with asubstrate temperature of 10 °C. The power used was set at 20 W, which resulted in a biasvoltage of approximately 170 V.

Silicon nitride substrates were prepared by plasma enhanced chemical vapordeposition (PECVD) onto silicon substrates. Typically thicknesses of approximately 100 nmwere used.

Atomic Force Microscopy (AFM) experiments were performed on a DigitalInstruments NanoScope IIIa, operated in the contact mode.

Both etched and unetched 60 nm thin films of PFS homopolymer were studied by X-ray Photoelectron Spectroscopy (XPS). Measurements were performed with a Kratos XSAM-800 spectrometer using a Mg Kα X-ray source (1253.6 eV). The input power was set at 150W (15 kV and 10 mA) and the hemispherical analyzer was placed perpendicular to thesample surface. Survey scans (0-1100 eV) were recorded to qualitatively determine theelements present. The atomic concentrations of carbon, oxygen, silicon and iron weredetermined from the relative peak areas calculated by numerical integration of the detailedscans (20-30 eV windows).

Auger electron spectroscopy (AES) was performed at an electron energy of 10 keV,with an electron current of 1 A. Sample sputtering for depth profiling was accomplished witha 3 keV argon ion beam over a 2x2 mm2 raster.


1 O. Sus, Liebig’s Annalen der Chemie 1944, 556, 65. Later a more complex reactionscheme was proposed: H. A. Levine, Am. Chem. Soc., Polym. Prepr. 1969, 101, 377.

2 H. Ito, C. G. Willson, Polym. Eng. Sci. 1983, 23, 1012.3 G. N. Taylor, T. M. Wolf, L. E. Stillwagon, Solid State Technol. 1984, 27, 145.4 S. A. MacDonald, H. Ito, C. G. Willson, Microelectron. Eng. 1983, 1, 269.5 T. M. Wolf, G. N. Taylor, J. Electrochem. Soc. 1984, 131, 1664.6 J. W. Labadie, S. A. MacDonald, C. G. Willson, J. Imag. Sci. 1986, 30, 169.7 (a) H. Fujioka, H. Nakajima, S. Kishimura, H. Nagata, Proc. SPIE 1990, 1262, 554.

(b) Y. Yoshida, H. Fujioka, H. Nakajima, S. Kishimura, H. Nagata, J. Photopolym.Sci. Technol. 1991, 4, 497. (c) Y. Yoshida, S. Kubota, H. Koezuka, H. Fujioka, J.Vac. Sci. Technol. B 1994, 12, 1402.

8 O. Nalamasu, F. A. Baiocchi, G. N. Taylor, ACS Symp. Ser. 1989, 412, 189.


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9 (a) S. A. MacDonald, H. Schlosser, H. Ito, N. J. Clecak, C. G. Willson, Chem. Mater.1991, 3, 435. (b) S. A. MacDonald, H. Schlosser, N. J. Clecak, C. G. Willson, J. M. J.Fréchet, Chem. Mater. 1992, 4, 1364.

10 (a) M. Suzuki, K. Saigo, H. Gokan, Y. Ohnishi, J. Electrochem. Soc. 1983, 30, 1962.(b) M. Morita, A. Tanaka, S. Immamura, T. Tamamura, O. Kogure, Jpn. J. Appl.Phys. 1983, 22, 659. (c) E. Reichmanis, G. Smolinski, C. G. Wilkins Jr., Solid StateTechnol. 1985, 28, 130.

11 M. A. Hartney, A. E. Novembre, F. S. Bates, J. Vac. Sci. Technol. B 1985, 3, 1346.12 (a) G. N. Taylor, T. M. Wolf, Polym. Eng. Sci. 1980, 20, 1087. (b) G. N. Taylor, T.

M. Wolf, J. M. Moran, J. Vac. Sci. Technol. 1981, 19, 872.13 N. J. Chou, C. H. Tang, J. Paraszczak, E. Babich, Appl. Phys. Lett. 1985, 46, 31.14 (a) H. G. Graighead, L. M. Schiavone, Appl. Phys. Lett. 1986, 48, 1748. (b) Y.

Ohmura, T. Shiokawa, K. Toyoda, S. Namba, Appl. Phys. Lett. 1987, 51, 1500. (c) M.E. Gross, W. L. Brown, L. R. Harriott, K. D. Cummings, J. Linnros, H. Funsten, J.Appl. Phys. 1989, 66, 1403. (d) P. Hoffmann, H. van dan Bergh, J. Flicstein, G. BenAssayas, J. Gierak, J.-F. Bresse, J. Vac. Sci. Technol. B 1991, 9, 3483. (e) A. A. Avey,R. H. Hill, J. Am. Chem. Soc. 1996, 118, 237. (f) M. T. Reetz, M. Winter, J. Am.Chem. Soc. 1997, 119, 4539.

15 D. M. Manos, D. L. Flamm, “Plasma Etching, An Introduction”, Academic Press, SanDiego, 1988.

16 A. J. van Roosmalen, J. A. G. Baggerman, S. J. H. Brader, “Dry Etching for VSLI”,Plenum Press, New York, 1991.

17 D. L. Flamm, V. M. Donnelly, Plasma Chem. Plasma Proc. 1981, 1, 317.18 J. W. Coburn, E. Kay, Solid State Technol. 1979, 22, 117.19 (a) M. Strobel, S. Corn, C. S. Lyons, G. A. Korba, J. Polym. Sci. A 1987, 25, 1295. (b)

M. Strobel, P. A. Thomas, C. S. Lyons, J. Polym. Sci. A 1987, 25, 3343.20 C. J. Mogab, A. C. Adams, D. L. Flamm, J. Appl. Phys. 1978, 49, 3796. (b) V. M.

Donnelly, D. L. Flamm, W. C. Dautremont-Smith, D. J. Werder, J. Appl. Phys. 1984,55, 242.

21 F. D. Egitto, L. J. Matienzo, H. B. Schreyer, J. Vac. Sci. Technol. A 1992, 10, 3060.22 For a recent review on soft lithography: Y. Xia, G. M. Whitesides, Angew. Chem. Int.

Ed. 1998, 37, 550.23 A. Kumar, G. M. Whitesides, Appl. Phys. Lett. 1993, 63, 2002.24 The escape depth of the electrons depends on the binding energy. For a higher binding

energy, the escape depth will decrease.25 (a) J. W. Coburn, J. Appl. Phys. 1979, 50, 5210. (b) J. W. Coburn, Plasma Chem.

Plasma Process. 1982, 2, 1. (c) G. S. Oehrlein, H. R. Williams, J. Appl. Phys. 1987,62, 662.

26 (a) E. Kim, Y. Xia, G. M. Whitesides, Nature 1995, 376, 581. (b) E. Kim, Y. Xia, G.M. Whitesides, J. Am. Chem. Soc. 1996, 118, 5722. (c) E. Kim, Y. Xia, G. M.Whitesides, Adv. Mater. 1996, 8, 245.

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68

Periodic Organic-Organometallic Microdomain Structures

69

&KDSWHU��

Periodic Organic-Organometallic MicrodomainStructures in Poly(styrene- block -ferrocenyl-

dimethylsilane) Copolymers and Blends withCorresponding Homopolymers #

A series of poly(styrene-block-ferrocenyldimethylsilane) copolymers (SF)with different relative molar masses of the corresponding blocks have beenprepared by sequential anionic polymerization. The bulk morphology of thesepolymers, studied by TEM and SAXS, showed well ordered lamellar andcylindrical domains, as well as disordered micellar structures. Temperaturedependent rheological measurements indicated an order-disorder transition forSF 17/8 (the numbers refer to the relative molar masses in 103 g/mol) between170 °C and 180 °C and an order-order transition for SF 9/19 between 190 °Cand 200 °C. The morphology of binary blends of these diblocks withhomopolymer was also investigated. In the blends the molar mass of thehomopolymer was always less than the molar mass of the matching block.Ordered spheres on a bcc lattice and the double gyroid morphology wereobserved for the blends. A double gyroid morphology was observed in F-richdiblock-homopolymer blends.

5.1 Introduction

Poly(ferrocenylsilanes) are organometallic polymers, consisting of alternatingferrocene and silane units in the main chain. The first synthesis of high molar masspoly(ferrocenylsilanes) was performed by thermal ring-opening polymerization using silicon-bridged ferrocenophanes, by Manners and coworkers.1 Since this successful synthesis a largerange of different, bridged ferrocenophanes were prepared with either symmetric orasymmetric substitution on the bridging atom.2 Later it was found that the ferrocenophanescan also undergo ring-opening polymerization in the solid state initiated by γ-irradiation3 and

# The work in this Chapter has been published: R. G. H. Lammertink, M. A. Hempenius, E. L. Thomas, G. J.Vancso, J. Polym. Sci. Part B, Polym. Phys. 1999, 37, 1009.

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70

in solution by anionic initiators4 or transition metal catalysts.5 The anionic ring-openingpolymerization allows one to synthesize well-defined poly(ferrocenylsilanes) as well as blockcopolymers.6 Organometallic-organic block copolymers are attractive materials, since theseform metal-containing nanostructures upon phase separation.

The morphology of microdomains formed by pure and simple diblock copolymers hasbeen an intensively researched and by now well understood area.7 In neat diblocks three“classic” ordered microphases are usually distinguished. These include alternating lamellae(LAM), hexagonally packed cylinders (HEX) and body-centered-cubic packed spheres(BCC). In addition, some other, more complex microstructures may appear especially nearthe order-disorder transition.8

Blends of block copolymers with a corresponding homopolymer give access to avariety of morphologies without having to synthesize a large series of block copolymers.From theory9 and experiment,10 it became clear that the relative molar mass of the addedhomopolymer compared with the molar mass of the corresponding block (α =MA,homo/MA,diblock) is a crucial parameter for block copolymer-homopolymer blends thatcontrols the morphology. If α > 0.5, the homopolymers have a disordering effect on themicrostructure. However, if α < 0.5, the homopolymers are much shorter than their matchingblock, and they will be more uniformly distributed throughout the domains. Consequently,the homopolymer solubility for low α is much higher compared to high α. Extensive work onbinary blends of polystyrene homopolymers with styrene-isoprene diblocks, or styrene-butadiene diblocks, demonstrated the importance of the relative homopolymer molar massand overall blend composition.10f-i

An important aspect about diblock/homopolymer blends found both by theory andexperiment is that the addition of homopolymer can stabilize new morphologies especiallybetween the lamellar and cylindrical phases. This is attributed to the relief of packingfrustration by the homopolymer.9b In neat diblocks, the junctions are placed on the interfaceand pull the monomers near these junctions thereby stretching them. The addedhomopolymer will preferentially occupy regions in the domains, which would otherwise beoccupied by stretched chains of the corresponding block, so some packing frustration can berelieved.

Most theories on microphase separated morphologies of diblock copolymers deal withconformationally symmetric diblocks. Conformational asymmetry can arise from differenceseither in density or Kuhn length, or both. The asymmetry ratio ε, which characterizes theconformational asymmetry of a system, can be defined by:11

2AA0

2BB0

2B,

2A,

BA

b

b

RR

ff

gg ρρ

ε ==

In this equation f is the volume fraction, Rg the radius of gyration, ρ0 the density of thepure component, and b is the Kuhn segment length. A few theoretical studies have


71

investigated the influence of conformational asymmetry on the order-order and order-disordertransitions in diblock copolymer melts.12 The order-disorder transitions were found to bealmost unchanged compared to conformational symmetric diblocks, but the order-ordertransitions were shifted towards the higher content of the block with the larger ρ0b

2.This chapter describes the microphase-separated structures of styrene-block-

ferrocenylsilane copolymers having molar masses between 25 and 40 kg/mol. In addition,blends of these block copolymers with corresponding homopolymer, either poly(styrene) orpoly(ferrocenylsilane), were studied. In order to prevent macrophase-separation within theblends, the molar mass of the added homopolymer was always less than about half of themolar mass of the corresponding block, and homopolymer was added that corresponded tothe majority phase of the diblock.

5.2 Sequential Anionic Polymerization of Styrene and 1,1’-Dimethylsilylferrocenophane

n-BuLi

Fe

Si

CH3

CH3

n-Bu CH2 CH Li

CHCH2n-BuSiFeCH3

CH3

H

1. THF2. MeOH

Chart 5.1 Sequential anionic polymerization of styrene and 1,1’-dimethylsilylferrocenophane. After theaddition of 1,1’-dimethylsilylferrocenophane to the living polystyrene, THF is added, which allows thepolymerization of the organometallic block. Finally, the anionic polymerization is terminated by the addition ofa few drops of methanol.

Poly(styrene-block-ferrocenyldimethylsilane) copolymers (SF) were synthesized bysequential anionic polymerization (Chart 5.1). Polymerization of styrene in ethylbenzene(~10 wt%) was initiated by n-butyllithium and allowed to proceed for 5 hours. After the

Chapter 5

72

styrene block formation was completed, 1,1’-dimethylsilylferrocenophane was added to thesolution and stirred for 5 minutes. Since 1,1’-dimethylsilylferrocenophane does notpolymerize in ethylbenzene, THF was added to the mixture, allowing the polymerization ofthe organometallic block to proceed. This method prevented reaction of the living polystyrylchains with THF that is known to occur at ambient temperatures.13 After 2 hours, the livingchains were terminated by adding a few drops of degassed methanol. The polymers wereprecipitated in methanol and dried under vacuum. The homopolymers and copolymers werecharacterized by gel permeation chromatography (GPC) and 1H NMR.

The notation used to identify the diblock copolymers of SF included thecorresponding molar masses (in 103 g/mol) of the two blocks, e.g. SF 18/19 is a 18,000 g/molpolystyrene -b- 19,000 g/mol polyferrocenylsilane diblock. Homopolymers are given as hPSand hPF, and blends are assigned by the diblock used and the weight percentage of thehomopolymer added, e.g. SF 18/19 with 9% 7.6 hPS is the SF diblock with 9 weight percentadded homopolystyrene, which has a molar mass of 7,600 g/mol. The sequential anionicpolymerization of styrene and 1,1’-dimethylsilylferrocenophane resulted in nearlymonodisperse neat diblock copolymers, without any trace of terminated polystyrene. Table5.1 shows the molecular characteristics, obtained by 1H NMR and GPC measurements.

Table 5.1 Molecular and morphological characteristics of poly(styrene-block-ferrocenyldimethylsilane)copolymers.

Mn

(kg/mol)

Mw

(kg/mol)

Mw/Mn φPS

(vol%)

φPF

(vol%)

microdomain

structureaNS+NF

b

SF 34/6 39.9 43.5 1.09 88 12 Dis 350

SF 24/7 31.0 32.5 1.05 81 19 Dis 257

SF 21/8 28.9 30.2 1.05 78 22 Dis 235

SF 27/12 38.9 40.6 1.04 73 27 Hex F 305

SF 17/8 25.6 27.3 1.07 71 29 Hex F 198

SF 18/19 36.8 38.8 1.05 53 47 Lam 247

SF 9/19 28.5 30.7 1.08 36 64 Lam 167

SF 12/25 36.8 39.3 1.07 36 64 Lam 216

SF 4/27 31.4 34.6 1.10 16 84 Dis 153

a Morphology according to TEM. Dis = disordered, Hex = hexagonally packed cylinders, Lam = lamellae.b NS and NF were calculated from the molar mass of each block, divided by the molar mass of the corresponding

monomer.


73

5.3 Bulk Morphology of Neat Diblocks

For the samples SF 34/6, SF 24/7, SF 21/8 and SF 4/27 disordered sphericalmorphologies were observed. A cylindrical morphology was observed for samples SF 27/12and SF 17/8, which contained 27 and 29 vol% PF, respectively. Figure 5.1 displays the TEMmicrograph of SF 27/12. As mentioned, due to the higher electron density of theorganometallic block, no staining was required and the organometallic domains appear dark.

Figure 5.1 Bright-field TEM micrograph of SF 27/12 (27 vol% PF) displaying hexagonally packed cylinders ina PS matrix.

0.02 0.04 0.06 0.08 0.10

102

103

104

105

106

Inte

nsity

SF 17/8

SF 27/12

√4

1

√3√7

√9√7√4

1

√3

q (Å-1)

Figure 5.2 SAXS patterns for SF 27/12 and SF 17/8. The arrows indicate the relative peak positions forhexagonally packed cylinders.

Chapter 5

74

The SAXS data confirmed the cylindrical morphology for both SF 27/12 and SF 17/8(Figure 5.2), with relative reflections at qn/q1 = 1, √3, √4, √7 and √9 (√4 is missing due toform factor cancellation). SAXS also provided additional information on the dimensions ofthe phase separated microdomains. The structural characteristics obtained from SAXSspacings, block densities and composition and are shown in Table 5.2. In this table a is theinterdomain distance, L the lamellar thickness, R the radius of the PF cylinders, Σ the area perjunction, and H corresponds to the surface mean curvature.

Figure 5.3 Bright-field TEM micrograph of SF 18/19 (47 vol% PF) showing a lamellar morphology.

0.02 0.04 0.06 0.08 0.1010-3

10-2

10-1

100

101

102

103

104

105

3

3

2

2

1

1

SF 12/25

Inte

nsity

SF 9/19

SF 18/19

32

1

q (Å-1)

Figure 5.4 SAXS patterns for SF 18/19, SF 12/25, and SF 9/19 (47, 64, and 64 vol5 PF, respectively). Thearrows indicate the relative peak positions for a lamellar morphology.

The samples SF 18/19, SF 9/19 and SF 12/25 (53, 36 and 36 vol% PS, respectively)all show a lamellar morphology in TEM. The SAXS pattern for SF 18/19 confirms the


75

lamellar morphology with relative peak positions at qn/q1 = 1, 2, 3, respectively. The domaincharacteristics were calculated from the peak positions (Table 5.2). Comparing SF 18/19 withSF 12/25, one can see that the lamellar spacing for SF 12/25 is slightly higher compared toSF 18/19, although the diblocks have about the same molar mass. Interestingly, SF 12/25 hasa larger volume fraction of the denser PF block than SF 18/19.

Table 5.2 Domain characteristics of microphase separated SF diblocks. In this table a is the interdomaindistance, L is the lamellar thickness, R is the radius of the PF cylinders, Σ is the area per junction, and Hcorresponds to the surface mean curvature.

φPS

(vol%)φPF

(vol%)a

(Å)LPS

(Å)LPF

(Å)R

(Å)Σ

(Å2)H

(1/µm)SF 27/12 73 27 261 71 501 70SF 17/8 71 29 205 58 427 86SF 18/19 53 47 244 129 115 372 0SF 9/19 36 64 214 77 137 402 0SF 12/25 36 64 262 94 168 394 0

5.4 Temperature Induced Transitions

It is known that microphase separation influences the flow behavior of diblockcopolymers to a great extent compared to a homogeneous melt. Rheology has been usedbefore to investigate order-order transitions,8 as well as order-disorder transitions7,14 indiblock copolymers. On the one hand, the microdomains, which exist at temperatures belowthe order-disorder temperature, act as physical crosslinks. On the other hand, in thedisordered state, the material exhibits flow behavior similar to conventional homopolymermelts. Thus the elastic modulus, G’, decreases when the order-disorder transition is crossed.Especially, the frequency dependence of elastic response at low reduced frequencies is highlysensitive for the microdomain morphology. In the ordered state the material behaves towardssolidlike (G’ ~ ω0.4-0.5), whereas it changes to near-terminal, liquidlike behavior in thedisordered state (G’ ~ ω2).14e,15

The rheological behavior of four neat diblocks was investigated including the samplesSF 34/6, SF 17/8, SF 18/19 and SF 9/19. As discussed in the morphology section, thesesamples display a disordered micellar, a cylindrical and two lamellar morphologies,respectively when annealed at 165 °C and quenched to room temperature.

Figure 5.5 A shows the mastercurve for G’ at the reference temperature (170 °C) forSF 34/6. The slope of the G’(ω) was found to be 1.5. As mentioned above, liquid likebehavior typical for homopolymer melts should result in a slope of 2. The difference betweenthe observed value and the theoretical value can be ascribed to compositional fluctuations.7

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76

The slope of G’(ω) in Figure 5.5 A therefore implies that SF 34/6 is disordered throughoutthe temperature range (150 - 200 °C) employed.

Figure 5.5 G’ moduli vs. frequency master curves for (A) SF 34/6, (B) SF 17/8, (C) SF 18/19, and (D) SF 9/19.For all mastercurves, 170 °C was taken as the reference temperature.

In Figure 5.5 B the G’ mastercurve of SF 17/8 is displayed. For temperatures between 150 °Cand 170 °C, relatively high G’ values were found compared to SF 34/6. Furthermore, G’ as afunction of ω has a slope of 0.3. This limiting behavior of the elastic modulus confirms thatSF 17/8 is microphase separated up to 170 °C. Above 180 °C however, the G’ moduluscollapses and a slope of 1.4 is found with frequency, indicating a more terminal-zone, liquidlike behavior. These findings suggest that an order-disorder transition takes place between170 °C and 180 °C for SF 17/8.

The G’ mastercurve for SF 18/19 is shown in Figure 5.5 C. The material displays alamellar morphology when annealed at 165 °C. Relatively high G’ moduli are observed witha slope of 0.6 through the entire temperature range. The limiting flow behavior suggests thatSF 18/19 is ordered at least up to 200 °C.

0.01 0.1 1 10 100 1000 1000010

2

103

104

105

106

A

slope = 1.5

150oC

160oC

170oC

180oC

190oC

200oC

G' (

dyne

s/cm

2 )

ω [rad/s]0.01 0.1 1 10 100 1000 10000

102

103

104

105

106

B

slope = 1.4

slope = 0.3

150oC

160oC

170oC

180oC

190oC

200oC

G' (

dyne

s/cm

2 )ω (rad/s)

0.01 0.1 1 10 100 1000 10000102

103

104

105

106

C

slope = 0.6 150oC

160oC

170oC

180oC

190oC

200oC

G' (

dyne

s/cm

2 )

ω (rad/s)

0.01 0.1 1 10 100 1000 10000102

103

104

105

106

D

slope = 0.5

slope = 0.1

150oC

160oC

170oC

180oC

190oC

200oC

G' (

dyn

es/c

m2 )

ω (rad/s)


77

The rheological curve for SF 9/19 displays a slope of 0.5 up to 190 °C (Figure 5.5 D)indicating that the sample was in the ordered state at these temperatures. At 200 °C however,the G’ modulus becomes higher than that at 190 °C and the slope becomes 0.1. This drasticchange in flow behavior represents an order-order transition. Previously, several groupsfound complex behavior of diblock copolymers near the order-disorder transition.8 Somestudies report similar rheological behavior for diblock copolymers as shown here.8c-f Atransition from the lamellar phase to a double gyroid or to a perforated layer structure uponheating was found to result in a higher value for G’. In addition, the frequency dependencebecame less pronounced.

To further study and identify the nature of the order-order transition, temperature-dependent SAXS measurements were performed. SF 9/19 was heated from 160 to 200 °Cusing the same heating rate that was used in the rheological experiments. Figure 5.6 displaysthe SAXS spectra for SF 9/19 at 30, 160, and 200 °C. No drastic changes can be observedother than a slight peak shift to lower q-values and peak broadening at higher temperatures.TEM images from the samples quenched from 165 and 200 °C, however, do displaydifferences between the two phases. Figure 5.7 shows the micrographs for SF 9/19 afterannealing at 165 °C (A) and followed by subsequent heating to 200 °C (B).

0.02 0.04 0.06 0.08 0.10100

101

102

103

104

105

106

1

1

2

2

2

1

200 oC

160 oC

30 oC

Inte

nsity

q (Å-1)

Figure 5.6 SAXS patterns for SF 9/19 at 30, 160, and 200 °C, showing no significant difference in peakpositions with increasing temperature. From Figure 5.5 D it was concluded that SF 9/19 exhibits an order-ordertransition between 190 and 200 °C.

The TEM images (Figure 5.7) indicate that the morphology at 200 °C consists ofalternating lamellae of PS and PF, in which the PS layers are perforated with PF. The

Chapter 5

78

perforations do not seem to possess long-range order, which might account for the absence ofany additional reflections in SAXS.

Previously, Hajduk et al.8f reported on a similar phase transition from lamellar (LAM)to perforated lamellar (PL) for several diblock systems. In their work as well as in our work,the similarity in SAXS pattern for the two phases (LAM and PL) in unoriented samples didnot allow the identification of these morphologies through scattering methods alone. Theauthors concluded that the PL phase is a long-lived nonequilibrium phase that gets convertedinto the bicontinuous gyroid morphology upon isothermal annealing, a transition that we didnot observe for our polymer, even after annealing for 24 h at 200 °C.

Figure 5.7 Bright-field TEM of SF 9/19 (64 vol% PF). (A) The regular lamellar morphology found for SF 9/19is shown. (B) The perforated lamellar (PL) morphology, present in this sample at 200 °C is displayed.

In conclusion, rheology measurements on neat SF diblocks allowed us to determinean order-disorder transition for SF 17/8 and an order-order transition for SF 9/19. Theseexperiments provide valuable additional information about the phase behavior of thesediblock copolymer materials.


79

5.5 Diblock-Homopolymer Blends

The addition of a relatively large molar mass homopolymer in diblock/homopolymerblends may cause macrophase separation. In order to circumvent this, we addedhomopolymer with a relatively low molar mass compared to the corresponding block (α =MA,homo / MA,diblock < 0.55 for nearly all blends). Homopolymer of the majority componentwas added to the diblock, and by using diblocks of different composition, it was possible tocover a broad overall blend composition, without having to add too much homopolymer(mostly below 30 wt%). In all blends studied, macrophase separation was never observed.

Using two homopoly(styrenes) (hPS) and three homopoly(ferrocenyldimethylsilanes)(hPF) of different molar mass, 30 blends were investigated in total by TEM and SAXS. Table5.3 shows the molecular characteristics of the homopolymers employed.

We will consider the blends for each diblock separately. The neat SF 27/12 diblockcontains 27 vol% PF and displays a cylindrical morphology (see Figure 5.1). Upon addinghomopolystyrene (7.6 hPS) to the diblock, the matrix swells and the morphology remainscylindrical. Adding 25% 7.6 hPS causes a phase transition to a spherical morphology (Figure5.8). The total PF content of the blend is then 20 vol%, which is close to the volume fractionin neat PI-PS diblocks at which the order-order transition from cylinders to spheres occurs(18 vol% PS or 24 vol% PI).16

Table 5.3 Molecular characteristics of poly(styrene) (hPS) and poly(ferrocenyldimethylsilane) (hPF)homopolymers.

Mn

(kg/mol)Mw

(kg/mol)Mw/Mn

4.0 hPS 4.0 4.2 1.067.6 hPS 7.6 8.1 1.06

6.5 hPF 6.5 7.1 1.0911.4 hPF 11.4 12.5 1.0924.2 hPF 24.2 26.2 1.09

For SF 27/12 25% 7.6 hPS up to 7 reflections are obtained in the correspondingSAXS experiment (Figure 5.9). This large number of reflections allows one to distinguishbetween the BCC and SC packing. The PF spheres are arranged in a BCC packing, with a cellparameter of 329 Å and a radius of 95 Å.

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Figure 5.8 Bright-field TEM micrograph on SF 27/12 with 25% 7.6 hPS (20 vol% PF in blend) showing aspherical morphology, in which the spheres are arranged on a BCC lattice.

0.02 0.04 0.06 0.08 0.10102

103

104

105

√5 √7√6√4√3

1

√2

Inte

nsity

q (Å-1)

Figure 5.9 SAXS pattern for SF 27/12 with 25% 7.6 hPS (20 vol% PF in blend). The arrows indicate therelative peak positions for a spherical morphology, in which the spheres are arranged on a BCC lattice.

SF 17/8 contains 29 vol% PF and displays a cylindrical morphology. The addition of13% 4.0 hPS to SF 17/8 results in a transition to a disordered cylindrical morphology. Fromthe rheological measurements we already found that SF 17/8 is close to the order-disordertransition at 165 °C (TODT: 170 °C - 180 °C). Therefore it can be expected that, despite theirequal PF content, SF 17/8 with 13% 4.0 hPS displays a disordered cylindrical morphologywhereas SF 27/12 with 8% 7.6 hPS displays an ordered cylindrical morphology. Thus theshape of the order-disorder transition line curves towards the middle of the phase diagram.This observation is in complete agreement with the findings from rheological measurements,which suggested that an order-disorder transition takes place close to the annealingtemperature.


81

Addition of 6.5 hPF to SF 27/12, causes the minority component (PF cylinders) toswell. Figure 5.10 displays the TEM micrograph of SF 27/12 with 11% 6.5 hPF. This blendshows a cylindrical morphology, with a total blend composition of 66 vol% PS. Previously itwas shown that a PS-PI diblock, blended with homopolymer and a volume fraction of 66%PS, displayed a double gyroid structure.10h Also, when SF 18/19 was blended with 27% hPS,giving a total PS content in the blend of 66% PS, a cylindrical morphology was observed. Itseems that the PS-rich double gyroid phase, if present, exists at a lower PS volume fraction inPS-PF diblocks compared to PS-PI diblocks. Such a shift of the order-order transitions can beexplained by the different conformational asymmetry of the two diblock copolymers.

Figure 5.10 Bright-field TEM micrograph of SF 27/12 with 11% 6.5 hPF (34 vol% PF in blend) showinghexagonally packed cylinders of PF in a PS matrix.

The neat diblock SF 18/19 (47 vol% PF) displays a lamellar morphology (Figure 5.3).The addition of 26% hPF, which results in a total PF content of 60 vol% in the blend, doesnot lead to any phase transition. However, the addition of only 9% hPS (43 vol% PF in blend)to this diblock results in a biphasic morphology comprised of PF cylinders and lamellae(Figure 5.11).

Figure 5.11 Bright-field TEM micrograph of SF 18/19 with 9% 7.6 hPS (43 vol% PF in blend) showing abiphasic morphology, consisting of hexagonally packed cylinders of PF and alternating lamellae.

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The TEM micrograph in Figure 5.11 clearly displays the coexistence of cylinders andlamellae. The upper part of the micrograph shows a hexagonal stacking of PF cylinders. Thelower part (left half) displays a near parallel cut to the lamellae. Biphasic morphologies areusually present in blends in between two pure phases.

The SAXS pattern of SF 18/19 with 9% 7.6 hPS (43 vol% PF), shown in Figure 5.12,confirms the existence of a biphasic morphology. Reflections of both the lamellar andcylindrical structures are present. Although a cylindrical morphology would result in a SAXSpattern with approximately the same peak positions, the relatively large intensity of thereflections from the lamellar morphology (qn/q1 = 1, 2, and 3) indicate the presence of bothstructures.

0.02 0.04 0.06 0.08 0.10102

103

104

105

√73

√32

1

Inte

nsity

q (Å-1)

Figure 5.12 SAXS pattern of SF 18/19 with 9% 7.6 hPS (43 vol% PF in blend). The blend displays a biphasicmorphology, as shown in Figure 5.11. The relatively large intensity of the peaks from lamellar scattering (qn/q1

= 1, 2, and 3) indicates the presence of both the lamellar and the cylindrical phases.

Biphasic morphologies may be present in blends in between two pure phases. Astriking feature in our case is that a PS-rich double-gyroid structure is not observed in ourdiblock-homopolymer blends with PS in the matrix and PF in the two interpenetratingnetworks. If a double-gyroid structure was present, it should exist between the cylindrical andlamellar phases.


83

Figure 5.13 Bright-field TEM micrograph of SF 9/19 with 6% 11.4 hPF (66 vol% PF in blend) displaying the[111] projection (“wagonwheel”) of the double-gyroid structure.

Figure 5.13 shows the morphology observed for the blend SF 9/19 with 6% 11.4 hPF.For this system as well as for SF 9/19 with 12 % 11.4 hPF and SF 12/25 with 6% 24.2 hPF aPF-rich double gyroid structure was observed. The total volume fractions PF in these blendsare 66 and 68 %, respectively. SAXS data confirms the double gyroid structure (Figure 5.14).The three observed reflections are consistent with the Ia 3d space group, showing reflectionsat relative positions of √3, √4, and √10.

0.02 0.04 0.06 0.08 0.10

102

103

104

105

Inte

nsity

SF 9/19 6% 11.4 hPF

SF 12/25 6% 24.2 hPF

√4

√3

√10

√4

√3

q (Å-1)

Figure 5.14 SAXS patterns for SF 9/19 with 6% 11.4 hPF and SF 12/25 with 6% 24.2 hPF (both blends have 66

vol% PF). The arrows indicate the relative peak positions for the Ia 3d space group.

The absence of higher order peaks is difficult to explain as a result of structure factorcancellation. The lattice parameters could be determined from the SAXS data and was foundto be 540 Å for the SF 9/19 6% 11.4 hPF blend, 541 Å for the SF 9/19 12% 11.4 hPF blend

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and 603 Å for the SF 12/25 6% 24.2 hPF blend. Although these are blends, we fitted theirlattice parameters, h, with the molar mass of the diblock in the blend. This resulted in h ~M0.4. Furthermore it is interesting to see that the addition of 6% and 12% 11.4 hPF to SF 9/19leads to double gyroid structures with almost equal dimensions.

Further addition of hPF to SF 9/19 resulted in a cylindrical morphology (Figure 5.15)at 18% 11.4 hPF (70 vol% PF in blend) and a disordered spherical micelle morphology at46% 6.5 hPF (80 vol% PF in blend). An ordered spherical morphology was not observed forthese blends. It must be noted that the ordered spherical structure is observed in a verynarrow composition range, so it may exist between 77 and 80 vol% PF in these blends.

Previously, experiments and theory both pointed out that the value of α determinespredominantly the topology of a diblock/homopolymer blend phase diagram. Recently, Janertand Schick9e studied the phase behavior of diblock/homopolymer blends in the weak tointermediate segregation limit within a mean-field approach. Their calculations were basedon a symmetric diblock (fA = fB = 0.5) and they varied the relative homopolymer molar mass,α, between 0.2 and 3.0.17 Homopolymers were found to have a disordering effect on anymicrostructure for α > 0.5. The increase of the entropy of mixing by this disordering effectcompetes with the ordering effect of relaxation of less favorable conformations by thehomopolymers within the microstructure. The stress relieving effect dominates for smallhomopolymers and low concentrations. Also, in our study the application of relatively lowmolar mass homopolymers in blends with copolymers did not result in macrophaseseparation. The addition of homopolymer was found to be a useful tool to give access to avariety of morphologies.

Figure 5.15 Bright-field TEM micrograph of SF 9/19 with 38% 6.5 hPF (77 vol% PF in blend) showinghexagonally packed PS cylinders in a PF matrix.

5.6 Conclusions

The results from TEM, SAXS and rheology experiments on the neat diblockcopolymers and diblock copolymer/hompolymer blends are combined in a phase diagram(Figure 5.16) for relatively low values of α (<0.55). In this diagram (NS+NF)/T, which is


85

proportional to χN (where χ is the Flory-Huggins interaction parameter), is plotted againstthe overall volume fraction of PF. In a blend, NS+NF was calculated from the molar mass andcomposition of the diblock used in the blend, and the monomer molar mass. Neat diblocksare given as solid symbols and blends are displayed as open symbols.

From Figure 5.16 it is obvious that the phase diagram is asymmetric with respect tovol% PF = 50 %. The order-order transitions seem to be placed along almost straight verticallines in the phase diagram for PF > 50 vol%. For volume fractions of PF below 50 %, theorder-disorder transition line curves towards the middle of the phase diagram. The order-order transitions in the intermediate-strong segregation regime for BCC - HEX - LAM - DG -HEX - BCC are located at PF volume fractions of approximately 0.22, 0.43, 0.65, 0.68 and0.79, respectively.

0 10 20 30 40 50 60 70 80 90 1000.3

0.4

0.5

0.6

0.7

0.8

0.90 10 20 30 40 50 60 70 80 90 100

Dis BCC HEX DG PL LAM

(N S+

N F)/

T

(K-1)

PF vol %

Figure 5.16 Phase diagram for SF diblock-homopolymer blends with α < 0.55. NS+NF indicate the total numberof monomers in the diblock. Order-order transitions are located at approximately 0.22, 0.43, 0.65, 0.68, and 0.79for BCC, HEX, LAM, DG, HEX, and BCC, respectively.

We can compare the location of the order-order transitions with theoreticalpredictions from Ohta and Kawasaki.18 Their transitions for BCC - HEX - LAM - HEX -BCC were symmetrical with respect to the volume fractions and were located at 0.215, 0.355,0.645, and 0.785, respectively. Our order-order transitions are shifted towards higher PFvolume fraction compared to their theoretical results. However, the order-order transitionsdepend strongly on the segregation regime, i.e. strong, intermediate or weak. For instance, the

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transition from cylinders to double gyroid shifts from 0.32 to 0.42 as the segregation becomesweaker (χN decreases from 40 to 12).19

The experimental transitions of Hasegawa et al.16 for SI diblocks were found to beshifted to higher PI volume fractions (0.24, 0.34, 0.39, 0.69, and 0.82 for the BCC - HEX -Tetrapod - LAM - HEX - BCC transitions, respectively). Although the shifts in their order-order transitions to higher PI volume fractions are comparable with our shifts to higher PFfractions, it is interesting to note that the Tetrapod structure was only observed in PI-poordiblocks whereas the DG was only observed in PF-rich diblock/homopolymer blends. A shiftof the order-order transitions is often considered as a result of the conformational asymmetry,present in the diblock. Therefore, it is questionable if the conformational asymmetry alsoaccounts for the absence of the double gyroid phase on one side of the phase diagram.

5.7 Experimental

N,N,N’,N’-Tetramethylethylenediamine (TMEDA), ferrocene, styrene, n-butyllithium(1.6 M in hexanes), dibutylmagnesium (1.0 M in heptane) and dichlorodimethylsilane wereobtained from Aldrich. All solvents were obtained from Merck.

TMEDA was distilled from sodium, heptane was distilled from calciumhydride, andferrocene was Soxhlet extracted with cyclohexane prior to use. 1,1’-Dimethylsilyl-ferrocenophane was prepared as described earlier20 and purified by sequential dissolution inheptane followed by vacuum sublimation. Styrene/ethylbenzene (~10 wt% solution) wasdried on dibutylmagnesium and distilled under vacuum. n-Butyllithium was diluted toapproximately 0.2 M with heptane, which was dried over n-butyllithium and distilled undervacuum. Tetrahydrofuran (THF) was dried over n-butyllithium and distilled under vacuum.

Poly(ferrocenyldimethylsilane) was prepared by anionic ring-opening polymerizationof 1,1’-dimethylsilylferrocenophane in THF, using n-butyllithium as an initiator. The livingchains were terminated after 2 hours by the addition of a few drops of degassed methanol.The polymers were precipitated in methanol and dried under vacuum.

Poly(styrene-block-ferrocenyldimethylsilane) copolymers (SF) were synthesized bysequential anionic polymerization carried out in an Mbraun glovebox purged with prepurifiednitrogen. Polymerization of styrene in ethylbenzene (~10 wt%) was initiated by n-butyllithium and allowed to proceed for 5 hours. After the styrene block formation wascompleted, 1,1’- dimethylsilylferrocenophane was added to the solution and stirred for 5minutes. Since 1,1’-dimethylsilylferrocenophane does not polymerize in ethylbenzene, THFwas added to the mixture, allowing the polymerization of the organometallic block to occur.This method prevented reaction of the living polystyryl chains with THF that is known totake place at ambient temperatures. After 2 hours, the living chains were terminated byadding a few drops of degassed methanol. The polymers were precipitated in methanol anddried under vacuum.


87

The homopolymers and copolymers were characterized by gel permeationchromatography (GPC) and 1H NMR. GPC measurements were carried out in THF usingmicrostyragel columns with pore sizes of 105, 104, 103 and 106 Å (Waters). The instrumentwas equipped with a dual detection system consisting of a differential refractometer (Watersmodel 410) and a differential viscometer (Viskotek model H502). Block ratios werecalculated from 1H NMR peak integrals. 1H NMR spectra were recorded in deuteratedchloroform at 250.1 MHz using a Bruker AC 250 spectrometer. The density ofpoly(ferrocenyldimethylsilane), 1.26 g/cm3, was obtained using a pycnometer.

Transmission electron microscopy experiments were performed on a JEOL 2000 FXelectron microscope, operating at 120 kV. Films of copolymers and copolymer/homopolymerblends (approx. 100 mg) were prepared by slowly casting a THF solution (5 % (wt/vol))during approximately one week. The films were then subsequently dried under vacuum at 50°C overnight, after which they were annealed at 165 °C for one week unless otherwisereported. After annealing the samples were quenched in an icewater bath, to avoidcrystallization of the ferrocenylsilane block and preserve their amorphous morphology. DSCexperiments on annealed and quenched samples did not show any melting endotherm asexpected, since a study on the crystallization kinetics of poly(ferrocenyldimethylsilane)showed that crystallization takes places within 10-20 minutes,21 whereas quenching in icewater only takes several seconds. Thin slices were cut from these films at room temperaturewith a Reichert-Jung Ultracut E microtome using a diamond knife. The slices were collectedon 600 mesh copper grids. Due to the large density difference between the two blocks (1.26g/cm3 vs. 1.05 g/cm3) and due to the large number of electrons in the iron atoms of theorganometallic domains no staining was required for TEM experiments.

Small angle X-ray scattering (SAXS) patterns were obtained at the Time-ResolvedDiffraction Facility (station X12B) at the National Synchrotron Light Source at BrookhavenNational Laboratory (BNL) using a custom-built two-dimensional detector (10 cm x 10 cm,512 x 512 pixels). The optical system provides a doubly focused (spot size, 0.5 mm x 0.5 mmfwhm) monochromatic X-ray beam (bandpass, ~5 x 10-4 ∆λ/λ) with a wavelength of λ = 1.54Å. The sample to detector distance was 245 cm.

Rheological measurements were performed on an Ares (Rheometric Scientific)rheometer, using a plate-plate configuration. The gap between the plates was 0.5 - 0.7 mmand a shear strain of 0.5 % was applied. Shear elastic moduli were measured at frequenciesfrom 0.1 - 100 Hz as a function of temperature between 150 °C to 200 °C with steps of 10 °C.Before each frequency sweep the sample was allowed to equilibrate for 10 minutes. G’ databelow 103 dynes/cm2 were removed, based on the limitations of the stress transducer.Mastercurves were constructed from the isothermal frequency sweeps using the principle oftime-temperature superposition. For all mastercurves, 170 °C was taken as the referencetemperature.

GPC measurements on samples that had been annealed or used in temperaturedependent rheological experiments did not show any sign of degradation.

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1 D. A. Foucher, B. Z. Tang, I. Manners, J. Am. Chem. Soc. 1992, 114, 6246.2 Since the first the ring-opening polymerization of 1,1’-dimethylsilylferrocenophane in

1992, numerous ferrocenylsilane polymers have been synthesized with differentsubstituents on the silicon atom and on the cyclopentadienyl-ring. For example, arylor longer alkyl chains on the silicon atom have been prepared: (a) M. T. Nguyen, A.F. Diaz, V. V. Dementev, K. H. Pannell, Chem. Mater. 1993, 5, 1389. (b) D. A.Foucher, R. Ziembinski, B. Z. Tang, P. M. Macdonald, J. Massey, C. R. Jaeger, G. J.Vancso, I. Manners, Macromolecules 1993, 26, 2878. J. Rasburn, D. A. Foucher, W.F. Reynolds, I. Manners, G. J. Vancso, Chem. Commun. 1998, 843. Dihydrosilanebridged poly(ferrocene): J. K. Pudelski, R. Rulkens, D. A. Foucher, A. J. Lough, P.M. Macdonald, I. Manners, Macromolecules 1995, 28, 7301. Silicon substituted withalkoxy, aryloxy, and amino: (c) P. Nguyen, A. J. Lough, I. Manners, Macromol.Rappid Commun. 1997, 18, 953. (d) P. Nguyen, G. Stojcevic, K. Kubalka, M. J.Maclachlan, X. H. Liu, A. J. Lough, I. Manners, Macromolecules 1998, 31, 5977.Silicon substituted with chlorine: (e) D. L. Zechel, K. C. Hultzsch, R. Rulkens, D.Balaishis, Y. Z. Ni, J. K. Pudelski, A. J. Lough, I. Manners, Organometallics 1996,15, 1972. Methylated cyclopentadienyl ring: (f) J. K. Pudelski, D. A. Foucher, C. H.Honeyman, A. J. Lough, I. Manners, S. Barlow, D. O’Hare, Organometallics 1995,14, 2470. (g) J. K. Pudelski, D. A. Foucher, C. H. Honeyman, P. M. Macdonald, I.Manners, S. Barlow, D. O’Hare, Macromolecules 1996, 29, 1894.

3 J. Rasburn, R. Petersen, T. Jahr, R. Rulkens, I. Manners, G. J. Vancso, Chem. Mater.1995, 7, 871.

4 (a) R. Rulkens, A. J. Lough, I. Manners, J. Am. Chem. Soc. 1994, 116, 797. (b) R.Rulkens, Y. Ni, I. Manners, J. Am. Chem. Soc. 1994, 116, 12121.

5 (a) Y. Ni, R. Rulkens, J. K. Pudelski, I. Manners, Macromol.Chem. Rapid Commun.1995, 16, 637. (b) P. Gómez-Elipe, P. M. Macdonald, I. Manners, Angew. Chem. Int.Ed. Engl. 1997, 36, 762.

6 (a) Y. Ni, R. Rulkens, I. Manners, J. Am. Chem. Soc. 1996, 118, 4102. (b) J. Massey,K. N. Power, I. Manners, M. A. Winnik, J. Am. Chem. Soc. 1998, 120, 9533.

7 F. S. Bates, J. H. Rosedale, G. H. Fredrickson, J. Chem. Phys. 1990, 92, 6255.8 (a) K. Almdal, K. A. Koppi, F. S. Bates, K. Mortensen, Macromolecules 1992, 25,

1743. (b) D. A. Hajduk, P. E. Harper, S. M. Gruner, C. C. Honeker, G. Kim, E. L.Thomas, L. J. Fetters, Macromolecules 1994, 27, 4063. (c) S. Förster, A. K.Khandpur, J. Zhao, F. S. Bates, I. W. Hamley, A. J. Ryan, W. Bras, Macromolecules1994, 27, 6922. (d) A. K. Khandpur, S. Förster, F. S. Bates, I. W. Hamley, A. J. Ryan,W. Bras, K. Almdal, K. Mortensen, Macromolecules 1995, 28, 8796. (e) M. F.Schulz, A. K. Khandpur, F. S. Bates, K. Almdal, K. Mortensen, D. A. Hajduk, S. M.Gruner, Macromolecules 1996, 29, 2857. (f) D. A. Hajduk, H. Takenouchi, M. A.Hillmyer, F. S. Bates, M. E. Vigild, K. Almdal, Macromolecules 1997, 30, 3788.

9 (a) M. D. Withmore, J. Noolandi, Macromolecules 1985, 18, 2486. (b) M. W. Matsen,Macromolecules 1995, 28, 5765. (c) H. Xi, S. T. Milner, Macromolecules 1996, 29,2404. (d) A. E. Likhtman, A. N. Semenov, Macromolecules 1997, 30, 7273. (e) P. K.Janert, M. Schick, Macromolecules 1998, 31, 1109.


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10 (a) R. J. Roe, W. C. Zin, Macromolecules 1984, 17, 189. (b) D. J. Kinning, K. I.Winey, E. L. Thomas, Macromolecules 1988, 21, 3502. (c) T. Hashimoto, H. Tanaka,H. Hasegawa, Macromolecules 1990, 23, 43787. (d) H. Tanaka, H. Hasegawa, T.Hashimoto, Macromolecules 1991, 24, 240. (e) H. Tanaka, T. Hashimoto,Macromolecules 1991, 24, 5398. (f) K. I. Winey, E. L. Thomas, L. J. Fetters, J. Chem.Phys. 1991, 95, 9367. (g) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules1991, 24, 6182. (h) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules 1992,25, 422. (i) K. I. Winey, E. L. Thomas, L. J. Fetters, Macromolecules 1992, 25, 2645.(j) R. J. Spontak, S. D. Smith, A. Ashraf, Macromolecules 1993, 26, 956. (k) M. M.Disko, K. S. Liang, S. K. Behal, R. J. Roe, K. J. Jeon, Macromolecules 1993, 26,2983. (l) K. J. Jeon, R. J. Roe, Macromolecules 1994, 27, 2439.

11 E. Helfand, Z. R. Wasserman, In “Developments in Block Copolymers”, I. Goodman,Applied Science, New York, 1982, vol. 1.

12 (a) J. D. Vavasour, M. D. Whitmore, Macromolecules 1993, 26, 7070. (b) M. W.Matsen, M. Schick, Macromolecules 1994, 27, 4014.

13 M. D. Glasse, Prog. Polym. Sci. 1983, 9, 133.14 (a) C. I. Chung, J. C. Gale, J. Polym. Sci. Part B, Polym. Phys. 1976, 14, 1149. (b) E.

V. Gouinlock, R. S. Porter, Polym. Eng. Sci. 1977, 17, 535. (c) C. I. Chung, M. I. Lin,J. Polym. Sci. Part B, Polym. Phys. 1978, 16, 545. (d) C. I. Chung, H. L. Griesbach,L. Young, J. Polym. Sci. Part B, Polym. Phys. 1980, 18, 1237. (e) F. S. Bates,Macromolecules 1984, 17, 2607. (f) K. I. Winey, D. A. Gobran, Z. Xu, L. J. Fetters,E. L. Thomas, Macromolecules 1994, 27, 2392.

15 J. H. Rosedale, F. S. Bates, Macromolecules 1990, 23, 2329.16 H. Hasegawa, H. Tanaka, K. Yamasaki, T. Hashimoto, Macromolecules 1987, 20,

1651.17 Note that the definition of α in the study by Janert and Schick9e (NA/NAB) is half the

value it would have according to our definition (NA,homo / NA,diblock) since a symmetricdiblock (f=0.5) was used. We prefer to use our definition because asymmetricdiblocks are used in blends in this study.

18 T. Ohta, K. Kawasaki, Macromolecules 1986, 19, 2621.19 M. W. Matsen, F. S. Bates, Macromolecules 1996, 29, 1091.20 A. B. Fischer, J. B. Kinney, R. H. Staley, M. S. Wrighton, J. Am. Chem. Soc. 1979,

101, 6501.21 (a) R. G. H. Lammertink, M. A. Hempenius, I. Manners, G. J. Vancso,

Macromolecules 1998, 31, 795. (b) R. G. H. Lammertink, M. A. Hempenius, G. J.Vancso, Macromol. Chem. Phys. 1998, 199, 2141. (c) M. A. Hempenius, R. G. H.Lammertink, G. J. Vancso, Macromol. Symp. 1998, 127, 161.

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Thin Films of PS-b-PFS

91

&KDSWHU��

Morphology and Surface Relief Structures ofAsymmetric Poly(styrene- block -ferrocenylsilane)

Thin Films #

Thin films of an asymmetric styrene-block-ferrocenyldimethylsilane (SF27/12, 27 vol% F) copolymer were studied both on solid silicon substrates aswell as free-standing. In the bulk, the diblock investigated here formspoly(ferrocenylsilane) cylinders in a poly(styrene) matrix. The thin films wereanalyzed by a variety of techniques including atomic force microscopy(AFM), transmission electron microscopy (TEM), x-ray reflectivity,secondary ion mass spectrometry (SIMS) and optical microscopy. From x-rayreflectivity measurements it was concluded that both blocks were present atthe substrate, whereas poly(styrene) wetted the free surface. The formation ofislands and holes were observed by optical microscopy and AFM, as a resultof a mismatch between the film thickness and the equilibrium spacingbetween layers of cylinders. Upon longer annealing of the films, the islandsdisappeared due to a change of the internal morphology of the thin film. Afterselective removal of the organic matrix by means of oxygen reactive ionetching the phase-separated structure could be visualized by AFM. Themorphology of the films altered upon annealing from an in-plane cylindricalto a hexagonal morphology. The structure in freestanding films could beanalyzed by TEM. After annealing, both types of orientations of cylinders(parallel and perpendicular with respect to the substrate) were found to bestable.

6.1 Introduction

In thick block copolymer films, the mismatch between the film thickness and the bulklattice period can be distributed over many layers. As the film thickness is decreased to avalue equal to a few domain periods, the frustration due to the mismatch becomes more

# The work described in this Chapter has been submitted for publication in Macromolecules.

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significant. Such thin films of block copolymers have received much attention in the lastdecade. In order to release the frustration in these films, several changes in morphology havebeen reported for lamellar systems, i.e. symmetric block copolymers. The morphologychange with decreasing thickness has been observed to depend on the specific interactionsthat the corresponding blocks have with their interfaces (either a free surface or a substrate).Three different situations can be distinguished: confined (between two substrates),unconfined (on a substrate), and freestanding thin films (no substrate).

External interfaces influence the nature of ordering in block copolymers. Generally, apreferential affinity of one of the blocks for an interface causes the block to preferentially wetone interface. This results in the ordering of the lamellae parallel to this interface. A blockcopolymer between two confining interfaces becomes frustrated if the confined thickness isincommensurate with the natural spacing of the lamellae. As a result, these films showdifferent periods in films compared to the bulk.1,2,3 However, there is a limit to the extent ofdistortion within the lamellae. Sometimes, even perpendicularly oriented structures have beenobserved.4 In such structures, both blocks are present at the interface. Theoreticalpredictions5,6,7 and simulations8 have supported such observations.

In non-confined films, specific surface segregation results in parallel orientation of thedomains with respect to the substrate. The frustration due to the mismatch between initialfilm thickness and bulk domain period can be released by the formation of islands, i.e.surface relief structures.9 Block selective segregation at the substrate will occur when thewetting component provides the lowest interfacial tension, or when it exhibits a specificaffinity for the substrate. This results in so-called substrate induced ordering,10 which mayeven be present at temperatures above the bulk order-disorder temperature. The surfacesegregation in symmetric block copolymers is believed to be governed by enthalpic factorsand can be tuned to obtain certain specificity for one block or make it neutral to bothblocks.11 However, if the enthalpic driving force is absent, the surface segregation can beentropy-driven resulting in surface segregation of the more flexible chain.12,13 The growth ofsurface relief structures has been studied in detail.14

Substrate-free block copolymer films allow one to study the thin film morphologywith a non-interacting interface. The surface tension of the corresponding blocks will theninduce preferential enrichment of the surface by the block with the lowest surface tension.15

Such effects of interfaces on the thin film morphology have been studied for symmetric blockcopolymer films.16

Asymmetric block copolymer films are extremely interesting if one wants to obtainlaterally structured surfaces. Recently, such films were explored to serve as lithographictemplates.17,18

Theoretical studies on the morphology in thin asymmetric block copolymer films alsoindicate interesting phenomena. In a very recent paper, a dynamic density functional theory(DDFT) was used to simulate microphases in an asymmetric block copolymer confined in athin film.19 Both the film thickness and the polymer-substrate interactions influence themorphologies in the corresponding thin films. Parallel morphologies were dominant if one ofthe blocks had a preference for both substrates. At certain thicknesses, however,


93

perpendicular orientation of cylinders was observed. Also some non-cylindricalmorphologies, like parallel lamellae and perforated lamellae, have been predicted. Anotherstudy concluded that a lamellar region can be formed at the substrate if the affinity differencefor the substrate of the two blocks is larger than some critical value.20 Interacting walls doposses a profound influence on domain formation in thin films.21 Suh et al.22 foundtheoretically that cylinders in triblock and diblock copolymers could orient parallel orperpendicular in confined geometries, depending on the film thickness and interfacialenergies.

Asymmetry in block copolymer composition results in interesting phenomena for thinfilm morphology when compared to symmetric cases. This is easy to envisage, if oneconsiders the case where the minority component wants to preferentially enrich the interface.The microdomain morphology of freestanding thin films of asymmetric diblocks has beenpreviously reported.23 Even if the surface enriching block is present in the minority phase, acontinuous wetting layer of that phase can still be present at the surface. Such a wetting layercan alter the near surface morphology.24,25 Layers of microdomains, cylinders or spheres, thatoriented parallel to the substrate in asymmetric block copolymer films have beenreported.26,27 In the case of a cylindrical morphology, the domains would lie parallel, buttransversely shifted with respect to the domains of the adjacent layer.

Fe

Si

CH3

CH3

H

CHCH2n-Bu

PS-b-PFS : SF

m

n

Chart 6.1 Chemical structure of styrene-block-ferrocenyldimethylsilane copolymer SF.

This chapter describes the thin film morphology of an asymmetric diblock copolymer,poly(styrene-block-ferrocenylsilane) as illustrated in Chart 6.1. The polymer studied was SF27/12 which is a 27,000 g/mol polystyrene -b- 12,000 g/mol polyferrocenylsilane diblock.The block copolymer contains 27 vol% of the organometallic block, and the bulk structureconsists of hexagonally packed cylinders of poly(ferrocenylsilane) in a poly(styrene)matrix.28 Thin films were studied both on silicon substrates, as well as without any substrate(freestanding).

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6.2.1 THIN FILMS ON SILICON SUBSTRATES

The block copolymer used for the thin film study, described in this chapter, possessesa cylindrical morphology in the bulk.28 If thin films are considered, the preferred aggregationof the blocks at the substrate and surface are crucial for the microdomain morphology. In thinasymmetric block copolymer films the domains are, in general, oriented parallel to thesubstrate and surface, if one block has a higher affinity for the particular substrate or surface.Even if the block, that has a preference for the interface, is in the minority, it may still form acontinuous wetting layer at the corresponding interface.24,25

Figure 6.1 Optical micrograph of a dust particle in a thin SF 27/12 film, after annealing for 24 hours at 150 °C.

After annealing SF 27/12 for 24 hours at 150 °C on a silicon substrate, the film formsislands and holes for film thicknesses that are incommensurate with the domain spacing. As itis displayed in Figure 6.1, gradual film thickness variations become “quantized” at specificthicknesses after annealing.29 The formation of plateaus is indicative of an oriented structure,with the nanodomains lying parallel to the film substrate and surface.

The diblock considered in this chapter, SF 27/12, forms a cylindrical phase in the bulkwith a lattice spacing of approximately 26 nm.30 The film forms islands and holes at specificthicknesses depending on which of the phases will be present at the surface and/or thesubstrate. Contact angle measurements performed on the corresponding homopolymers couldnot distinguish between poly(styrene) and poly(ferrocenyldimethylsilane).31 Thus, theobviously low difference in surface energy results in a small driving force for one of the


95

phases to wet the surface. From XPS measurements, however, it was determined thatpoly(styrene) covers the surface of thin PS-b-PFS films after annealing.

At a film thickness of approximately 52 nm (as measured by ellipsometry beforeannealing), which is about equal to two times the lattice spacing, the film forms an almostbicontinuous islands and holes structure after annealing for four days at 160 °C (see Figure6.2). The height of the terraces however, is similar to the lattice spacing of the diblockcopolymer. So, from the corresponding AFM image, we can already conclude that themajority phase, poly(styrene), is not present at both the surface and the substrate. Ifpoly(styrene) was present at both the substrate and the surface, such a structure would exactlyfill a 52 nm thin film with two layers of cylinders (since each layer is about 26 nm thick).XPS measurements indicated that poly(styrene) wets the free surface of the block copolymerfilm. Therefore, the organometallic block is probably present at the substrate. Fromtheoretical studies it was already shown that a lamellar region might be present at thesubstrate, if the difference between the affinity of the two phases for the substrate is largeenough. Otherwise, both phases might be present at the substrate.20

The area occupied by islands is approximately 50% of the total film surface asdisplayed in Figure 6.2. Together with an island height of about 26 nm, this implies that thethickness of the film in a valley is circa 39 nm. A film with an initial thickness of 39 nm doesnot form any islands, but remains continuous after annealing.

To obtain information about the internal microstucture of the thin diblock film, wehave to “zoom in” in order to visualize the phase separated domains. Typically, transmissionelectron microscopy (TEM) is used to study phase separation in block copolymer systems.For these thin films, however, the silicon substrate does not allow the use of TEM since it isnot transparent for electrons. Another possibility for thin film studies is the use of atomicforce microscopy (AFM). The difference in mechanical response of the different domainsallows one to visualize these domains by AFM tapping mode phase imaging. The use ofAFM in imaging block copolymer thin films has proven to be a successful tool to study thelateral morphology on a nanometer level.32

However, the applicability of AFM to study the phase separated morphology of thinpoly(styrene-block-ferrocenylsilane) copolymer films was found to be limited. Due to thesurface coverage of the poly(styrene) phase, the underlying structure could not be visualizeddirectly. If a soft phase would form the matrix of the film, the AFM tip could penetratethrough the covering layer and image the underlying structure. However, this is not possibledue to the glassy nature of poly(styrene).

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Figure 6.2 AFM height image of an annealed 52 nm thin film of SF 27/12 on a silicon substrate. The terracesare approximately 25-27 nm higher compared to the valleys, as can be observed in the line scan on the right.

To overcome this problem, we selectively removed the organic poly(styrene) matrixby means of oxygen reactive ion etching (O2-RIE). In oxygen plasma, organic substances arequickly removed. As we already discussed in Chapter 4, the organometallic polymer forms athin protective oxide layer upon exposure to an oxygen plasma. The removal of thisorganometallic polymer is then slowed down due to the competition between the oxideformation by the plasma and the ion sputtering effect. As a result, a high etching ratedifference, up to 50 : 1, is obtained between the organic and organometallic phases, and theorganic phase can be selectively removed.33

Figure 6.3 displays the AFM images of annealed SF 27/12 thin films of differentthicknesses, after the selective removal of the organic phase by means of O2-RIE. A full layerof in-plane cylinders can be observed for a 40 nm thick film (Figure 6.3 C). The filmthickness is then about equal to the thickness of the film in the valleys from Figure 6.2.Probably the film consists of a single layer of cylinders on top of a thin wetting layer at thesubstrate. If the initial thickness is lower than approximately 40 nm, the single layer ofcylinders can not be completely formed. As can be observed in Figure 6.3 B for a filmthickness of 35 nm, in such cases areas are formed where a hexagonal structure is present.This becomes even more pronounced at a film thicknesses of 30 nm (Figure 6.3 A). Thus themismatch between the initial film thickness and the equilibrium lattice spacing can bebalanced by a change in morphology.


97

Figure 6.3 AFM height images, taken in tapping mode, of thin SF 27/12 films of different thicknesses afterannealing for less than two days at 150 °C. The organic matrix phase has been selectively removed by means ofO2-RIE treatment for 10 s, leaving the organometallic structure behind. (A) 30 nm film tickness, (B) 35 nm filmthickness, (C) 40 nm film thickness.

From the Fourier transforms of the images in Figure 6.3, the lattice spacing could bedetermined. The spacing was found to be 26 nm for image A, 27 nm for image B, and 29 nmfor image C. In the bulk, the cylinders have a lattice spacing of 26 nm, which results in acylinder to cylinder distance of about 30 nm. The spacing from the Fourier transform ofimage A, can be directly related to the spacing from SAXS measurements performed on bulksamples, since both posses a hexagonal structure. Furthermore, the spacing from image C isin quite good agreement with the bulk distance between cylinders. Since image B consists of

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a mixed morphology, the average spacing from the Fourier transform has an intermediatevalue.

Figure 6.4 AFM height images of a 47 nm thin SF 27/12 film after annealing at 150 °C for different times: (A)24 hours, (B) 48 hours, (C) 96 hours.

Provided that the thickness of film is not commensurate with the domain period,islands or holes are formed. This is illustrated in Figure 6.4, for a 47 nm thin film that hasbeen annealed at 150 °C for different times. After 24 hours, the film is still quite smooth, andthe morphology consists of cylinders oriented parallel to the surface. Near the slightlyelevated structures, some defects in the morphology (“dots”) are frequently observed. A well-developed island can be observed in Figure 6.4 B, after annealing for 48 hours. Within theisland and the valleys, the cylinders are oriented parallel to the surface, whereas a sort of


99

corona is present around the island that displays a hexagonal, dot-like structures. Defectstructures near island edges have been previously observed. For example, the steps betweenterraces in symmetric block copolymer films were found to consist of perpendicularlyoriented lamellae.10h In asymmetric diblock films, forming a cylindrical morphology, edgedefects were interpreted as a layer of spheres above a perforated lamella, with the spherespositioned above the perforations.34 Such domain edge structures play an important roleduring the growth and coarsening of islands. Surprisingly, after further annealing, the islandsslowly disappear and the complete film possesses a hexagonal structure (Figure 6.4 C). So,after annealing for more than 3 to 4 days the surface of the film becomes smooth again, witha uniform thickness that is incommensurate with the lattice period.

0 5 10 15 20 25 30 35 40 45 50 55 600.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0.40

0.45

0.50

φ PF

S

Distance from substrate [nm]

Figure 6.5 Volume fraction of PFS plotted against distance from the substrate, obtained from the SIMSspectrum for a 50 nm thin film of SF 27/12, annealed for 20 hours at 150 °C.

Fourier transforms of the AFM images were determined to estimate average repeatperiods. Image A and B have spacings according to Fourier analysis, of approximately 28nm, which is close to the value of 30 nm corresponding to the bulk cylinder to cylinderdistance. Perhaps the cylinders are slightly compressed for image A and B, since the filmthickness is not commensurate with the domain spacing, i.e. the film is too thick. The Fouriertransform of image C indicated a spacing of about 35 nm, which results in a domain todomain distance of approximately 40 nm for a hexagonal structure. The large discrepancybetween the domain to domain distance in the bulk (30 nm) compared to the thin film (40nm) suggests that the structure in image C is not a result of perpendicularly orientedcylinders. Spheres on a BCC lattice have approximately the same spacing for their (110)planes than the spacing of (100) planes in a hexagonal cylindrical morphology.35 Thus, a

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transition from cylinders to spheres near the film surface can not account for the largedeviation in the period spacing.

To gain information about the in-depth morphology of the thin diblock films, x-rayreflectivity and secondary ion mass spectroscopy (SIMS) measurements were performed. Alayered morphology, with its layers parallel to the substrate and surface, gives rise to changesin the density perpendicular to the film. Such changes can be measured by x-ray reflectivityexperiments.

From secondary ion mass spectrometry (SIMS) a first guess about the internal filmstructure in the perpendicular direction can be obtained. The depth profile displays anoscillating signal for iron and silicon, which are present in the organometallic block. At thesubstrate the silicon signal shows a sharp increase due to the silicon substrate. The SIMSresults were converted to volume fraction of PFS as a function of film depth and aredisplayed in Figure 6.5. The volume fraction of PFS oscillates as a result of the alignment ofthe domains parallel to the substrate. The poly(ferrocenylsilane) block is present at the siliconsubstrate. The volume fraction of PFS at the silicon substrate is comparable to the bulkvolume fraction of the diblock (0.27) whereas the volume fraction in the layers, which areoriented parallel to the substrate, exceeds the bulk volume fraction. The distance between thetwo layers is approximately 20 nm, which is less than the period observed in the bulk sample(26 nm).

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

-8

-7

-6

-5

-4

-3

-2

-1

0

1

Exp

Calc

log

(re

flect

ivity

)

q [Å-1]

Figure 6.6 X-ray reflectivity data of a thin SF 27/12 film that has been annealed for 20 hours at 150 °C.

The information obtained from the SIMS experiments can be used as a firstapproximation for the calculation of a density profile from x-ray reflectivity measurements.These measurements were performed on the same films as used for the SIMS experiments. In


101

general, the electron density contrast between different phases in block copolymer thin filmsis quite low, and therefore neutron scattering is usually used to investigate such polymerlayers, of which one block is deuterated. However, deuteration may change the phasebehavior of the polymers. A novel analysis method has been recently developed to obtainaccurate information about the density depth profile of block copolymer thin films utilizingx-rays (i.e. no deuteration is needed).36

If one would be using the well-known Parratt algorithm to obtain the density depthprofile,37 only substrate-film and film-air interfaces could be subtracted from the reflectivitydata in Figure 6.6. By using the new Fourier method,36 which takes into account differentlimits of integration in q-space, the interfacial parameters can be determined with highaccuracy although the difference in the electron density (the contrast) of the two phases isstill extremely small.

Several small peaks are visible besides the large peak around d = 500 Å in Figure 6.7.These small peaks contain the information about the low contrast interface. The simulatedreflectivity and Fourier transferred intensities have been calculated from a refractive indexdepth profile. The initial guess for the iteration process that leads to the final profile wastaken from the SIMS data.

0 100 200 300 400 500 600 700 8004

5

6

7

8

9 Exp

Calc

Fo

urie

r in

tens

ity

d [Å]

Figure 6.7 Fourier transforms of the x-ray reflectivity curves as displayed in Figure 6.6.

The refractive index (related to the electron density) depth profile is shown in Figure6.8. In this figure, the substrate is situated at z = 0, and the film has a total thickness of about50 nm. From this profile, it can be concluded that the organometallic ferrocenylsilane phaseis present at the substrate, whereas the organic styrene phase enriches the surface (as was alsoshown by XPS measurements). This is in full agreement with the observations from AFM,

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which indicated a so-called asymmetric wetting. However, a lamellar wetting layer of theorganometallic phase at the substrate would result in a relatively large volume fraction PFS(close to 1). SIMS and x-ray reflectivity data clearly indicated that this is not the case.Moreover, the results suggested that both phases are located near the substrate. Obviously,the affinity for one of the two phases is not large enough to induce the lamellar ordering atthe substrate, but rather both blocks are in contact with the substrate.20

Interestingly, the thickness of the film is comparable to the thickness of the filmdisplayed in the AFM image in Figure 6.2. So, if a film of this thickness (approximately 50nm) is annealed for only one day, the film remains flat and the internal structure evolves. Thespacing between the two layers in the density profile is approximately 20 nm. This impliesthat the morphology of the film is frustrated, since the equilibrium lattice spacing in the bulkis approximately 26 nm. After longer annealing, i.e. for several days, the film becomes moreand more unsatisfied with its initial thickness, and surface relief structures are formed.

0 100 200 300 400 500 6001.7x10-6

1.8x10-6

1.9x10-6

2.0x10-6

2.1x10-6

2.2x10-6

2.3x10-6

2.4x10-6

2.5x10-6

1 -

Ref

ract

ive

Inde

x

z [Å]

Figure 6.8 Simulated density depth profile (expressed in terms of refractive index variation, inset displays entireprofile), as obtained from x-ray reflectivity data, using the “new” Fourier method from ref. 36.

6.2.2 FREESTANDING DIBLOCK COPOLYMER FILMS

Diblock copolymer thin films can be annealed without the presence of a substrate. Inorder to explore the morphology of such films, a thin film was spin-coated on a silicon wafer,after which the film was removed by using a lift-off technique in a water bath. The floatingcopolymer film was then picked up with a TEM copper grid. During subsequent annealingthe film did not experience contact with any substrate.

0 100 200 300 400 500 6000.0

1.0x10-6

2.0x10-6

3.0x10-6

4.0x10-6


103

In Figure 6.9, a TEM image of a freestanding film of SF 27/12 is displayed. Thesample was annealed for 24 hours at 150 °C. The majority of the film consists of cylindersthat are oriented parallel to the surface, since the majority component, poly(styrene), wetsboth free surfaces (symmetric wetting). A few small areas were observed where theorganometallic cylinders have a perpendicular orientation.

Figure 6.9 TEM image of a freestanding film of SF 27/12, after annealing for 24 hours at 150 °C. Themorphology consists of cylinders of poly(ferrocenylsilane) (dark) in a poly(styrene) matrix (bright). A few smallareas contain perpendicular oriented cylinders.

The orientation of the cylinders within the film can be nicely illustrated by tilting thespecimen in the electron microscope. This is displayed in Figure 6.10. The film displayed inthis figure was annealed for 4 days, which resulted in an increase of the areas with aperpendicular orientation.

Figure 6.10 TEM micrographs of a freestanding SF 27/12 film. Image A was taken without any tilt, showing agrain boundary. Within this grain the domains are hexagonally arranged. Image B was taken at 30° tilt,illustrating that the hexagonally packed domains are perpendicularly oriented cylinders. Furthermore, theparallel oriented cylinders become visible when the sample is tilted.

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Upon longer annealing times, the areas with perpendicularly oriented cylinders wereobserved to become larger. Perpendicularly oriented cylinders have been observed before, butthey were presumably a result of the film preparation, i.e. quick solvent evaporation.38 Inother words, perpendicularly oriented cylinders were found to be metastable thus far, sincethe higher surface free energy block would also be present at the surface.39 The perpendicularorientation in freestanding films of poly(styrene-block-ferrocenyldimethylsilane) seems to bea stable morphology.

An explanation for the perpendicular orientation may be that the difference in surfacefree energy between the two phases is small enough to allow both blocks to be present at thesurface. In doing so, the frustration within the film due to the mismatch between the filmthickness and the domain period can be released. From contact angle measurements thesurface energies of poly(styrene) and poly(ferrocenyldimethylsilane) were found to beapproximately equal within the accuracy of the measurements. In a theoretical study, it wasmentioned that if the difference between tendency for surface segregation is below a certainvalue, perpendicular orientation can be favorable, since it releases the frustration by allowingboth blocks at the same interface.22 For a perpendicular orientation of cylinders, theequilibrium period of the block copolymer system can be realized, whereas this is notpossible in the parallel morphology without the formation of islands or holes.

6.3 Conclusions

The thin film morphology of an asymmetric styrene-block-ferrocenyldimethylsilanecopolymer was studied. The microdomain structure could be visualized by AFM afterselective removal of the organic matrix by means of oxygen reactive ion etching. The thinfilm morphology after annealing consists of cylinders that are oriented parallel to the surface.Upon annealing, the films form surface relief structures if the film thickness isincommensurate with the equilibrium spacing period. From SIMS and x-ray reflectivitymeasurements, it was concluded that both poly(styrene) and poly(ferrocenylsilane) arepresent at the silicon substrate whereas poly(styrene) preferentially wets the free surface.After longer annealing periods, usually more than four days, the surface relief structuresdisappeared due to a rearrangement of the microdomains.

For freestanding films both orientations of the cylinders were found to be present inannealed samples. The perpendicular orientation of cylinders appears to be stable as well,since the amount of perpendicularly oriented cylinders is increasing in time during annealing.The perpendicular orientation allows the film thickness to be incommensurate with theequilibrium spacing while maintaining the equilibrium period, so the stress can be relievedwithout the formation of surface relief structures (islands or holes).


105

6.4 Experimental

N,N,N’,N’-Tetramethylethylenediamine (TMEDA), ferrocene, styrene, n-butyllithium(1.6 M in hexanes), dibutylmagnesium (1.0 M in heptane) and dichlorodimethylsilane wereobtained from Aldrich. All solvents were obtained from Merck.

TMEDA was distilled from sodium, heptane was distilled from calciumhydride, andferrocene was Soxhlet extracted with cyclohexane prior to use. 1,1’-Dimethylsilyl-ferrocenophane was prepared as described earlier and purified by sequential dissolution inheptane followed by vacuum sublimation. Styrene/ethylbenzene (~10 wt% solution) wasdried on dibutylmagnesium and distilled under vacuum. n-Butyllithium was diluted toapproximately 0.2 M with heptane that was dried over n-butyllithium and distilled undervacuum. Tetrahydrofuran (THF) was dried over n-butyllithium and distilled under vacuum.

Poly(styrene-block-ferrocenyldimethylsilane) copolymers (SF) were synthesized bysequential anionic polymerization carried out in an Mbraun glovebox purged with prepurifiednitrogen. Polymerization of styrene in ethylbenzene (~10 wt%) was initiated by n-butyllithium and allowed to proceed for 5 hours. After the styrene block formation wascompleted, 1,1’- dimethylsilylferrocenophane was added to the solution and stirred for 5minutes. Since 1,1’-dimethylsilylferrocenophane does not polymerize in ethylbenzene, THFwas added to the mixture, allowing the formation of the organometallic block. This methodprevented reaction of the living polystyryl chains with THF that is known to occur at ambienttemperatures. After 2 hours, the living chains were terminated by adding a few drops ofdegassed methanol. The polymers were precipitated in methanol and dried under vacuum.

GPC measurements were carried out in THF using microstyragel columns with poresizes of 105, 104, 103 and 106 Å (Waters). The instrument was equipped with a dual detectionsystem consisting of a differential refractometer (Waters model 410) and a differentialviscometer (Viskotek model H502). Block ratios were calculated from 1H NMR peakintegrals. 1H NMR spectra were recorded in deuterated chloroform at 250.1 MHz using aBruker AC 250 spectrometer. The density of poly(ferrocenyldimethylsilane), 1.26 g/cm3, wasobtained by means of a pycnometer.

Thin films were prepared by spin-coating on silicon wafers from toluene solutions. Byvarying the spinning speed and concentration of the copolymer solution the film thicknesscould be controlled. The film thickness was measured by ellipsometry prior to annealing. Thefilms were annealed under vacuum at 150-160 °C.

The organic poly(styrene) matrix could be selectively removed by means of oxygenreactive ion etching. Oxygen reactive ion etching (O2-RIE) experiments were carried out inan Elektrotech PF 340 apparatus. The pressure inside the etching chamber was 10 mTorr, thesubstrate temperature was set at 10 °C, and an oxygen flow rate of 20 cm3/min wasmaintained. The power was set at 75 W. Etching times of only 10 seconds were found to beenough to expose the organometallic domains, so they could be visualized by AFM.

The morphology of unetched and etched films was studied by atomic forcemicroscopy (AFM) operating in the tapping mode. AFM experiments were performed using a

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106

NanoScope III A instrument (Digital Instruments) utilizing NanoSensors tapping tips. Etchedfilms could be studied best at low amplitudes of oscillation at free vibration (A0 = 1.0 V) andrelatively high operating setpoint ratios (A/A0 ~ 0.9). Height and phase images were recordedat a scanning rate of 1 Hz.

Freestanding films were prepared by spin-coating an approximately 70 nm thin filmonto a silicon wafer. The film was floated off in a water bath and picked up with a copperTEM grid. The films were then annealed under vacuum at 150 °C.

Transmission electron microscopy (TEM) images of freestanding films wereperformed on a JEOL 200 CX and Philips CM30 electron microscope. The sample holderscould be tilted. No staining was required, due to the larger electron density of theorganometallic phase.

X-ray reflectivity measurements were carried out at the beamline X10B of theNational Synchrotron Light Source at Brookhaven National Laboratory using a wavelengthof 1.127 Å. Analysis of the reflectivity data was performed using a new Fourier method,which allows one to evaluate even low contrast layered systems.36

Depth profiling was performed on a SIMS Atomika instrument with an Ar beam at 3keV rastered over a 1x1 mm2 spot. The samples were covered by a sacrificial thin film ofdeuterated poly(styrene) by floating it onto the diblock film. This ensured that steady-statesputtering conditions were achieved before the sputtering crater reached the originalcopolymer surface. Positive Fe and Si and negative deuterium and carbon were monitored asa function of sputtering time. The thickness of the film was measured with ellipsometry andthe depth scale was obtained by dividing the carbon trace by the total thickness.


1 P. Lambooy, T. P. Russell, G. J. Kellogg, A. M. Mayes, P. D. Gallagher, S. K. Satija,Phys. Rev. Lett. 1994, 72, 2899.

2 N. Koneripalli, N. Singh, R. Levicky, F. S. Bates, P. D. Gallagher, S. K. Satija,Macromolecules 1995, 28, 2897.

3 N. Koneripalli, R. Levicky, F. S. Bates, J. Ankner, H. Kaiser, S. K. Satija, Langmuir1996, 12, 6681.

4 G. J. Kellogg, D. G. Walton, A. M. Mayes, P. Lambooy, T. P. Russell, P. D.Gallagher, S. K. Satija, Phys. Rev. Lett. 1996, 76, 2503.

5 D. G. Walton, G. J. Kellogg, A. M. Mayes, P. Lambooy, T. P. Russell,Macromolecules 1994, 27, 6225.

6 G. T. Pickett, A. C. Balazs, Macromolecules 1997, 30, 3097.7 M. W. Matsen, J. Chem. Phys. 1997, 106, 7781.8 (a) M. Kikuchi, K. Binder, Europhys. Lett. 1993, 21, 427. (b) M. Kikuchi, K. Binder,

J. Chem. Phys. 1994, 101, 3367.9 G. Coulon, D. Auserre, T. P. Russell, J. Phys. (France) 1990, 51, 777.


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10 Surface-induced ordering in thin films has been studied by several techniques, such asneutron reflectivity, secondary ion mass spectroscopy (SIMS) and transmissionelectron microscopy (TEM). (a) S. H. Anastasiadis, T. P. Russell, S. K. Satija, C. F.Majkrzak, Phys. Rev. Lett. 1989, 62, 1852. (b) S. H. Anastasiadis, T. P. Russell, S. K.Satija, C. F. Majkrzak, J. Chem. Phys. 1990, 92, 5677. (c) A. Menelle, T. P. Russell,S. H. Anastasiadis, S. K. Satija, C. F. Majkrzak, Phys. Rev. Lett. 1992, 68, 67. (d) A.M. Mayes, T. P. Russell, P. Bassereau, S. M. Baker, G. S. Smith, Macromolecules1994, 27, 749. (e) N. Gadegaard, K. Almdal, N. B. Larsen, K. Mortensen, Appl. Surf.Sci. 1999, 142, 608. A selection of papers on thin block copolymer films, that havebeen studied by SIMS: (f) G. Coulon, T. P. Russell, V. R. Deline, P. F. Green,Macromolecules 1989, 22, 2581. (g) T. P. Russell, G. Coulon, V. R. Deline, D. C.Miller, Macromolecules 1989, 22, 4600. TEM can be used to elucidate themorphology within thin films: (h) B. L. Carvalho, E. L. Thomas, Phys. Rev. Lett.1994, 73, 3321. (i) B. L. Carvalho, R. L. Lescanec, E. L. Thomas, Macromol. Symp.1995, 98, 1131. (j) Y. Liu, M. H. Rafailovich, J. Sokolov, S. A. Schwarz, S. Bahal,Macromolecules 1996, 29, 899.

11 (a) P. Mansky, T. P. Russell, C. J. Hawker, M. Pitsikalis, J. Mays, Macromolecules1997, 30, 6810. (b) P. Mansky, T. P. Russell, C. J. Hawker, J. Mays, D. C. Cook, S.K. Satija, Phys. Rev. Lett. 1997, 79, 237.

12 (a) M. D. Foster, M. Sikka, N. Singh, F. S. Bates, S. K. Satija, C. F. Majkrzak, J.Chem. Phys. 1992, 96, 8605. (b) M. Sikka, N. Singh, A. Karim, F. S. Bates, S. K.Satija, C. F. Majkrzak, Phys. Rev. Lett. 1993, 70, 307.

13 A theoretical analysis of blends of stiff and flexible chains agreed with theexperimental observations that the surface is enriched by the more flexible chain: G.H. Frederickson, J. P. Donley, J. Chem. Phys. 1992, 97, 8941. However, simulationshave pointed out that surface segregation can be entropy driven, resulting in surfaceenrichment by the stiffer polymer: (a) A. Yethiraj, S. Kumar, A. Hariharan, K. S.Schweizer, J. Chem. Phys. 1994, 100, 4691. (b) S. K. Kumar, A. Yethiraj, K. S.Schweizer, F. A. M. Leermakers, J. Chem. Phys. 1995, 103, 10332. (c) A. Yethiraj,Phys. Rev. Lett. 1995, 74, 2018.

14 (a) G. Coulon, B. Collin, D. Ausserre, D. Chatenay, T. P. Russell, J. Phys. (France)1990, 51, 2801. (b) M. Maaloum, D. Ausserre, D. Chatenay, G. Coulon, Y. Gallot,Phys. Rev. Lett. 1992, 68, 1575. (c) B. Collin, D. Chatenay, G. Coulon, D. Ausserre,Y. Gallot, Macromolecules 1992, 25, 1621. (d) Z. Cai, K. Huang, P. A. Montano, T.P. Russell, J. M. Bai, G. W. Zajac, J. Chem. Phys. 1993, 98, 2376. (e) R. Mutter, B.Stühn, Macromolecules 1995, 28, 5022. (f) P. C. M. Grim, I. A. Nyrkova, A. N.Semenov, G. ten Brinke, G. Hadziioannou, Macromolecules 1995, 28, 7501. (g) G.Vignaud, A. Gibaud, G. Grübel, S. Joly, D. Ausserre, J. F. Legrand, Y. Gallot, Phys.B 1998, 248, 250. (h) J. Heier, E. Sivaniah, E. J. Kramer, Macromolecules 1999, 32,9007.

15 A. K. Rastogi, L. E. St. Pierre, J. Coll. Inter. Sci. 1969, 31, 168.16 Y. Liu, M. H. Rafailovich, J. Sokolov, S. A. Schwarz, S. Bahal, Macromolecules

1996, 29, 899.17 M. Park, C. Harrison, P. M. Chaikin, R. A. Register, D. H. Adamson, Science 1997,

276, 1401.

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18 (a) R. G. H. Lammertink, M. A. Hempenius, J. E. van den Enk, V. Z.-H. Chan, E. L.Thomas, G. J. Vancso, Adv. Mater. 2000, 12, 98. (b) R. G. H. Lammertink, M. A.Hempenius, G. J. Vancso, V. Z.-H. Chan, E. L. Thomas, Abstract ACS, PMSEPreprint, 1999, 81, 14.

19 H. P. Huinink, J. C. M. Brokken-Zijp, M. A. van Dijk, G. J. A. Sevink, J. Chem. Phys.2000, 112, 2452.

20 M. S. Turner, M. Rubinstein, C. M. Marques, Macromolecules 1994, 27, 4986.21 G. Brown, A. Chakrabarti, J. Chem. Phys. 1994, 101, 3310.22 K. Y. Suh, Y. S. Kim, H. H. Lee, J. Chem. Phys. 1998, 108, 1253.23 L. H. Radzilowski, B. L. Carvalho, E. L. Thomas, J. Polym. Sci. Part B, Polym. Phys.

1996, 34, 3081.24 H. Yokoyama, E. J. Kramer, M. H. Rafailovich, J. Sokolov, S. A. Schwarz,

Macromolecules 1998, 31, 8826.25 Y. Liu, W. Zhao, X. Zheng, A. King, A. Singh, M. H. Rafailovich, J. Sokolov, K. H.

Dai, E. J. Kramer, S. A. Schwarz, O. Gebizlioglu, S. K. Sinha, Macromolecules 1994,27, 4000.

26 A. Karim, N. Singh, M. Sikka, F. S. Bates, W. D. Dozier, G. P. Felcher, J. Chem.Phys. 1994, 100, 1620.

27 C. S. Henkee, E. L. Thomas, L. J. Fetters, J. Mater. Sci. 1988, 23, 1685.28 R. G. H. Lammertink, M. A. Hempenius, E. L. Thomas, G. J. Vancso, J. Polym. Sci.

Part B, Polym. Phys. 1999, 37, 1009.29 The use of interference microscopy to investigate the free surface topology of thin

copolymer films has been presented previously: see ref. 9.30 See chapter 5 and ref. 28.31 The contact angles of liquids with the poly(styrene) and poly(ferrocenylsilane)

surfaces were approximately 90° for water, 55° for diiodomethane, and 78° glycerol.Using the acid-base interaction theory (C. J. van Oss, L. Ju, M. K. Chaudhury, R. J.Good, J. Coll. Interface Sci. 1989, 128, 313), this resulted in a surface energy γ = 33mN/m.

32 Ph. Leclère, R. Lazzaroni, J. L. Brédas, J. M. Yu, Ph. Dubois, R. Jérôme, Langmuir1996, 12, 4317.

33 See Chapter 4 and ref. 18.34 C. Harrison, M. Park, P. Chaikin, R. A. Register, D. H. Adamson, N. Yao,

Macromolecules 1998, 31, 2185.35 (a) K. A. Koppi, M. Tirrell, F. S. Bates, J. Rheol. 1994, 38, 999. (b) S. Sakurai, T.

Hashimoto, L. J. Fetters, Macromolecules 1996, 29, 740.36 O. H. Seeck, I. D. Kaendler, M. Tolan, K. Shin, M. H. Rafailovich, J. Sokolov, R.

Kolb, Appl. Phys. Lett., in press.37 L. G. Parratt, Phys. Rev. 1954, 95, 359.38 (a) G. Kim, M. Libera, Macromolecules 1998, 31, 2569. (b) G. Kim, M. Libera,

Macromolecules 1998, 31, 2670.39 P. Mansky, P. Chaikin, E. L. Thomas, J. Mater. Sci. 1995, 30, 1987.

Thin Films of PI-b-PFS

109

&KDSWHU��

Morphology and Crystallization of Thin Films ofAsymmetric Organic-Organometallic Diblock Copolymers

of Isoprene and Ferrocenyldimethylsilane #

The morphology of thin films of asymmetric block copolymers ofpoly(isoprene-block-ferrocenyldimethylsilane) was studied using atomic forcemicroscopy, transmission electron microscopy, and optical microscopy. Blockcopolymers with the organometallic (ferrocenylsilane) phase between 20 and28 vol% were investigated. At these compositions the copolymers formedcylindrical structures in the bulk. Thin films, spin-cast on silicon wafers,possessed different morphologies depending on the composition and the filmthickness. The block copolymers exhibited a surface structure consisting of aworm-like pattern for film thicknesses exceeding the domain spacing. Whenthe film thickness matched the inter domain spacing, a surface morphologyconsisting of hexagonally packed domains was obtained for the diblockcontaining 20 vol% of the organometallic phase, and a complete worm-likestructure was observed for the diblock containing 28 vol% of theferrocenylsilane phase. The diblock containing 24 vol% of the organometallicphase displayed a mixed morphology at this thickness consisting of a worm-like structure and hexagonally packed domains. Hole formation, due toincompatibility between the film thickness and the domain spacing, wasobserved for films thinner than a single domain layer. The formation ofislands, or holes, in thicker films was not observed because lattice distortionscan relieve the stress that is generated by excess material. Crystallization ofthe PFS phase took place at room temperature and resulted in large hedriticstructures over the whole surface of the diblock copolymer films. Themicrodomain morphology was completely destroyed in the films consisting ofcrystallized organometallic domains. From DSC measurements on bulksamples, the melting temperature and enthalpy of the diblock copolymersindicated well-advanced crystallization compared to ferrocenyldimethylsilanehomopolymers.

# The work described in this Chapter has been published: R. G. H. Lammertink, M. A. Hempenius, G. J. Vancso,Langmuir 2000, in press.

Chapter 7

110

7.1 Introduction

The use of self-assembly of macromolecules into patterns with typical structure sizesthat are an order of magnitude smaller compared to patterns obtained by conventional opticallithography seems a very promising approach.1,2 These and other applications requireadequate understanding of the morphology and stability of thin block copolymer films.

The presence of a substrate and free surface strongly influence the behavior of blockcopolymers in thin films compared to their behavior in the bulk. In the bulk, the equilibriummorphology is mainly determined by the volume fractions of the corresponding blocks, theFlory-Huggins interaction parameter (χ) and the degree of polymerization.3 Generally,symmetric block copolymers form lamellar structures. With increasing asymmetry, themorphology exhibits cylinders and spheres of one phase embedded in the matrix of the otherphase, and co-continuous morphologies. In thin films, the presence of a substrate and surfacecan induce orientation of the microphase structure,4 and can result in changes in domaindimensions or phase transitions5,6 due to segregation of one of the blocks at the substrate orsurface.

The ordering in thin films of asymmetric block copolymers, with spherical orcylindrical morphologies has been investigated to a much smaller extent compared tosymmetric block copolymers. In one theoretical study, the surface induced ordering of acylindrical phase was treated.7 The authors found that a lamellar region (covering layer) atthe surface existed if the surface tension difference between the two phases and the surface isabove a certain critical value. Suh et al.8 found theoretically that cylinders in triblock anddiblock copolymers could orient parallel or perpendicular in confined geometry, dependingon the film thickness and surface tensions. Recently, dynamic density functional theory wasused to study microdomain formation in confined thin films of asymmetric blockcopolymers.9 It was concluded that the film thickness and the polymer-substrate interactionare the main parameters that govern thin film morphologies of asymmetric block copolymers.

Parallel morphologies were found to be dominant when one of the blocks had apreference for the substrates. However, at certain thicknesses perpendicularly orientedcylinders were observed. In addition, non-cylindrical morphologies, such as parallel lamellaeand perforated lamellae were observed for films between selective substrates. Experimentsshow that a strong preference of one of the blocks for the substrate leads to preferentialwetting of the substrate and therefore in an orientation of the domains parallel to thesubstrate.10,11 When there was no clear affinity for one of the two blocks to wet the substrate,a parallel orientation was observed for spherical and cylindrical domains in a single domainlayer.12 In films of asymmetric poly(styrene-block-butadiene) copolymers, annealed withoutthe presence of a substrate, a transition from the bulk cylindrical morphology was observedas the film thickness decreased.13 This was induced by surface segregation of the lowersurface tension PB component. The morphology of spherical and cylindrical PS-b-PB andPS-b-PI copolymer thin films on silicon wafers was investigated by Harrison et al.14 In asingle domain layer, the cylinders oriented parallel to the substrate, which was wetted by alayer of PB. The wetting of the silicon by the PB was ascribed to the lower interfacial tension


111

between unsaturated hydrocarbons and polar substrates, compared to aromatic hydrocarbonsand polar substrates.15 In a monolayer of spheres, the domains were arranged in a hexagonallattice, oriented parallel to the substrate. Islands and holes were observed for both sphericaland cylindrical morphologies, if the initial film thickness was not commensurable with thedomain spacing. Lattice distortions in the bulk of the films, however, could prevent thecreation of islands and holes.16

In this study, AFM observations of the morphology of thin films of asymmetricpoly(isoprene-block-ferrocenyldimethylsilane) copolymers on a silicon substrate aredescribed. Recently, thin films of these polymers that form lateral structures on solidsubstrates were found to be interesting candidates to function as lithographic templates in aone-step reactive ion etching process.2 The composition of the diblock copolymers studiedranged between 20 vol% to 28 vol% for the organometallic phase.


7.2.1 ANIONIC SYNTHESIS OF POLY (ISOPRENE-BLOCK-FERROCENYLSILANES )

The poly(isoprene-block-ferrocenyldimethylsilane) copolymers used for this studywere synthesized by sequential anionic polymerization of isoprene and 1,1’-dimethylsilyl-ferrocenophane. Isoprene was polymerized in toluene, which results in a relatively high(~94%) amount of 1,4 addition, according to the 1H NMR spectra (not shown).17

SiFe

n-BuLi

Fe

Si

CH3

CH3

H

1. THF

2. MeOH

n-Bu

n-Bu

CH3

CH3

Li

nm

Chart 7.1 Sequential anionic polymerization of isoprene and 1,1’-dimethylsilylferrocenophane.

Chapter 7

112

The molecular characteristics of the polymers synthesized are given in Table 7.1. Thediblocks are denoted as IF referring to isoprene and ferrocenyldimethylsilane respectively,followed by their molar masses in kg/mol.

Table 7.1 Molecular characteristics of poly(isoprene-block-ferrocenylsilane) copolymers (PI-b-PFS).

Mn (kg/mol) Mw (kg/mol) Mw/Mn PFS vol%IF 29/15 44.1 46.3 1.05 28IF 31/13 44.9 47.2 1.05 24IF 36/12 48.4 49.5 1.02 20

Mn number average molar mass, Mw weight average molar mass, Mw/Mn polydispersity index

7.2.2 THIN FILM MORPHOLOGY AND STABILITY

Figure 7.1 Optical micrograph of an isoprene-block-ferrocenylsilane copolymer thin film, after annealing at 150°C.

Thin films of PI-b-PFS diblock copolymers, spin-cast from toluene, are smooth andcompletely cover the silicon substrate. Figure 7.1 shows an optical micrograph of a film ofPI-b-PFS after annealing for 20 hours at 150 °C. The block copolymer dewetted the substrateduring heat treatment, and exhibited patterns very similar to the final stage of dewetting ofhomopolymer thin films.18,19 The patterns were also similar to thin block copolymer filmsthat have been annealed above the orded-disorder transition temperature.20 The bulk order-


113

disorder transition temperatures of the PI-b-PFS copolymers used in this study were between120 and 130 °C, as determined by rheological experiments. Since annealing below 120 °Cmight result in crystallization of the PFS phase,21 we studied the morphology of unannealeddiblock films at room temperature. Although unannealed films display morphologies that areprobably in non-equilibrium states, they give valuable information about the formation oflateral nanodomains. Following spin-casting we observed that the order of the morphologyincreased upon storage at room temperature. PI is a viscoelastic “liquid” polymer at roomtemperature and the glass transition temperature of the PFS block was approximately 18 °C,as determined from DSC experiments. The decrease of the glass transition temperature of thePFS block, compared to the bulk value for PFS homopolymer (around 25 °C for a similarmolar mass),21 can be ascribed to the partial mixing of the two blocks at the microdomaininterface. Probably some mobility within the thin diblock film is still present to allow forsmall-scale rearrangements resulting in increased order.

AFM phase images22 of relatively thick films (> 100 nm) of PI-b-PFS copolymers aredisplayed in Figure 7.2. Due to the lower surface free energy of the PI blocks, PI segregatesat the surface, which was confirmed by contact angle and XPS measurements.23 IF 29/15 andIF 31/13 display a disordered worm-like morphology (Figure 7.2 A and B). A thick IF 36/12film, containing only 20 vol% of the organometallic block, also forms a worm-likemorphology, along with some areas containing dots (Figure 7.2 C). In a previous paper onpoly(styrene-block-ferrocenylsilane) copolymers and blends with correspondinghomopolymers, we obtained a transition from cylinders to spheres in the bulk at acomposition of 22 vol% of the organometallic block.24 Transmission electron microscopyexperiments on freestanding isoprene-block-ferrocenylsilane copolymer films, however,indicated a cylindrical structure for all three diblocks investigated.

In the images captured in Figure 7.2 A-C the Fourier transform of the image intensityindicated a repeat distance of the domains of approximately 33 nm for IF 29/15 and IF 31/13,and 31 nm for IF 36/12. Thus it was of interest to study films with thicknesses close to thisvalue, as one would anticipate a frustrated behavior for these films. AFM phase images ofthin films, approximately 30 nm, are displayed in Figure 7.3. For IF 29/15 an in-plane worm-like structure can still be observed (Figure 7.3 A). Perpendicular orientation for films of IF29/15 has not been observed. A 30 nm thin film of IF 31/13 displays a mixed morphology,consisting of in-plane cylinders and circular domains that are hexagonally arranged (Figure7.3 B) with an interdomain distance of approximately 33 nm, as determined from the Fouriertransform of the image shown in the inset. Films that are slightly thicker than the one shownin Figure 7.3 B, display a pure in-plane cylindrical morphology.

Chapter 7

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Figure 7.2 Tapping mode AFM phase images of relatively thick PI-b-PFS copolymer films (> 100 nm). A. IF29/15 and B. IF 31/13 showing an in-plane worm-like structure. C. IF 36/12 contains areas with circular andworm-like features. The Fourier transforms, with center and side corresponding to a reciprocal distance d-1

equals 0 and 62.5 µm-1 respectively, are displayed in the top right corner of each image.

Harrison et al.14 observed a transition from in-plane cylinders to spheres at edges ofislands in thin PS-b-PB films on silicon. The substrate and surface in their films were wettedby PB, which was present in the cylinders. Due to the PB wetting layer the amount of PBinside the film was reduced and a transition from cylinders to spheres resulted. In our case,the surface is wetted by the PI phase, which is the majority phase. The observed change fromworm-like domains to hexagonally arranged domains in IF 31/13 films can be understood ifone considers that the microstructure is consisting of a single layer of cylinders in the organicmatrix. If the film thickness is reduced, the diameter of the cylinder must decrease whichresults in an increase of the curvature of the cylinders. In order to relieve the resulting stress,


115

the microstructure changes. One possibility is the formation of spheres with a lower meancurvature, the other possibility is to orient the cylinders perpendicular to the substrate andmaintain the equilibrium cylinder diameter. Van Dijk and Van den Berg studied PS-b-PB-b-PS triblock films of different thickness containing 24 vol% PS.25 They stated (based on AFMobservations) that for films that are thinner than two repeat distances, the cylinders orientperpendicular to the surface whereas parallel orientation only existed at certain thicknessesthat were compatible with the repeat distance.

Figure 7.3 Tapping mode AFM phase images of thin (30 nm) PI-b-PFS copolymer films. A. IF 29/15 consistingof a worm-like structure B. IF 31/13 contains areas with hexagonally packed domains as well, due to the relieveof stress induced by the incommensurate film thickness C. AFM phase image of IF 36/12 displayinghexagonally arranged domains. The Fourier transforms, with center and side corresponding to a reciprocalspacing d-1 equals 0 and 62.5 µm-1 respectively, are displayed in the inset of each image.

Chapter 7

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IF 36/12 (Figure 7.3 C) displays a morphology, completely consisting of hexagonallypacked domains. From the Fourier transform, shown in the inset, the domain spacing wasdetermined to be 30 nm (2/√3 times the 26 nm spacing for the Fourier transform), which isapproximately equal to the film thickness. A weaker ring can also be observed in the Fouriertransform of Figure 7.3 C at 15 nm, which confirms the hexagonal packing (hexagonalarrangement displays intensities at relative distances of 1 and √3 in q-space). The existence ofthis morphology strongly depends on the film thickness such as it was observed for IF 31/13.The value of the film thickness must match the interdomain distance in order to form thehexagonal structure. A slightly thicker film results in worm-like features, mixed with thehexagonal structure.

Theoretical studies predict the transition from parallel to perpendicular cylinders forvery thin films under confinement.8 If differences between the surface tensions of the twoblocks is small and the thickness is incommensurate with the domain spacing, a morphologywith vertically oriented cylinders can be obtained according to this theory. The transition isexplained as a result of the stress generated by the mismatch between the film thickness andthe domain distance.

Figure 7.4 Tapping mode AFM height image (25 µm scan size) of a thin film (< 30 nm) of IF 29/15. The filmthickness is incommensurate with the domain spacing, which results in the formation of holes.

A further decrease of the film thickness to values below the repeat distance results inthe formation of holes in the copolymer film. In block copolymer films on a solid substrate,the stress induced by the shortage of material is relieved by the formation of holes in such a


117

way that the remaining film has a thickness that is commensurate with the domaindimensions. Figure 7.4 displays an AFM height image of a large area of an IF 29/15 film. In-between the holes, the worm-like morphology, is still present (Figure 7.5 A) for IF 29/15.Similar hole formation is found for films of IF 31/13 and IF 36/12 of thicknesses that areincommensurate with their domain spacing. The AFM image of an IF 36/12 film in Figure7.5 B shows a dot-like pattern near a hole, followed by a region of the worm-like structure,and hexagonal packed domains further away from a hole. This structure is in agreement withthe previous observation for IF 36/12 films, and can be rationalized by considering thegradual increase of the film thickness near a hole. The thickness was found to be larger thanthe average thickness of the remaining film in areas showing the worm-like structure. Furtheraway from the hole, the hexagonal structure is observed. The hexagonal structure is thereforestable enough to prevent the hole from further growing which would otherwise result in astructure as observed in Figure 5A. Dewetting resulting in hole formation for film thicknessesless than 30 nm and occurs quickly after spin-casting, usually within hours.

Figure 7.5 A. Tapping mode AFM phase image (2 µm scan size) at higher magnification of a thin (< 30 nm) IF29/15 film. The holes are formed so the film can retain the thickness that is commensurate with a single domainlayer. B. Tapping mode AFM phase image (2 µm scan size) at higher magnification of a thin (< 30 nm) IF 36/12film. An in-plane worm-like morphology can be observed in areas where the film thickness is higher than theaverage thickness of the remaining film. Further away from the hole, the hexagonal structure is still present.

It is interesting to note that the formation of islands on top of the film, due to materialexcess, has not been observed. Such absence of island formation was previously observed forasymmetric diblock copolymers.10,16,26 The formation of islands and holes in thicker films canbe suppressed by lattice distortions in the film. Changes in surface topography generallyoccur over long time scales for asymmetric copolymers. In films that are thinner than thedomain spacing, lattice distortions are not possible and the film forms holes in order torelieve stress.

Chapter 7

118

Figure 7.6 displays an AFM image of a 29 nm thin film of IF 36/12, 50 hours afterspin-casting. The two holes in this film, which are visible in Figure 7.6, were formed withinhours after spin-casting due to incompatibility between the film thickness and the domainspacing as was discussed before. On the other hand, the large “hedrite-like” feature in the topof this image is slowly formed during storage at room temperature. Its shape suggests that thestructure evolves via a nucleation and growth mechanism. Similar features can be observedfor all three diblock copolymer films for all thicknesses studied, several days after spin-casting. In the smooth film region of the image, the microstructure is still preserved.

Figure 7.6 Tapping mode AFM height image of a thin IF 36/12 film after storage at room temperature for twodays. Similar structures were observed for all three diblock copolymers and all thicknesses studied after severaldays of storage at room temperature.

Figure 7.7 A and B show the edge of such a growing hedritic structure in an IF 31/13film at a higher magnification. The scan shown in Figure 7.7 B was taken from the same area,approximately 20 minutes after capturing the first scan. It is clearly visible that the cylindersbreak up in front of the slowly growing structure. The square inside the image contains thesame reference area, visible in both images. Eventually the entire film consists of such“hedrite-like” structure, as shown in Figure 7.8. However, the film does not dewet to anextent as observed for annealed films (Figure 7.1), even after storage for several months.


119

Figure 7.7 Subsequent tapping mode AFM images of a thin IF 31/13 film at the edge of a structure, similar tothat displayed in Figure 7.6. The time gap between the two images is approximately 20 minutes. The squareinside the images contains a reference area.

Figure 7.8 AFM height image of a thin PI-b-PFS film after storage for several months at room temperature. Thecomplete film surface is transformed into a hedrite like structure.

DSC experiments on bulk samples stored for long periods of time at roomtemperature revealed crystallization of the PFS block. Figure 7.9 displays the DSC heatingtraces of IF 29/15 after one week storage at room temperature, as well as the second heatingtrace. All three diblock copolymers show similar results. The endotherm at 138 °C indicatesthe melting of the crystallized PFS phase. In the second heating trace this melting is absent,

Chapter 7

120

which indicates that the crystallization did not take place during the heating scan. Themelting endotherm has a magnitude of approximately 7.4 J/g for IF 29/15. Since the blockcopolymer consists of 35 wt% of the crystallizing PFS block, the melting enthalpy for thePFS phase is approximately 21 J/g. This is a relatively large value compared to the meltingenthalpies for PFS homopolymers usually observed (<15 J/g), indicating high levels ofcrystallinity.21 Also the melting temperature of 138 °C indicates that the crystallization is welladvanced, since the temperature is close to the equilibrium melting temperature obtained forthe PFS homopolymer (143 °C). Finally, it is worth noting that the size of the hedriticstructure as observed in Figure 7.8 is much larger compared to the typical size of spherulitesof crystallized homopolymer (typically 20 µm).21

50 60 70 80 90 100 110 120 130 140 150

1 week at RT

2nd

heating

End

o >

Temperature [oC]

Figure 7.9 DSC heating traces of IF 29/15, after storage for 1 week at room temperature, showing the meltingof the PFS phase. The second heating trace does not show the melting transition, which indicates thatcrystallization takes place upon storage at room temperature and is not a result of crystallization during theheating scan.

Crystallization of asymmetric diblock copolymers from phase separated melts hasbeen described in the literature. The strength of the segregation was found to have a profoundeffect on the structure after crystallization. The reports regarding weakly segregated blockcopolymers in which the crystallizing component is present in the minority phase, areespecially relevant for a comparison.27,28 Ethylene block copolymers, that form a hexagonallypacked cylindrical structure in an amorphous matrix have been studied in the weaksegregation limit. After crystallization of the ethylene block, the materials form alternatinglamellae. The spacing between these lamellae in the corresponding morphology depends onthe thermal history. The original microphase separated morphology is often found to bedestroyed due to the chain folding of the crystallizing block. Avrami analysis on the


121

crystallization kinetics of ε-caprolactone-butadiene block copolymers have indicated that theprimary stage of crystallization was hardly influenced by the phase separation.29 Spheruliticmorphologies were observed after crystallization of block copolymers, with crystallizingblock volume fractions as low as 12 %.30 Our observations of diblock copolymer thin films,consisting of amorphous isoprene and crystalline ferrocenyldimethylsilane, are consistentwith these previous studies.

Figure 7.10 displays a TEM image of a freestanding IF 36/12 diblock copolymer film.An amorphous region, exhibiting the phase separated microdomain structure, can be observednear the left bottom quarter of this image. On the right-hand side extended crystals are clearlyvisible that appear dark in the TEM image and have inter-crystal spacings that arecomparable to the microdomain spacing. The breakup of the block copolymer microdomainmorphology near the boundary of the hedritic morphology is observable. This TEM imageconfirms our conclusions made on the basis of AFM experiments.

Figure 7.10 Transmission electron micrograph of a freestanding film of IF 36/12. The edge of the hedriticfeature is visible. In the amorphous region the microdomain structure can still be distinguished.

7.3 Conclusions

Thin films of asymmetric PI-b-PFS copolymers display a complex phase andmorphology behavior when cast on a silicon substrate. Because of the unfavorable interactionbetween the block copolymer film and the substrate, the copolymer completely dewets thesubstrate after annealing at 150 °C, which is above the order-disorder temperature of thecopolymers studied. At room temperature the films are stable for a longer time (days),allowing the study of surface pattern formation using atomic force microscopy. Although allthree diblock copolymers studied exhibit a cylindrical structure, the influence of the filmthickness on the morphology was distinct for each copolymer composition. Block copolymers

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122

that contained relatively less ferrocenylsilane (IF 31/13 and IF 36/12, containing 24 and 20vol% respectively) were able to form hexagonally arranged domains for “frustrating”thicknesses, whereas IF 29/15 (28 vol% ferrocenylsilane) displayed a worm-like structure andhole-formation. There is however a limit to the amount of stress that can be released by thechange in morphology. For all films that are thinner than a certain thickness (typically <30nm) hole formation was observed. Upon storage over longer periods, the PFS phase of theblock copolymer crystallizes at room temperature. During crystallization large hedrite likestructures slowly evolve and destroy the microdomain morphology. This is in accordancewith previous observations on weakly segregated asymmetric block copolymers that have acrystallizing block in the minority phase. Interestingly, the melting enthalpy, the meltingtemperature, and the hedritic feature indicate that the crystals must be well-ordered andrelatively perfect compared to the crystals of PFS homopolymer.

7.4 Experimental

Isoprene/toluene (10/90 vol%) was dried over dibutylmagnesium and distilled undervacuum. 1,1’-dimethylsilylferrocenophane was prepared as described earlier,31 and waspurified by subsequent vacuum sublimation and recrystallization in heptane at -60 °C. n-Butyllithium was diluted with heptane (dried from BuLi and distilled under vacuum) toapproximately 0.2 M. THF was dried over butyllithium and distilled under vacuum.

Poly(isoprene-block-ferrocenylsilanes) were prepared by subsequent anionicpolymerization of isoprene and ferrocenylsilane. The polymerization of isoprene (in toluene)was initiated by n-butyllithium and allowed to proceed for 16 hours. After the polymerizationof isoprene was complete, 1,1’-dimethylsilylferrocenophane was added and the solution wasstirred for 10 minutes. THF was added since the polymerization of 1,1’-dimethylsilylferrocenophane does not take place in toluene. After two more hours, thereaction was terminated by the addition of a few drops of degassed methanol. The polymerswere precipitated in methanol, dried under vacuum and subsequently analyzed by GPC and1H NMR.

Thin film samples were prepared by spin-casting a polymer solution in toluene on 3-inch silicon wafers, covered by the corresponding native oxide layers. Prior to spin-casting,the wafers were cleaned in 100% nitric acid at room temperature for 10 minutes and 70%nitric acid at 95 °C for 15 minutes. After cleaning, the silicon substrates were thoroughlyrinsed with water and dry-spinned. Film thicknesses were varied by changing the polymerconcentration and spinning speed and were measured by ellipsometry.

The development of the morphology was studied by atomic force microscopy (AFM)operating in the tapping mode. AFM experiments were performed using a NanoScope III Ainstrument (Digital Instruments) utilizing NanoSensor tapping tips. The amplitude ofoscillation at free vibration, A0, was set to 5.0 V. The operating setpoint ratio (A/A0) was setto relatively low values (A/A0 ~ 0.6) to obtain best contrast. For visualization of the phase


123

separation in thin copolymer films (typically 1 or 2 µm scan sizes), the phase image gave bestcontrast.32 Height images were used for larger scan sizes (typically 5 to 50 µm) for capturingthe hole formation and dewetting behavior of the thin films.

Transmission electron microscopy (TEM) experiments were performed on a JEOL200 CX microscope, operating at 160 kV. Freestanding diblock copolymer films wereprepared by placing a drop of polymer solution in toluene on a distilled water bath. Afterevaporation of the solvent, the film was picked up on a copper TEM grid and placed invacuum for several hours. No staining was required, due to the larger electron scatteringpower of the organometallic phase compared to the organic phase.

Rheological experiments, to determine the bulk order-disorder transition temperature(TODT), were performed on a Paar Physica UDS 200 rheometer, using the plate-plateconfiguration. While applying a strain of 1% at a frequency of 1 Hz, the temperature wasincreased from 25 to 150 °C. The TODT was taken as the temperature at which the shearelastic modulus collapsed.

A Perkin-Elmer DSC-7 was used for thermal analysis. In DSC experiments a scan rateof 10 K/min was employed. Tg values quoted in this study correspond to the onset of the heatcapacity change as determined by the software of the equipment used.


1 M. Park, C. Harrison, P. M. Chaikin, R. A. Register, D. H. Adamson, Science 1997,276, 1401.

2 (a) R. G. H. Lammertink, M. A. Hempenius, J. E. van den Enk, V. Z.-H. Chan, E. L.Thomas, G. J. Vancso, Adv. Mater. 2000, 12, 98. (b) R. G. H. Lammertink, M. A.Hempenius, G. J. Vancso, V. Z.-H. Chan, E. L. Thomas, Abstract ACS, PMSEPreprint 1999, 81, 14.

3 L. Leibler, Macromolecules 1980, 13, 1602.4 S. H. Anastasiadis, T. P. Russell, S. K. Satija, C. F. Majkrzak, Phys. Rev. Lett. 1989,

62, 1852.5 J. P. Spatz, S. Sheiko, M. Möller, Adv. Mater. 1996, 8, 513.6 J. P. Spatz, M. Möller, M. Noeske, R. J. Behm, M. Pietralla, Macromolecules 1997,

30, 3874.7 M. S. Turner, M. Rubinstein, C. M. Marques, Macromolecules 1994, 27, 4986.8 K. Y. Suh, Y. S. Kim, H. H. Lee, J. Chem. Phys. 1998, 108, 1253.9 H. P. Huinink, J. C. M. Brokken-Zijp, M. A. van Dijk, G. J. A. Sevink, J. Chem. Phys.

2000, 112, 2452.10 Y. Liu, W. Zhao, X. Zheng, A. King, A. Singh, M. H. Rafailovich, J. Sokolov, K. H.

Dai, E. J. Kramer, S. A. Schwarz, O. Gebizlioglu, S. K. Sinha, Macromolecules 1994,27, 4000.

11 H. Yokoyama, E. J. Kramer, M. H. Rafailovich, J. Sokolov, S. A. Schwarz,Macromolecules 1998, 31, 8826.

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12 C. S. Henkee, E. L. Thomas, L. J. Fetters, J. Mater. Sci. 1988, 23, 1685.13 L. H. Radzilowski, B. L. Carvalho, E. L. Thomas, J. Polym. Sci. Part B, Polym. Phys.

1996, 34, 3081.14 (a) C. Harrison, M. Park, P. M. Chaikin, R. A. Register, D. H. Adamson, N. Yao,

Polymer 1998, 30, 2733. (b) C. Harrison, M. Park, P. M. Chaikin, R. A. Register, D.H. Adamson, N. Yao, Macromolecules 1998, 31, 2185.

15 J. Israelachvili, “Intermolecular and Surface Forces”, 2nd ed., Academic Press,London, 1991, 315.

16 A. Karim, N. Singh, M. Sikka, F. S. Bates, W. D. Dozier, G. Felcher, J. Chem. Phys.1994, 100, 1620.

17 M. Morton, L. J. Fetters, Rubber. Chem. Technol. 1975, 48, 359.18 G. Reiter, Phys. Rev. Lett. 1992, 68, 75.19 A. Sharma, G. Reiter, J. Coll. Interf. Sci. 1996, 178, 383.20 R. Limary, P. F. Green, Langmuir 1999, 15, 5617.21 (a) R. G. H. Lammertink, M. A. Hempenius, I. Manners, G. J. Vancso,

Macromolecules 1998, 31, 795. (b) R. G. H. Lammertink, M. A. Hempenius, G. J.Vancso, Macromol. Chem. Phys. 1998, 199, 2141.

22 Care should be taken in identifying the phases by their relative contrast in height andphase mode images, since the contrast is not unique to a particular phase but issensitive to changes in the operating conditions of the microscope: J. P. Pickering, G.J. Vancso, Polym. Bull. 1998, 40, 549.

23 The contact angle of a droplet of water with a block copolymer film wasapproximately 102°, which is near the value for pure poly(isoprene). Forpoly(ferrocenyldimethylsilane) this contact angle is about 96°. Furhtermore, XPS datain ref. 2, show almost no Fe and Si signal from the poly(ferrocenyldimethylsilane)block in a thin block copolymer film, but only a large C signal. This confirms that thefree surface is wetted by the organic poly(isoprene) phase.

24 R. G. H. Lammertink, M. A. Hempenius, E. L. Thomas, G. J. Vancso, J. Polym. Sci.,Polym. Phys. 1999, 37, 1009.

25 M. A. Van Dijk, R. van den Berg, Macromolecules 1995, 28, 6773.26 N. Singh, A. Kudrle, M. Sikka, F. S. Bates, J. Phys. II (France) 1995, 5, 377.27 D. J. Quiram, R. A. Register, G. R. Marchand, A. J. Ryan, Macromolecules 1997, 30,

8338.28 A. J. Ryan, I. W. Hamley, W. Bras, F. S. Bates, Macromolecules 1995, 28, 3860.29 S. Nojima, H. Nakano, Y. Takahashi, T. Ashida, Polymer 1994, 35, 3479.30 P. Rangarajan, R. A. Register, L. J. Fetters, Macromolecules 1993, 26, 4640.31 A. B. Fischer, J. B. Kinney, R. H. Staley, M. S. Wrighton, J. Am. Chem. Soc. 1979,

101, 6501.32 Ph. Leclère, R. Lazzaroni, J. L. Brédas, J. M. Yu, Ph. Dubois, R. Jérôme, Langmuir

1996, 12, 4317.

Nanostructured Surfaces

125

&KDSWHU��

Nanostructured Surfaces by a Combination of Self-Assembly of Block Copolymers and Reactive Ion

Etching #

Thin organic-organometallic diblock copolymer films have been employed fornanolithographic applications. Combining self-assembly of the blockcopolymer in thin films with the specific properties of the organometallicphase, regularly arranged nanostructures were obtained. The diblockcopolymer films consisted of a styrene or isoprene organic block and aferrocenyldimethylsilane organometallic block. Due to the high resistance ofthe organometallic phase towards oxygen reactive ion etching (RIE) theorganic phases could be selectively removed. Upon etching, theorganometallic domains were converted into an inorganic oxide that proved tobe very stable towards fluorocarbon plasma etching. This allowed to transferthe pattern, generated by phase separation of the block copolymer, intounderlying silicon and silicon nitride substrates.

8.1 Introduction

Various approaches have been employed to enhance the etch selectivity in thin blockcopolymer films. A common feature of these approaches is to selectively load one of thephases with suitable inorganic components. If one block contains double bonds, transitionmetal containing compounds, which possess reactivity for unsaturated double bonds, can beused to load this block and thus will selectively aggregate in the corresponding phase. Thistechnique is widely used for contrast enhancement in transmission electron microscopy.1 Parket al.2 used an OsO4-stained microphase-separated thin film of poly(styrene-block-butadiene),PS-b-PB, which produced holes upon RIE in silicon nitride substrates, resulting in an etch

# Parts of the work described in this Chapter have been published: (a) R. G. H. Lammertink, M. A. Hempenius,G. J. Vancso, V. Z.-H. Chan, E. L. Thomas, ACS Fall Meeting, New Orleans, PMSE Preprint 1999, 81, 14. (b)R. G. H. Lammertink, M. A. Hempenius, J. E. Van den Enk, V. Z.-H. Chan, E. L. Thomas, G. J. Vancso, Adv.Mater. 2000, 12, 98.

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selectivity of only 2:1. Möller et al. in a series of papers discussed the use of poly(styrene-block-2-vinylpyridine), PS-b-P2VP, to prepare masks for nanolithography either by loadingthe P2VP domains with gold particles,3 or by selective growth of Ti on top of PS domains.4

Aksay et al.5 used PS-b-PB-b-PS triblock copolymers and selectively loaded a bariumtitanate precursor into hydroxylated thin films of this polymer into the trans-1,2 polybutadiolmatrix. These approaches achieve etching contrast by selective introduction of inorganiccomponents in one of the phases. The advantage of our approach is that the inorganiccomponents are inherent in our block copolymer system thereby eliminating a loading step bya metal and allowing for simple one-step lithography.

One-step approaches have already been considered where “imageable” blockcopolymers are used, of which one of the blocks intrinsically contains elements that providean etch barrier.6,7 Silicon containing polymers seem to be good candidates, especially whenetch resistance in an oxygen plasma is required. Such polymers form a thin SiOx layer on thepolymer surface when exposed to an oxygen plasma.8 After the formation of such a layer, theetch rate is determined by the competition between the oxide growth and the oxide removalby ion sputtering. This results in relatively low etch rates and therefore in a high etchselectivity, as high as 50:1, compared to organic compounds.

It is known that in thin films of block copolymers with laterally separatedmicrodomains, surface free energy differences between the blocks and the affinity of one ofthe blocks for the substrate can lead to in-plane orientation of cylinders or lamellae.9,10

Microdomain orientations in films make these unsuitable as templates for etching arrays ofregularly sized and spaced dots. If regularly packed cylinders are oriented upright withrespect to the substrate in the film, their top view in the film would exhibit circular regions.Such perpendicular orientation of cylinders was observed for poly(styrene-block-butadiene)when cast on a water surface,11 after quick solvent evaporation,12 and for very thin spin-coated films of poly(styrene-block-butadiene-block-styrene).13 However, the former approachyielded millimeter size films with variable thickness and is therefore less useful in potentialnanolithographic applications.

In this chapter the use of organometallic microdomain formation in block copolymerfilms for lithographic applications is considered. Self-assembled thin block copolymer filmswere used to serve as lithographic templates that can be used to generate laterallynanostructured surfaces.


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8.2.1 PATTERNS BY SELF-ASSEMBLY OF DIBLOCK COPOLYMER THIN FILMS

A thin film (approximately 30 nm) of isoprene-block-ferrocenylsilane copolymer (IF36/12, 36 kg/mol PI -block- 12 kg/mol PFS) forms lateral patterns when casted on siliconsubstrates.14 For the approach reported here, it is of importance that at given volume fractionsand film thicknesses it is possible to obtain a regular morphology, which consists ofhexagonally packed organometallic domains (see Figure 8.1). After spin-casting the entiresilicon wafer is covered by this regular morphology, without the need for subsequentannealing.

Figure 8.1 Tapping mode AFM phase image of 30 nm thin spin cast film of IF 36/12. The film consists oflaterally separated poly(ferrocenylsilane) domains in an organic matrix.

Table 8.1 shows the atomic concentrations of carbon (1s), oxygen (1s), silicon (2p),and iron (2p) for a PI-b-PFS diblock and PFS homopolymer, before and after oxygen reactiveion etching (O2-RIE) as determined by XPS experiments. For the unetched PFShomopolymer we find atomic concentration ratios for C:Si:Fe of 12:0.86:0.97, which is closeto the theoretical relative amounts of 12:1:1 (C12H14SiFe for a repeat unit). After O2-RIE theoxygen concentration increased significantly, which was accompanied by a decrease in thecarbon concentration. A large amount of carbon is removed by RIE compared to the relative

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amounts of silicon and iron. Upon etching the PFS homopolymer, the Si/C ratio increasedfrom 0.07 to 0.24. At the same time the Fe/C ratio increased from 0.08 to 0.51. It isinteresting to see that following O2-RIE the relative amount of Fe with respect to Siincreased. This suggests that Fe is more stable towards RIE treatment than Si by usingoxygen plasma. From the corresponding XPS spectra shown in Figure 8.2 and Figure 8.3, itcan be seen that both the Si2p and Fe2p binding energies increase following the oxygenplasma treatment, indicating the conversion of Si and Fe into a complex oxide.15 Theseresults demonstrate that poly(ferrocenylsilane) forms an etch barrier when exposed to anoxygen plasma due to the formation of an oxide layer at the surface of the polymer.16 Thisleads to a very low etch rate of the organometallic part compared to organic polymers, likepoly(isoprene), since the oxide protects the underlying polymer from the plasma. Under theconditions employed here, we observed an etch selectivity for PI:PFS of approximately 40:1,as determined from the corresponding homopolymer etch rates (ca. 7.3 nm/s and 0.18 nm/sfor PI and PFS respectively). This very high etch selectivity allows one to remove the organicphase, leaving domains behind that are converted into an iron-silicon-oxide.

Table 8.1 Atomic concentrations of PI-b-PFS diblock and PFS homopolymer films before and after O2-RIE.

PFS homopolymer PFS homopolymerO2 etched

PI-b-PFS diblock PI-b-PFS diblockO2 etcheda

Element atom % atom % atom % atom %C1s 85.1 33.3 98.7 22.1O1s 1.9 41.8 0.4 49.6Si2p 6.1 8.0 0.7 14.4Fe2p 6.9 16.9 0.2 13.8a Due to the open structure of the etched PI-b-PFS diblock film, additional silicon from the substrate is observed

and thus the relative silicon concentration in the film is overestimated.

The XPS data (Table 8.1) of a PI-b-PFS diblock film reveals that a large amount ofcarbon is present in the original film, as expected. As mentioned, after exposure to the O2-plasma, the carbon concentration decreased and oxygen, silicon, and iron concentrationsincreased. This is in agreement with the removal of the organic phase, and with the oxideformation of the silicon and iron, which are present in the organometallic domains. Due to theopen, porous structure, which is formed when the organic phase is removed from the diblockfilm, the plasma also reached the underlying silicon substrate in some areas. A film thicknessprior to etching of approximately 60 nm and an etch time of 15 s (during which it is possibleto etch away approximately 100 nm PI) confirms the observation that the substrate wasreached. Etching all the way through the film explains the relatively large amount of silicondetected by XPS for the etched diblock film compared with the etched PFS homopolymer.The additional peak at 99.8 eV, caused by silicon from the wafer, can be observed in the Si2pXPS spectrum of the O2 etched PI-b-PFS diblock. This peak is not observed for very thick(>200 nm) PI-b-PFS diblock films after O2-RIE.


129

110 108 106 104 102 100 98 96 94 92 90

1x104

2x104

3x104

4x104

5x104

Si2p

IF diblock IF O

2 etched

PFS homopolymer PFS O

2 etched

Cou

nts

Binding Energy (eV)

Figure 8.2 XPS spectra of Si2p for the IF diblock copolymer and the PFS homopolymer, before and afteroxygen reactive ion etching.

735 730 725 720 715 710 705 7005.0x104

1.0x105

1.5x105

2.0x105

2.5x105

3.0x105

3.5x105

Binding Energy (eV)

Fe2p IF diblock IF O

2 etched

PFS homopolymer PFS O2 etched

Cou

nts

Figure 8.3 XPS spectra of Fe2p for the IF diblock copolymer and the PFS homopolymer, before and afteroxygen reactive ion etching.

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Figure 8.4 shows a representative AFM height image of a thin film (~30 nm) of IF36/12 (see Figure 8.1) after O2-RIE for 10 s. The organometallic domains that were convertedinto an oxide by O2-RIE, are still arranged in their original hexagonal structure. The heightcontrast in the AFM image increased significantly upon etching, due to the removal of theorganic matrix. The line scan in Figure 8.4 (top right) indicates a domain height ofapproximately 7 nm. According to the Fourier transform of the AFM image (bottom right inFigure 8.4), no change in domain spacing occurred upon etching. This simple one-stepetching process results in a nanostructured surface with an area density of approximately 1.2 ·1011 dots/cm2. After etching the nanostructured surface is stable when stored at roomtemperature or rinsed with common organic solvents, in contrast to the thin diblockcopolymer film.

Figure 8.4 Tapping mode AFM height image (left) of a 30 nm thin film of IF 36/12 after O2-RIE (as shown inFigure 8.1 before etching). A single line scan (top right) and the 2D Fourier transform (bottom right)demonstrate the regularity of the structure. The height difference between the two markers shown was 7.0 nm.

The aspect ratio of the domains, however, is limited. Due to the removal of materialfrom the organometallic phase, the features shrink, which leads to somewhat flatter domains.To obtain a good visualization of the domains, a cross-sectional transmission electronmicrograph is shown in Figure 8.5. Since the cross-section contains more than one row ofdomains, the sample had to be tilted so that the corresponding rows were aligned. The dotsare regularly spaced, which confirms the AFM observations. Furthermore, selected areadiffraction indicated that the inorganic nanodomains were essentially amorphous.


131

Figure 8.5 Cross sectional TEM of a thin film of IF 36/12 after O2-RIE treatment. The Fe-Si-oxide domainsappear dark in the TEM, as does the silicon oxide layer at the surface of the silicon substrate.

Figure 8.6 AFM height images of a thin IF 36/12 film on a silicon nitride substrate. The left image was takenafter O2-RIE of the film, similar to the film displayed in Figure 8.1. The film displayed on the right, wassubsequently treated for 30 seconds with a CF4/O2 plasma. The line scans below each image clearly demonstratethe increase in aspect ratio of the domains.

For potential applications, a higher aspect ratio of the nanodomains is highlydesirable.17 In order to obtain such an increase in domain height, a CF4/O2 reactive iontreatment was applied to the substrate. Such a plasma is frequently used in semiconductorindustry for etching silicon, silicon oxide, and silicon nitride. Due to the presence of iron inthe nanodomains, the domains are more stable towards a fluorocarbon plasma compared tothe substrate. More precisely, the etch rate of silicon and silicon nitride under the conditions

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of the experiment was found to be approximately 20 nm/min, whereas no significant changein film thickness was observed for poly(ferrocenylsilane) after 1 minute of etching.

Due to the presence of iron in the inorganic domains, the pattern can be transferredinto the underlying substrate by the action of CF4/O2 reactive ion etching. Figure 8.6 showsthe AFM images of a thin IF 36/12 film after O2-RIE treatment (left) followed by a CF4/O2-RIE treatment (right) for 30 seconds. Longer etching times resulted in destruction of theregular pattern. As can be concluded from the image scale and the line scans, the height ofthe features has increased upon the etching treatment. This is a result of the removal of thesubstrate in between the domains combined with the high etch resistance of the inorganicdomains.

8.2.2 PATTERNS BY SELF-ASSEMBLY ON DIFFERENT LENGTH SCALES

Self-assembly strategies can be used to obtain regular patterns on a variety of lengthscales. Larger dimensions may be extremely interesting if one wants to obtain structures forphotonic applications, e.g. photonic band-gap materials.18 The use of block copolymers withlarger molar masses will result in larger domain spacings, however, the extent to which thedimensions can be increased is quite limited. It is known that the domain spacings do notscale linearly with molar mass, but rather as L ~ M2/3, where L is the lamellar spacing and Mis the molar mass.19 There is, however, also a limit to the maximum attainable molar mass ofdiblock copolymers. Furthermore, if the copolymer chains become very long, it will be moredifficult to obtain very regular equilibrium structures, as the diffusion dynamics are sloweddown considerably.

A styrene-block-ferrocenyldimethylsilane copolymer (SF 101/24) with a relativelylarge molar mass was used to obtain a regular pattern with larger domain spacings. Aftercasting and annealing the organic phase was selectively removed by means of an O2-RIEtreatment. Figure 8.7 displays AFM height images of diblock copolymer thin films afterselective removal of the organic phase.

An alternative technique uses block copolymers of which one block has a highpreference for the substrate. A pattern is formed by the strong adsorption of one block whilethe other block dewets this adsorbed layer. This has been investigated extensively for styrene-block-vinylpyridine copolymers on hydrophilic substrates.20,21 Lateral structures fromhydrophilic-hydrophobic block copolymers can also be obtained at the air-water interface.Aggregates, so-called surface micelles, have been observed for a number of hydrophilic-hydrophobic diblock copolymers and are therefore believed to be general phenomena.22 Thinfilms that float on a water bath can be picked up on solid substrates by standard Langmuir-Blodgett (LB) techniques so they can be studied by atomic force microscopy.


133

Figure 8.7 AFM height images of diblock copolymer thin films. A. Low molar mass isoprene-block-ferrocenylsilane copolymer IF 36/12. B. High molar mass styrene-block-ferrocenylsilane copolymer SF 101/24.As can be seen from the images, a higher molar mass results in larger domain dimensions, but also in a lessregular morphology.

Trough area (cm 2)

0 100 200 300 400 500 600 700

Sur

face

pre

ssur

e (π

) (dy

ne/c

m)

0

5

10

15

20

Figure 8.8 Surface pressure – trough area isotherm of a PVP-b-PS / PS-b-PFS blend on water.

We employed this technique to obtain lateral structures with relatively largedimensions. For such an approach, a vinylpyridine-block-ferrocenylsilane copolymer wouldbe suitable. We used a PVP-b-PS (105/104) and a PS-b-PFS diblock (SF 54/52) blend (mixedin a ratio of 1:4 respectively). By using a blend of two diblock copolymers, theorganometallic phase could be “anchored” onto the hydrophobic PS block. In Figure 8.8 thepressure-area isotherm of the blend on water is displayed. Upon reducing the area the surface

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aggregates become increasingly compressed. This leads to an increase in the surface pressure,which becomes evident in Figure 8.8 at an area of approximately 400 cm2. Films wereprepared in a Langmuir-Blodgett trough by casting the blend from a chloroform solution ontothe water bath. During the transfer of the film onto the solid substrate, a constant surfacepressure was maintained.

Figure 8.9 displays the AFM images of an LB film that was picked up at a constantsurface pressure of 10 dyne/cm. The affinity of the PVP block for the water surface results inspreading of this block, whereas the hydrophobic block forms aggregates. Rodlike structuresare formed which have typical domain distances of approximately 200 nm. Within theaggregates, the PS-b-PFS copolymer phase separates, which results in the porous structurewithin the rods after the O2-RIE treatment. By varying the surface pressure, which is appliedduring film transfer, different morphologies, like circular aggregates and pancakes, can beobtained.

Figure 8.9 AFM images of a LB film prepared at a constant surface pressure of 10 dyne/cm consisting of ablend of PVP-b-PS and PS-b-PFS block copolymers. A. After deposition from water onto a silicon substrate. B.Followed by an O2-RIE treatment, which removed the organic phase.

8.3 Conclusions and Outlook

The combination of self-assembly of thin films of organic-organometallic blockcopolymers with reactive ion etching has been used to obtain regular nanostructured surfaces.The presence of iron and silicon in the ferrocenylsilane block allows for the selective removalof the other organic block by means of oxygen-reactive ion etching. We demonstrated that anisoprene-block-ferrocenyldimethylsilane diblock copolymer on a silicon substrate resulted inthe generation of a nanoperiodic surface with approximately 1.2 · 1011 dots/cm2. Asubsequent reactive ion etch treatment using CF4/O2 feed gas allows the transfer of thenanoperiodic structure into the underlying silicon or silicon nitride substrate. For manyapplications the patterned surfaces have to be defect-free over large areas and the orientation


135

of the domains should be controlled. Control over domain orientation has been achieved byusing for example electric fields23 and directional solidification.24

Surface aggregates can be obtained at the air-water interface using hydrophilic-hydrophobic block copolymers. It was demonstrated by blending PVP-b-PS and a PS-b-PFScopolymers, lateral structures were obtained at the air-water interface which could betransferred via standard Langmuir-Blodgett techniques onto various substrates. Rodlikefeatures were obtained that display an internal structure after oxygen reactive ion etching as aresult of the phase separation within the rods.

8.4 Experimental

Poly(isoprene-block-ferrocenylsilane) and poly(styrene-block-ferrocenylsilane) wereprepared by subsequent anionic polymerization of isoprene or styrene and ferrocenylsilane.The polymerization of isoprene (in toluene) was initiated by n-butyllithium and allowed toproceed for 16 hours. Polymerization of styrene (in ethylbenzene) was initiated with n-butyllithium and allowed to proceed for 6 hours. After the polymerization of the organicblock was complete, 1,1’-dimethylsilylferrocenophane was added and the solution was stirredfor 10 minutes. THF was added since the polymerization of 1,1’-dimethylsilylferrocenophanedoes not take place in toluene. After two more hours, the reaction was terminated by theaddition of a few drops of degassed methanol. The polymers were precipitated in methanol,dried under vacuum and subsequently analyzed by GPC and 1H NMR. The polymers aredenoted IF (Isoprene-Ferrocenylsilane) and SF (Styrene-Ferrocenylsilane), followed by theirrespective molar mass in kg/mol.

Thin films (approximately 30 nm) were prepared by spin-coating of a polymersolution in toluene (~1 % wt/vol) onto a silicon wafer, containing its native oxide, at 4000rpm for 30 s. The wafers were cleaned prior to spin-coating. The cleaning involved treatmentfor 10 minutes in 100% nitric acid, rinsing subsequently with deionized water, followed byadditional treatment for 15 minutes in 70% nitric acid at 95 °C and finally rinsing withdeionized water and dry spinning.

Langmuir-Blodgett (LB) films were prepared by spreading a blend of PVP-b-PS andPS-b-PFS (SF) onto a clean water surface. The polymer film was transferred onto a siliconsubstrate at constant surface pressure via standard LB techniques.

Film thicknesses were varied by spin-casting from different concentrations of polymersolutions and were determined by ellipsometry. The thickness was determined on 25 spots ona wafer, with a standard deviation less than 1%.

Oxygen reactive ion etching (O2-RIE) experiments were carried out in an ElektrotechPF 340 apparatus. The pressure inside the etching chamber was 10 mTorr, the substratetemperature was set at 10 °C, and an oxygen flow rate of 20 cm3/min was maintained. Thepower was set at 75 W. The 30 nm PI-b-PFS diblock thin film, used for the morphologicalstudy, was etched for 10 s. The 60 nm thin PI-b-PFS diblock and PFS homopolymer films

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(used for XPS study) were etched for 15 s. To determine the etch ratio of the two polymers,PFS and PI homopolymer films were etched under the conditions mentioned above for 60 sand 15 s respectively. The decrease in film thickness was determined by ellipsometry.

Both etched and unetched 60 nm thin films of PFS homopolymer and PI-b-PFSdiblock copolymer were studied by X-ray Photoelectron Spectroscopy (XPS). Measurementswere performed with a Kratos XSAM-800 spectrometer using a Mg Kα X-ray source (1253.6eV). The input power was set at 150 W (15 kV and 10 mA) and the hemispherical analyzerwas placed perpendicular to the sample surface. Survey scans (0-1100 eV) were recorded toqualitatively determine the elements present. The atomic concentrations of carbon, oxygen,silicon and iron were determined from the relative peak areas calculated by numericalintegration of the detailed scans (20-30 eV windows). For a 30 nm thin film of PI-b-PFS thesame peaks were observed after O2-RIE compared to a 60 nm film of PI-b-PFS. Only the Sisignal, caused by silicon from the wafer is relatively more intense for the 30 nm etched PI-b-PFS film.

The morphology of the spin-coated films (unannealed) and the etched films wasstudied by atomic force microscopy (AFM) operating in the tapping mode. AFM experimentswere performed using a NanoScope III A instrument (Digital Instruments) utilizingNanoSensors tapping tips. For unetched films, the amplitude of oscillation at free vibration,A0, was set at 5.0 V. The operating setpoint ratio (A/A0), which is inversely proportional tothe damping of the cantilever, was set to relatively low values (A/A0 ~ 0.6) to obtain bestcontrast. Etched films could be studied best at low amplitudes of oscillation at free vibration(A0 = 1.0 V) and relatively high operating setpoint ratios (A/A0 ~ 0.9). Height and phaseimages were recorded at a scanning rate of 1 Hz.

Cross sectional TEM was performed on a Philips CM30 electron microscope,operating at 300 kV.


1 D. W. Schwark, D. L. Vezie, J. R. Reffner, E. L. Thomas, B. K. Annis, J. Mater. Sci.Lett. 1992, 11, 352.

2 M. Park, C. Harrison, P. M. Chaikin, R. A. Register, D. H. Adamson, Science 1997,276, 1401.

3 J. P. Spatz, T. Herzog, S. Mössmer, P. Ziemann, M. Möller, Adv. Mater. 1999, 11,149.

4 J. P. Spatz, P. Eibeck, S. Mössmer, M. Möller, T, Herzog, P. Ziemann, Adv. Mater.1998, 10, 849.

5 T. Lee, N. Yao, I. A. Aksay, Langmuir 1997, 13, 3866.6 A. H. Gabor, E. A. Lehner, G. Mao, L. A. Schnegenburger, C. K. Ober, Chem. Mater.

1994, 6, 927.


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7 (a) A. Avgeropoulos, V. Z.-H. Chan, V. Y. Lee, D. Ngo, R. D. Miller, N.Hadjichristidis, E. L. Thomas, Chem. Mater. 1998, 10, 2109. (b) V. Z.-H. Chan, J.Hoffman, V. Y. Lee, H. Iatrou, A. Avgeropoulos, N. Hadjichristidis, R. D. Miller, E.L. Thomas, Science 1999, 286, 1716.

8 (a) G. N. Taylor, T. M. Wolf, Polym. Eng. Sci. 1980, 20, 1087. (b) G. N. Taylor, T.M. Wolf, J. M. Moran, J. Vac. Sci. Technol. 1981, 19, 872.

9 C. S. Henkee, E. L. Thomas, L. J. Fetters, J. Mater. Sci. 1988, 23, 1685.10 T. P. Russell, G. Coulon, V. R. Deline, D. C. Miller, Macromolecules 1989, 22, 4600.11 P. Mansky, P. Chaikin, E. L. Thomas, J. Mater. Sci. 1995, 30, 1987.12 (a) G. Kim, M. Libera, Macromolecules 1998, 31, 2569. (b) G. Kim, M. Libera,

Macromolecules 1998, 31, 2670.13 M. A. van Dijk, R. van den Berg, Macromolecules 1995, 28, 6773.14 Several morphologies can be obtained when varying the film thickness and/or the

block copolymer composition. See also Chapter 7, and: R. G. H. Lammertink, M. A.Hempenius, G. J. Vancso, Langmuir 2000, in press.

15 Detailed results regarding the thickness and composition of the oxide are reported inChapter 4.

16 Auger electron spectroscopy (AES) data obtained for oxygen etched PFShomopolymer are consistent with the XPS results. Depth profiling (argon ionsputtering) indicated an approximately 10 nm thin layer rich of iron, silicon and oxideat the surface of the oxygen etched polymer film. See also Chapter 4.

17 For instance, silicon nanopillars show interesting fluorescence properties and areinteresting to investigate quantum effects in semiconductors: L. T. Canham, Appl.Phys. Lett. 1990, 57, 1046.

18 E. Yablonovitch, J. Opt. Soc. Am. B. 1993, 10, 283.19 A. M. Mayes, T. P. Russell, V. R. Deline, S. K. Satija, C. F. Majkrzak,

Macromolecules 1994, 27, 7447.20 J. C. Meiners, A. Ritzi, M. H. Rafailovich, J. Sokolov, J. Mlynek, G. Krausch, Appl.

Phys. A 1995, 61, 519.21 (a) J. P. Spatz, S. Sheiko, M. Möller, Adv. Mater. 1996, 8, 513. (b) J. P. Spatz, M.

Möller, M. Noeske, R. J. Behm, M. Pietralla, Macromolecules 1997, 30, 3874. (c) E.Y. Kramarenko, I. I. Potemkin, A. R. Khokhlov, R. G. Winkler, P. Reineker,Macromolecules 1999, 32, 3495. (d) I. I. Potemkin, E. Yu. Kramarenko, A. R.Khokhlov, R. G. Winkler, P. Reineker, P. Eibeck, J. P. Spatz, M. Möller, Langmuir1999, 15, 7290. (e) J. P. Spatz, P. Eibeck, S. Mössmer, M. Möller, E. Yu.Kramarenko, P. G. Khalatur, I. I. Potemkin, A. R. Khokhlov, R. G. Winkler, P.Reineker, Macromolecules 2000, 33, 150. (f) J. P. Spatz, S. Mössmer, C. Hartmann,M. Möller, T. Herzog, M. Krieger, H.-G. Boyen, P. Ziemann, B. Kabius, Langmuir2000, 16, 407.

22 (a) J. Zhu, A. Eisenberg, R. B. Lennox, J. Am. Chem. Soc. 1991, 113, 5583. (b) J.Zhu, R. B. Lennox, A. Eisenberg, Langmuir 1991, 7, 1579. (c) J. Zhu, A. Eisenberg,R. B. Lennox, J. Phys. Chem. 1992, 96, 4727. (d) S. Li, S. Hanley, I. Khan, S. K.Varshney, A. Eisenberg, R. B. Lennox, Langmuir 1993, 9, 2243. (e) S. A. Rice, B.Lin, J. Chem. Phys. 1993, 99, 8308. (f) M. Meszaros, A. Eisenberg, R. B. Lennox,Faraday Discuss. 1994, 98, 283. (g) Z. Li, W. Zhao, J. Quinn, M. H. Rafailovich, J.Sokolov, R. B. Lennox, A. Eisenberg, X. Z. Wu, M. W. Kim, S. K. Sinha, M. Tolan,

Chapter 8

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Langmuir 1995, 11, 4785. (h) S. Li, C. J. Clarke, R. B. Lennox, A. Eisenberg,Colloids. Surf. A 1998, 133, 191. (i) J. K. Cox, K. Yu, A. Eisenberg, R. B. Lennox,Phys. Chem. Chem. Phys. 1999, 1, 4417. (j) J. K. Cox, K. Yu, B. Constantine, A.Eisenberg, R. B. Lennox, Langmuir 1999, 15, 7714.

23 T. L. Morkved, M. Lu. A. M. Urbas, E. E. Ehrichs, H. M. Jaeger, P. Mansky, T. P.Russell, Science 1996, 273, 931.

24 C. De Rosa, C. Park, E. L. Thomas, B. Lotz, Nature 2000, in press.

Summary

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The work described in this thesis concerns the synthesis, characterization, andproperties study of ferrocenyldimethylsilane homopolymers and block copolymers. Due tothe presence of iron and silicon in the polymer main chain, these macromolecules possesscharacteristics that are very distinctive from common organic polymers. The objective of thisthesis was to investigate and combine the unprecedented features of these organometallicpolymers with block copolymer self-assembly, as discussed in Chapter 1.

A brief introduction was presented in Chapter 2 in topics that were relevant for thisthesis. Attention was given to crystallization and melting behavior of semi-crystallinepolymers. Furthermore, block copolymer self-assembly was described with special interest inthin films of block copolymers, blends of block copolymers with correspondinghomopolymer, and conformational asymmetry in block copolymers. The foundation of thetechniques that were employed throughout the work was also presented.

The anionic ring-opening polymerization of strained dimethylsilylferrocenophaneallowed the synthesis of well-defined polymers with targeted molar masses. Chapter 3describes the anionic synthesis, characterization, and thermal behavior of a range offerrocenylsilane homopolymers. These semi-crystalline macromolecules displayed a melting-recrystallization process during heating from the crystallized state. The reorganizationprocess involved the melting of thin lamellae that rapidly crystallized into thicker ones, whichmelted subsequently at somewhat higher temperatures. The occurrence of this reorganizationprocess was investigated by differential scanning calorimetry experiments that displayedmultiple melting transitions during a heating scan from the crystallized state. Results fromsmall-angle x-ray scattering measurements complemented these conclusions.

In Chapter 4, the action of reactive ion etching on poly(ferrocenylsilane) wasdiscussed. Treatment with an oxygen plasma resulted in the formation of a thin oxide layer,rich in iron and silicon, at the surface of the polymer. Due to the formation of this non-volatile oxide layer, the etch rate of this organometallic polymer was much lower comparedto the etch rate of organic polymers. In addition, the non-volatile oxide layer was also veryresistant towards fluorocarbon plasma treatment. It was demonstrated that substrates, likesilicon and silicon nitride, were selectively removed by the action of a fluorocarbon plasmawhile the iron-silicon-oxide remained.

Chapter 5 described the synthesis and self-assembly of styrene-block-ferrocenylsilanecopolymers. Block copolymers were synthesized by the sequential anionic polymerization ofstyrene and silylferrocenophane that resulted in near monodisperse materials with tailoredcompositions and molar masses. Upon phase separation, these block copolymers as well as

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their blends with homopolymer formed morphologies such as spheres on a bcc lattice,hexagonally packed cylinders, double gyroid, and alternating and perforated lamellae.Rheological experiments were performed to give insight into the order-disorder transitions.

Thin films, on solid substrates or freestanding, of a block copolymer of styrene andferrocenylsilane were studied which was discussed in Chapter 6. This diblock formed acylindrical morphology in the bulk, consisting of hexagonally packed ferrocenylsilanecylinders in an organic polystyrene matrix. In thin films on solid silicon substrates, thesurface free energies and the affinity for the substrate of the two blocks induced a parallelorientation of the domains with respect to the substrate. The films became frustrated when thefilm thickness was not commensurate with the equilibrium domain spacing. This resulted in achange in the domain spacing. However, if the frustration was large enough, the film formedislands or holes after longer annealing times. In addition, a change in morphology wasobserved after extensive annealing times, which resulted in the disappearance of the surfacerelief structures. The microdomain morphology in thin films was visualized by atomic forcemicroscopy, after the selective removal of the organic matrix phase by means of an oxygenplasma etch. A transition was observed from parallel oriented cylinders to hexagonallypacked domains when the annealing time was exceeding several days.

Asymmetric isoprene-block-ferrocenylsilane copolymers were considered in Chapter7. Three diblocks were synthesized, consisting of ferrocenylsilane volume fraction of 20, 24,and 28, respectively. All three copolymers exhibited a cylindrical structure in the bulk. Thinfilms of these copolymers displayed unique morphological changes for each composition as afunction of film thickness. The block copolymer consisting of 28 vol% cylinders displayed awormlike structure for all film thicknesses. When the film thickness was below the domainperiod, holes were formed so that the remaining film still exhibited the wormlikemorphology. The diblock with 24 vol% of the organometallic block also displayed awormlike structure, unless the film thickness was approximately equal to the domain spacing.In that case also areas containing hexagonally packed domains were observed. For thediblock containing 20 vol% ferrocenylsilane, a film thickness approximately equal to thedomain spacing, resulted in a surface morphology completely composed of the hexagonalstructure. For thicker films of all three compositions, a wormlike morphology was observedat the surface. The crystallization of the ferrocenylsilane block occurred even at roomtemperature. This resulted in the slow development of large hedritic features.

The use of ferrocenylsilane polymers combined with self-assembly as a strategytowards nanostructuring surfaces was described in Chapter 8. As thin organic-block-organometallic copolymer films formed lateral patterns on solid substrates, they can beemployed to serve as lithographic masks with dimensions of sub-100 nm level. Oxygenreactive ion etching selectively removed the organic material and converted theorganometallic material into a non-volatile oxide. A subsequent fluorocarbon plasmatreatment etched the unexposed substrate areas whereas the organometallic nanodomainswere insensitive towards removal. It was demonstrated that the thin block copolymer filmscould be used as templates for nanolithography.

Samenvatting

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Het onderzoek dat dit proefschrift beschrijft, omvat de synthese en karakterisering vande relatief nieuwe ferrocenyldimethylsilaan homopolymeren en blok copolymeren. Door deaanwezigheid van ijzer en silicium in de polymere hoofdketen bezitten deze materialeneigenschappen die duidelijk verschillen van organische polymeren. Hoofdstuk 1 beschrijft hetonderwerp van deze thesis, namelijk het onderzoeken en combineren van de nieuweeigenschappen van deze organometaal polymeren met blok copolymeer self-assembly.

Hoofdstuk 2 geeft een korte inleiding in de onderwerpen die relevant zijn voor ditpromotieonderzoek. Het kristallisatie- en smeltgedrag van semi-kristallijne polymerenalsmede de fasenscheiding en self-assembly van blok copolymeren wordt beschreven.Speciale aandacht wordt besteed aan dunne films van blok copolymeren, blends van blokcopolymeren met overeenkomstige homopolymeren en conformatie asymmetrie in blokcopolymeren. Eveneens wordt de achtergrond van de toegepaste technieken belicht.

Hoofdstuk 3 beschrijft de anionische ring-opening polymerisatie van dimethylsilyl-ferrocenophaan. Deze polymerisatie leidde tot de synthese van een serie goed-gedefinieerdeferrocenylsilaan homopolymeren met gecontroleerde molecuulmassa’s. De karakterisering enhet thermische gedrag van deze serie werd eveneens onderzocht. De semi-kristallijnemacromoleculen vertoonden een smelt-herkristallisatie proces gedurende het opwarmenvanuit de gekristalliseerde toestand. Het optreden van dit proces werd aangetoond met behulpvan differential scanning calorimetry experimenten, welke meerdere smeltovergangen lietenzien. Resultaten van kleine-hoek röntgen-verstrooiingsexperimenten bevestigden dat tijdenshet proces dunne lamellen herkristaliseren tot dikkere.

Hoofdstuk 4 beschrijft het effect van reactive ion etching op poly(ferrocenylsilaan).Een zuurstof plasma behandeling resulteerde in de vorming van een dunne oxide laag, rijkaan ijzer en silicium, aan het oppervlak van het polymeer. Door de vorming van deze oxidelaag was de etssnelheid van het organometaal polymeer veel lager dan de etssnelheid vanorganische polymeren. De oxide laag bleek verder ook zeer resistent tegen fluorkoolstofplasma behandelingen, welke vaak gebruikt worden om substraten als silicium te etsen. Doorde hoge etsselectiviteit is het mogelijk om het polymeer als masker te gebruiken om eenpatroon in het substraat over te brengen.

Hoofdstuk 5 beschrijft de synthese en self-assembly van styreen-blok-ferrocenylsilaancopolymeren. Blok copolymeren werden gesynthetiseerd door middel van anionischepolymerisatie van achtereenvolgens styreen en silylferrocenophaan. Deze syntheseresulteerde in bijna monodisperse materialen met gecontroleerde samenstelling enmolecuulmassa. Fasenscheiding van deze blok copolymeren en hun blends met

142

homopolymeer leidde tot regelmatige morfologieën (bollen in een bcc structuur, hexagonaalgepakte cylinders, double gyroid en alternerende lamellen). Om inzicht te krijgen in order-disorder overgangen werden ten slotte rheologische experimenten uitgevoerd. De self-assembly van styreen-blok-ferrocenylsilaan copolymeren bleek een geschikte strategie te zijnvoor het verkrijgen van regelmatige organometale structuren op een nanometer schaal.

Hoofdstuk 6 beschrijft de bestudering van dunne films van een styreen-blok-ferrocenylsilaan copolymeer. De morfologie van het diblok bestaat uit hexagonaal gepakteferrocenylsilaan cylinders in een organische polystyreen matrix. De oppervlakte vrije energieen de affiniteit van de twee blokken voor het substraat resulteerden in een parallelle orientatievan de domeinen ten opzichte van het substraat. Indien de filmdikte niet overeenstemde metde evenwichts domeinafstand ontstond er frustratie in de copolymeer film. Hierdoor werd dedomeinafstand veranderd. Echter bij een te grote frustratie vormde de film na langereannealingstijden eilanden of gaten. Door de vorming van deze reliëfstructuren konden beideblokken bij het gewenste grensvlak aggregeren en bleef tegelijkertijd de evenwichtsdomeinafstand gehandhaafd. Na zeer lange annealingstijden werd een verandering in demicrostructuur waargenomen, welke gepaard ging met de verdwijning van de eilanden engaten. Nadat de organische fase selectief was verwijderd door middel van een zuurstofplasma ets stap kon de microstructuur worden bestudeerd met behulp van atomic forcemicroscopy (AFM). Hieruit bleek dat de parallel liggende cylinders transformeerden inhexagonaal geplaatste domeinen.

Hoofdstuk 7 beschrijft drie asymmetrische isopreen-blok-ferrocenylsilaancopolymeren, bestaande uit ferrocenylsilaan volume fracties van respectievelijk 20, 24 en28%. Alle drie copolymeren vertoonden een cylindrische structuur in de bulk, terwijl voorelke compositie unieke morfologische veranderingen als functie van de filmdikte kondenworden waargenomen. Het diblok met 28% van het organometaal blok vertoonde eenwormachtige structuur voor alle filmdiktes. Bij een filmdikte kleiner dan de domein afstandvormde de film gaten zodat de wormachtige structuur in de rest van de film kon wordengehandhaafd. Het diblok met 24% van het organometaal blok liet eveneens een wormachtigestructuur zien, tenzij de filmdikte ongeveer gelijk was aan de domeinafstand. In dat gevalwerden ook gebieden met hexagonaal gepakte domeinen waargenomen. Het diblok met 20%van het organometaal blok vertoonde een oppervlakte morfologie volledig bestaand uithexagonaal gepakte domeinen bij een filmdikte die gelijk was aan de domein afstand. Vooralle drie copolymeren werd een wormachtige morfologie waargenomen bij dikkere films.Uiteindelijk verstoorde de kristallisatie van het ferrocenylsilaan blok de morfologie van defilms.

Hoofdstuk 8 beschrijft het gebruik van ferrocenylsilaan polymeren gecombineerd metblok copolymeer self-assembly om nano-gestructureerde oppervlakken te verkrijgen. Doordatde organisch-blok-organometaal copolymere films laterale patronen vormden op substraten,konden ze als nanolithografische maskers dienen. Met behulp van zuurstof reactive ionetching werd de organische fase selectief verwijderd en het organometaal materiaal tot eenoxide omgezet. Doordat de geoxideerde nanodomeinen ongevoelig zijn voor fluorkoolstofplasma behandeling konden regelmatige patronen in substraten worden overgebracht.

Publications List

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Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J. Langmuir 2000, in press.“Morphology and crystallization of thin films of asymmetric organic-organometallic diblockcopolymers of isoprene and ferrocenylsilane”

Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J. Abstract ACS, PMSE Preprint, ICIStudent Award competition 2000, in press.“Poly(ferrocenylsilanes) at the interface of chemistry, materials science, andnanotechnology”

Péter, M.; Hempenius, M. A.; Lammertink, R. G. H.; van Os, M. T.; Vancso, G. J. AbstractACS, Polymer Preprint, 2000, in press.“Electrochemical AFM on surface grafted poly(ferrocenylsilanes)”

Lammertink, R. G. H.; Hempenius, M. A.; Van den Enk, J. E.; Chan, V. Z.-H.; Thomas, E.

L.; Vancso, G. J. Advanced Materials, 2000, 12, 98-103.

“Nanostructured thin films of organic-organometallic block copolymers: one-step lithography

with poly(ferrocenylsilanes) by reactive ion etching”

Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J.; Chan, V. Z.-H.; Thomas, E. L.

Abstract ACS, PMSE preprint, 1999, 81, 14-15.

“Organometallic nanodomains from block copolymers for lithographic applications”

Peter, M.; Lammertink, R. G. H.; Hempenius, M. A.; van Os, M.; Beulen, M. W. J.;

Reinhoudt, D. N.; Knoll, W.; Vancso, G. J. Chemical Communications 1999, 359-360.

“Synthesis, characterization and thin film formation of end- functionalized organometallic

polymers”

Lammertink, R. G. H.; Hempenius, M. A.; Thomas, E. L.; Vancso, G. J. Journal of Polymer

Science Part B-Polymer Physics 1999, 37, 1009-1021.

“Periodic organic-organometallic microdomain structures in poly(styrene-block-

ferrocenyldimethylsilane) copolymers and blends with corresponding homopolymers”

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Lammertink, R. G. H.; Hempenius, M. A.; Manners, I.; Vancso, G. J. Macromolecules 1998,31, 795-800.

“Crystallization and melting behavior of poly(ferrocenyldimethylsilanes) obtained by anionic

polymerization”

Lammertink, R. G. H.; Versteeg, D. J.; Hempenius, M. A.; Vancso, G. J. Journal of Polymer

Science Part A-Polymer Chemistry 1998, 36, 2147-2150.

“Morphology control of organometallic domains in phase-separated poly(styrene-block-

ferrocenyldimethylsilanes)”

Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J. Macromolecular Chemistry and

Physics 1998, 199, 2141-2145.

“Crystallization kinetics and morphology of poly(ferrocenyldimethylsilane)”

Hempenius, M. A.; Lammertink, R. G. H.; Vancso, G. J. Macromolecular Symposia 1998,127, 161-163.

“Crystallization and melting behaviour of poly(ferrocenylsilanes)”

Hempenius, M. A.; Lammertink, R. G. H.; Vancso, G. J. Macromolecules 1997, 30, 266-272.

“Side-chain liquid-crystalline polysiloxanes via anionic polymerization: (n-

undecyloxy)arenecarboxylic acid mesogens linked to poly(dimethylsiloxane-co-

methylvinylsiloxane)”

Hempenius, M. A.; Lammertink, R. G. H.; Vancso, G. J. Macromolecular Rapid

Communications 1996, 17, 299-303.

“Well-defined side-chain liquid-crystalline polysiloxanes”

SUBMITTED

Hempenius, M. A.; Robins, N. S.; Lammertink, R. G. H.; Vancso, G. J. MacromolecularRapid Communications, submitted.“Organometallic polyelectrolytes: synthesis, characterization and layer-by-layer deposition ofcationic poly(ferrocenyl(3-ammoniumpropyl)methylsilane)”

Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J.; Shin, K.; Sokolov, J.; Rafailovich,M. H. Macromolecules, submitted.“Morphology and surface relief structures of asymmetric poly(styrene-block-ferrocenylsilane) thin films”

Publications List

145

Lammertink, R. G. H.; Hempenius, M. A.; Vancso, G. J. Chemistry of Materials, submitted.“Poly(ferrocenyldimethylsilanes) for reactive ion etch barrier applications”

Wübbenhorst, M.; van Turnhout, J.; Hempenius, M. A.; Lammertink, R. G. H.; Vancso, G. J.Macromolecules, submitted.“Molecular and cooperative dynamics of low-Tg amorphous and liquid crystallinepolysiloxanes studied with broadband dielectric spectroscopy”

Dankwoord

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Ondanks dat een dankwoord relatief het minste werk van het proefschrift in beslagneemt, wordt het waarschijnlijk het meest gelezen. Daarom zal ik ook via deze weg demensen bedanken die mijn werk gesteund en verbeterd hebben, danwel aangenamer hebbengemaakt.

Voor mijn promotor, professor Vancso, geldt dat hij heeft bijgedragen aan alle driebovengenoemde punten. Julius, ontzettend bedankt voor je enthousiasme en begeleiding; jegaf me alle vrijheid en je deur stond altijd voor me open. Ik vond het fantastisch dat je me demogelijkheid gaf om ervaring op te doen bij buitenlandse universiteiten. Dit laatste heeft eenenorme impact gehad op het verloop van het onderzoek en zal ook mijn verdere loopbaanbeïnvloeden.

Mark Hempenius was er altijd voor de directe begeleiding. Zonder jouw chemischekennis en doorzettingsvermogen zou dit werk ondenkbaar geweest zijn. Het hak- en bakwerkin het lab was van uitzonderlijke kwaliteit.

During 1998 I had the fortune to visit Prof. E. L. Thomas and his group at theMassachusetts Institute of Technology in Cambridge, MA. I still look back on that projectwith great pleasure, also thanks to my MIT-colleagues Chinedum, Cheolmin, Sam,Augustine, Benita, Yoel, Stephane, and Vanessa. Daarnaast wil ik Joost, Raoul, Sam Adams,Jim en Jill bedanken. Zonder jullie zou mijn verblijf in Boston lang niet zo gezellig zijngeweest.

A short visit to the State University of New York in Stony Brook in the group of Prof.M. H. Rafailovich proved to be very fruitful as well. I especially thank Kwanwoo Shin withwhom I spend a lot of beamtime at the NSLS in Brookhaven.

Voor het nodige papierwerk kon ik altijd rekenen op de hulp van Karin, Sandra,Joyce, Gerda, en Geneviève. Clemens en John ben ik veel verschuldigd voor de technischeassistentie en hulp bij allerlei problemen. Ook binnen het MESA+ onderzoeks instituuthebben vele mensen mij de afgelopen jaren geholpen. Met name wil ik Rico en Markbedanken voor de vele TEM sessies, Albert voor de XPS metingen en de cleanroom mensenvoor hun assistentie. De goede werksfeer zorgde ervoor dat ik geen dag ‘s ochtends mettegenzin op de fiets stapte. Hiervoor wil ik al mijn (ex)-collega’s in de MTP groep bedanken.

Naast alle hulp die ik ontving op het wetenschappelijke vlak, zijn de “social events”natuurlijk ook van enorme waarde. Voor de nodige ontspanning en fysieke inspanning was erhet zaalvoetbal op de vrijdagmiddag met Luuk, Mark, Gerard, Wilco en Ype. Ook destandaard uit de hand lopende MTP-borrels zullen mij altijd bij blijven. Mijn oud-huisgenoten Nathan en Rob en de rest van de “buurt” ben ik dankbaar voor hetveraangenamen van mijn studie- en promotietijd met ieder jaar weer het zeilweekend in

148

Friesland als hoogtepunt. Dan zijn er nog mijn paranimfen Jason en Pieter, met wie ik tweehobbies deel, respectievelijk “spicy food” en motoren. Jason, bedankt voor het grondigdoorlezen van dit proefschrift en de vele interessante koffie-discussies. Pieter, bedankt voorje inzet tijdens de vele paranimfen-overleggen en je avontuurlijke emails op demaandagochtend.

Vaak worden de belangrijkste mensen het laatst genoemd. Mijn ouders wil ikbedanken voor alles wat ze voor mij gedaan hebben. Jullie interesse voor mijn werk was vanenorme waarde voor mij. Voor alles en nog veel meer kon ik altijd terugvallen op Cindy. Jijzorgde voor het nodige tegengas en liet me zien wat echt belangrijk is.

Curriculum Vitae

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Rob Lammertink werd geboren op 3 april 1972 te Enter. Na het behalen van het VWOdiploma in 1996 aan het Pius X college te Almelo, begon hij met de studie ChemischeTechnologie aan de Universiteit Twente. Gedurende deze studie heeft hij twee stageopdrachten uitgevoerd, respectievelijk bij BASF in Ludwigshafen en ATO-DLO inWageningen. Het onderwerp van zijn afstudeerproject in de groep van prof. dr. G. J. Vancsoomvatte de synthese en karakterisering van vloeibaar kristallijne poly(siloxanen). Na hetvoltooien van zijn studie op 3 april 1996, begon hij in diezelfde groep als assistent inopleiding. Onder de supervisie van prof. dr. G. J. Vancso en dr. M. A. Hempenius onderzochthij de synthese en karakterisering van poly(ferrocenylsilanen). Gedurende dit promotieonderzoek is er met diverse binnen- en buitenlandse universiteiten samengewerkt. Vanjanuari tot juni 1998 verbleef Rob Lammertink aan het Massachussetts Institute ofTechnology (MIT) te Cambridge waar hij in de groep van prof. dr. E. L. Thomas defasenscheiding en self-assembly van blok copolymeren met behulp van transmissieelectronen micriscopie en kleine hoek röntgen verstrooiing onderzocht. Voor dit werkbezoekontving hij een NWO beurs. Tijdens een kort werkbezoek in 1999 in de groep van prof. M.H. Rafailovich aan de State University of New York (SUNY) in Stony Brook heeft hij metbehulp van röntgen reflectiviteit dunne blok copolymere films onderzocht. Rob Lammertinkis als finalist geselecteerd voor de ICI Student Award in Applied Polymer Science competitie2000.

POLY(FERROCENYLDIMETHYLSILANES) AT THE ...IV Poly(ferrocenyldimethylsilanes) at the Interface of...

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