An Investigation Into the Structural Design of Polymer Thin Films: from Stimuli Responsive Polyampholyte Brushes to Polymer- Grafted Nanocomposites

University of Tennessee, Knoxville TRACE: Tennessee Research and Creative Exchange

Doctoral Dissertations Graduate School

8-2018

An Investigation into the Structural Design of Polymer Thin Films: From Stimuli Responsive Polyampholyte Brushes to Polymer- Grafted Nanocomposites

Rachel Irene Ramirez University of Tennessee, [email protected]

Follow this and additional works at: https://trace.tennessee.edu/utk_graddiss

Recommended Citation Ramirez, Rachel Irene, "An Investigation into the Structural Design of Polymer Thin Films: From Stimuli Responsive Polyampholyte Brushes to Polymer-Grafted Nanocomposites. " PhD diss., University of Tennessee, 2018. https://trace.tennessee.edu/utk_graddiss/5020

This Dissertation is brought to you for free and open access by the Graduate School at TRACE: Tennessee Research and Creative Exchange. It has been accepted for inclusion in Doctoral Dissertations by an authorized administrator of TRACE: Tennessee Research and Creative Exchange. For more information, please contact [email protected]. To the Graduate Council:

I am submitting herewith a dissertation written by Rachel Irene Ramirez entitled "An Investigation into the Structural Design of Polymer Thin Films: From Stimuli Responsive Polyampholyte Brushes to Polymer-Grafted Nanocomposites." I have examined the final electronic copy of this dissertation for form and content and recommend that it be accepted in partial fulfillment of the equirr ements for the degree of Doctor of Philosophy, with a major in Chemistry.

S. Micheal Kilbey II, Major Professor

We have read this dissertation and recommend its acceptance:

Ampofo K. Darko, Emmanouil Doxastakis, Bin Zhao

Accepted for the Council:

Dixie L. Thompson

Vice Provost and Dean of the Graduate School

(Original signatures are on file with official studentecor r ds.) An Investigation into the Structural Design of Polymer Thin

Films: From Stimuli Responsive Polyampholyte Brushes to

Polymer-Grafted Nanocomposites

A Dissertation Presented for the

Doctor of Philosophy

Degree

The University of Tennessee, Knoxville

Rachel Irene Ramirez

August 2018

Dedication

This dissertation is dedicated to my loving parents, Ada Caro and Rafael

Ramirez, my fiancée Joshua Moncada, my brothers, Luis and Miguel Santiago,

and my dog, Maximus.

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Acknowledgements

Firstly, I would like to extend my sincerest and everlasting gratitude to my advisor, Prof. S. Michael Kilbey II. His determination, guidance, and encouragement, have molded me into a responsible researcher that can express clear, concise ideas and results. My years in Dr. Kilbey’s lab have been instrumental in my development as a materials scientist, not only filled with professionalism but filled with laughter, as well, that has led to fostering friendships with my group members and with Dr. Kilbey. I will forever be grateful to Dr. Kilbey for his mentorship and friendship as I continue my career.

I would also like to acknowledge my family, friends, and colleagues who have supported and encouraged me during my entire journey to receiving my

Ph.D. in polymer chemistry at The University of Tennessee. First, I would like to thank my parents, Ada Caro and Rafael Ramirez, for their full support. Even when the road became rocky and I became very discouraged, their encouragement through the hard times and unconditional love are what kept me motivated to move forward. Both of you are not only fantastic parents but incredible people and I can honestly say that I would not be the person I am today without you. I would also like to thank my brothers Luis and Miguel, for offering both words of advice and a kick in the pants when I needed it. I would also like to thank my fiancé Joshua

Moncada, who has kept me grounded during the most difficult times in graduate school. His love and support mean more to me than he will ever know.

Special thanks go to Dr. Jeremiah Woodcock for providing insight and background on polyampholyte brushes. Dr. John Dunlap of the Advanced

Microscopy and Imaging Center at The University of Tennessee-Knoxville is gratefully acknowledged for his assistance with atomic force microscopy studies. I truly thank Dr. Xu Wang for establishing the foundation for many polymeric synthetic techniques.

Finally, I am grateful to all of the members (past and present) of the Kilbey research group for their friendship, support, advice, and engaging conversations.

These members include Jesse Davis, Zach Seibers, Graham Collier, Kamlesh

Bornani, Bethany Aden, Dayton Street, Elizabeth O’Connell, Will Ledford, Natalie

Czarnecki, and Sina Sabury. This work would not be possible without the financial support from the National Science Foundation and Honeywell. Thank you all for your help.

Abstract

Understanding how the structural design of polymer thin films impacts their response to stimuli, such as pH, is crucial for developing systems with targeted attributes to further expand the scope of their applications. These polymer systems can be modified based on the choice of monomers, composition, and structural design, which provides a tunable source for both functionality and tailorability. This tailorability in design gives rise to a wide array of molecular properties that have a significant impact on the macroscopic properties of thin films. This dissertation work aims to provide insight into how the structural design can impact properties on two categories of polymer thin films: copolymer-grafted nanocomposites and polymer brushes.

Copolymer nanocomposites were investigated on how the miscibility of the copolymer-grafted nanoparticles can be tuned by using the enthalpic interactions between the graft and the polymer matrix. Changing the overall composition of the copolymer allowed us to drive dispersion of the resulting nanocomposites in the matrix. The copolymer-grafted nanoparticles were synthesized using surface- initiated activators regenerated via electron transfer atom transfer radical polymerization in which poly(methyl methacrylate-r-cyclohexyl methacrylate) was grown from the silica nanoparticle surface and dispersed in a chemically dissimilar polystyrene matrix. An investigation into how the thermomechanical properties of the resulting copolymer-grafted nanocomposites was conducted using fused deposition modeling.

The second category of polymer thin films examined was polyampholyte brushes in which the impact of modifying the composition on swelling behavior was investigated. Polyampholytes are comprised of charge-positive and charge- negative repeat units, which directly contributes to trade-offs between charge which is externally regulated by solution pH and added salt, and structure. A series of swelling studies were performed to examine how copolymer composition affects structural response of random polyampholyte brushes as pH is changed and betaine, a zwitterion, is added.

The work in this dissertation involves the investigation of several types of polymer thin films, the common theme is clarifying how the structural design and composition affects the properties of polymer brushes, both as copolymer-grafted nanocomposites and on planar surfaces. In total, this research provides insight into how polymer design, polymer structure, and behavior responses are associated.

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Table of Contents

1.1: Introduction ...... 1

1.1 Universal Importance of Polymer Thin Films ...... 2

1.2 Research Objectives ...... 4

1.3 Thin Film Copolymer-grafted Nanocomposites ...... 8

1.3.1 Miscibility of copolymer-grafted nanoparticles in a chemically dissimilar

matrix ...... 8

1.3.2 Overview of fused deposition modeling ...... 12

1.3.3 Scaling of polymer grafted-nanoparticles via the grafting “through”

approach ...... 15

1.4 Polyampholyte Brushes Synthesized via ARGET ATRP ...... 17

1.5 Organization of Dissertation ...... 19

Chapter 2: Dispersion of Polymer-grafted Nanoparticles Poly(methyl methacrylate- r-cyclohexyl methacrylate) Synthesized via Surface-initiated ARGET ATRP in a

Chemically Dissimilar Matrix ...... 21

2.1 Abstract ...... 22

2.2 Introduction ...... 23

2.3 Experimental Section ...... 28

2.3.1 Materials and preparations ...... 28

2.3.2 General characterization ...... 28

2.3.3 Synthesis of poly(methyl methacrylate-r-cyclohexyl methacrylate)

(P(MMA-r-CHMA)) via ARGET ATRP ...... 29

2.3.4 Synthesis of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane ..... 30

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2.3.5 Activation of nanoparticles using 3-(2-bromoisobutyramido)propyl

(trimethoxy)silane ...... 31

2.3.6 Synthesis of polymer-grafted nanoparticles via surface-initiated

ARGET ATRP ...... 32

2.3.7 Preparation of copolymer nanocomposites ...... 33

2.4 Results and Discussion ...... 33

2.4.1 Kinetics of ARGET-ATRP of P(MMA-r-CHMA) in solution ...... 33

2.4.2 Synthesis of P(MMA-r-CHMA)-grafted nanoparticles ...... 39

2.4.3 Miscibility of copolymer-grafted nanoparticles in a chemically distinct

matrix ...... 45

2.5 Conclusions ...... 49

Chapter 3: Fused Deposition Modeling of Poly(methyl methacrylate-r-cyclohexyl methacrylate) Copolymer-grafted Nanoparticles in a Chemically Dissimilar Matrix

...... 51

3.1 Abstract ...... 52

3.2 Introduction ...... 52

3.3 Experimental Section ...... 55

3.3.1 Materials and preparations ...... 55

3.3.2 General characterization ...... 55

3.3.3 Synthesis of 3-(2-Bromoisobutyramido)propyl(trimethoxy)silane .... 56

3.3.4 Immobilization of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane

onto silica nanoparticles ...... 56

3.3.5 Synthesis of polymer grafted nanoparticles via surface initiated

ARGET ATRP ...... 57

3.3.6 Preparation of the copolymer nanocomposite filament and FDM

printing ...... 58

3.4 Results and Discussion ...... 59

3.4.1 Optimizing printing conditions for pure polystyrene and copolymer-

grafted nanocomposites ...... 59

3.5 Conclusions ...... 68

Chapter 4: Scalable Preparation of Poly(methyl methacrylate)-grafted

Nanoparticles via the Grafting “Through” Approach ...... 69

4.1 Abstract ...... 70

4.2 Introduction ...... 70

4.3 Experimental section ...... 73

4.3.1 Materials ...... 73

4.3.2 General characterization techniques ...... 73

4.3.3 Immobilization and activation of nanoparticles using 3-

(trimethoxysilyl)propyl methacrylate ...... 74

4.3.4 Synthesis of poly(methyl methacrylate)-grafted nanoparticles via

conventional free radical polymerization ...... 75

4.4 Results and Discussion ...... 77

4.4.1 Synthesis of PMMA-grafted nanoparticles by using SPM and VTS as

activators ...... 77

4.5 Conclusions ...... 86

Chapter 5: ARGET-ATRP Synthesis and Swelling Response of Compositionally

Varied Poly(methacrylic acid-co-N, N-diethylaminoethyl methacrylate)

Polyampholyte Brushes ...... 87

5.1 Abstract ...... 88

5.2 Introduction ...... 89

5.3 Experimental Section ...... 95

5.3.1 Materials ...... 95

5.3.2 General characterization methods ...... 95

5.3.3 Synthesis of poly(tert-butyl methacrylate-co-diethylaminoethyl

methacrylate) (P(tBMA-co-DEAEMA)) via ARGET ATRP ...... 97

5.3.4 Synthesis of 3-(2-Bromoisobutyramido)propyl(trimethoxy)silane .... 98

5.3.5 Immobilization of the ARGET ATRP silane initiator onto silicon

surfaces...... 98

5.3.6 Synthesis of polymer brushes via surface initiated ARGET ATRP .. 99

5.3.7 Generation of P(MAA-co-DEAEMA) brushes via post-polymerization

modification ...... 101

5.3.8 Ellipsometric swelling studies of P(MAA-co-DEAEMA) brushes ... 101

5.4 Results and Discussion ...... 102

5.4.1 Kinetic studies of solution synthesis P(tBMA-co-DEAEMA)

copolymers ...... 102

5.4.2 Synthesis of P(MAA-co-DEAEMA) polymer brushes ...... 107

5.4.3 Swelling studies of polyampholyte brushes ...... 111

5.4.4 Swelling response in the presence of betaine zwitterion ...... 120

5.5 Conclusions ...... 122

Chapter 6: Summary, Conclusions, and Future Work ...... 124

6.1 Summary and Conclusions ...... 125

6.2 Future Work ...... 127

List of Referances ...... 130

Appendices ...... 162

Appendix A - Chapter 3: Fused deposition modeling of Poly(methyl methacrylate- r-cyclohexyl methacrylate) Copolymer-grafted Nanoparticles in a Chemically

Dissimilar Matrix ...... 163

Appendix B - Chapter 5: ARGET-ATRP Synthesis and Swelling Response of

Compositionally Varied Poly(methacrylic acid-co-N, N-diethylaminoethyl methacrylate) Polyampholyte Brushes ...... 171

Vita ...... 178

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List of Tables

Table 2-1. Macromolecular characteristics of P(MMA-r-CHMA) polymerized in solution...... 35

Table 2-2. Characteristics of copolymer-grafted nanoparticles of P(MMA-r-CHMA)

...... 42

Table 3-1. Average Young’s Modulus of pure PS printed at various temperatures

...... 63

Table 3-2. Average Young’s modulus of 1wt % bare nanocomposites printed at various temperatures ...... 65

Table 3-3. Average values of Young’s modulus of 1 wt % 90:10 CHMA:MMA copolymer nanocomposites printed at various temperatures ...... 67

Table 4-1. Characteristics of PMMA-grafted nanoparticles ...... 79

Table 5-1. Characteristics of P(tBMA-co-DEAEMA) copolymers ...... 104

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Table 5-2. Characterization data for P(MAA-co-DEAEMA) and dry film thicknesses ...... 109

Table 5-3. Comparison of “parent” and polyampholyte brush thickness and molecular weight ratios ...... 111

Table B-1. Amounts of buffer species and sodium chloride required to make 100 mL buffer solution of 30 mM ionic strength at pH values given ...... 174

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List of Figures

Figure 1-1. Examples of polymer thin films created from homopolymer and mixed polymer brushes, films formed by layer-by-layer deposition, block copolymer films, films containing functional particles, and polymer nanocomposites...... 3

Figure 1-2. Schematic of a typical FDM setup ...... 13

Figure 1-3. A general depiction of different surface modification techniques of grafting polymers to nanoparticles: (A) “grafting to”, (B) “grafting from”, and (C)

“grafting through” methods ...... 16

Figure 2-1. Kinetic plots for ARGET ATRP of P(MMA-r-CHMA) at various comonomer ratios. For plots A and B, the copolymer was targeted to be 10% MMA

(90% CHMA); for C and D the copolymer was targeted to be 75% MMA (25%

CHMA); and E and F the copolymer was targeted to be 90% MMA (10% CHMA).

All reactions performed in anisole with A and B at 50 ⁰C while the others were done at 70 ⁰C ...... 36

Figure 2-2. 1H NMR spectra for ARGET ATRP of P(MMA-r-CHMA) at a composition of 75/25 MMA/CHMA. An aliquot was taken every 1 hr and a total of

7 aliquots were acquired. The scheme above the spectra indicates the signature

xv peaks for the monomers and the copolymer: A 4.85 ppm (monomer peak for

CHMA); B 3.76 ppm (monomer peak for MMA); C 4.65 ppm (CHMA in polymer), and D 3.55 (MMA in polymer) ...... 38

Figure 2-3. Representative 1H NMR spectrum of the free random copolymer comprised of 75/25 MMA/CHMA. C and D correlate to the protons located on the polymer chain, C represents the hydrogen in CHMA and D represents the hydrogen in MMA...... 41

Figure 2-4. Thermogravimetric analysis data for bare nanoparticles, nanoparticles activated with 3-(2-bromoisobutyramido)propyl(trimethoxy)silane and 10/90

MMA/CHMA copolymer-grafted nanoparticles ...... 44

Figure 2-5. AFM images of p(MMA-r-CHMA) nanocomposites. To form the nanocomposites, copolymer-grafted nanoparticles were incorporated into 50k polystyrene matrix at 5 wt %. The copolymer compositions of MMA:CHMA are as follows: 10:90 (top), 75:25 (top middle), 80:20 (bottom middle), 90:10 (bottom).

Each composition was imaged as cast and then as a function of annealing time.

Samples were annealed at 150 ⁰C under vacuum for up to 120 hr ...... 47

Figure 3-1. Luzbot mini FDM printer ...... 53

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Figure 3-2. Thermogravimetric analysis of pure PS pellets indicates onset of degradation is at ~300 ⁰C...... 60

Figure 3-3. Representative stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS across temperatures ranging from 230 ⁰C to 270 ⁰C in 5 degree increments ...... 62

Figure 3-4. Representative stress-strain curves measured by DMA used to determine the Young’s modulus of the printed nanocomposites comprised of 1 wt

% bare silica nanoparticles in a PS matrix at 235 ⁰C, 265 ⁰C, and 270 ⁰C ...... 65

Figure 3-5. Representative stress-strain curves for the printed nanocomposites comprised of 1 wt % 90:10 CHMA:MMA grafted nanoparticles in a PS matrix at

265 ⁰C and 270 ⁰C ...... 67

Figure 4-1. A general depiction of different surface modification techniques used to graft polymers to nanoparticles: (A) “grafting to”, (B) “grafting from”, and (C)

“grafting through” may be used for the attachment of polymer chains to surfaces

...... 72

Figure 4-2. Thermogravimetric analysis of bare nanoparticles, activated nanoparticles decorated with (A) SPM and (B) VTS, and the resulting PMMA- grafted nanoparticles made by grafting through ...... 81

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Figure 4-3. FTIR-ATR spectra of bare nanoparticles (blue), PMMA grafted nanoparticles (red), and pure PMMA (black). The spectra are offset vertically

(stacked) for clarity ...... 83

Figure 4-4. Distribution of hydrodynamic radius, Rh, measured at four different angles for (A) bare nanoparticles, (B) SPM-activated nanoparticles, and (C)

PMMA-grafted nanoparticles in THF at a concentration of 0.5 mg/mL...... 85

Figure 5-1. Kinetic plots for ARGET ATRP of P(tBMA-co-DEAEMA) at various comonomer ratios. Plots A and B are for a copolymer targeted to be 80 mol % tBMA and 20 mol % DEAEMA, C and D are for the copolymer targeted to be 50 mol % tBMA and 50 mol % DEAEMA, and E and F are for the copolymer targeted to be 20 mol % tBMA and 80 mol % DEAEMA. All reactions were run in anisole at

T=35 °C for 8 h...... 105

Figure 5-2. Swelling response of p(DEAEMA-co-MAA) brushes as a function of pH. Amine content varies, with (A) representing 20 mol % DEAEMA, (B) 50 mol %

DEAEMA, and (C) 80 mol % DEAEMA. In each case, the swollen thickness measured by in situ ellipsometry is normalized by the thickness of the brush measured in ambient conditions...... 113

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Figure 5-3. Swelling behavior of p(MAA-co-DEAEMA) brushes of different composition as a function of pH. The swelling behavior of the original curves to the curves obtained from the addition of the zwitterion are compared with the amine content in the polyampholytes increasing from 20% (left) to 80% (right)...... 121

Figure A-1. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 230 ⁰C...... 164

Figure A-2. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 235 ⁰C...... 164

Figure A-3. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 240 ⁰C...... 165

Figure A-4. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 245 ⁰C...... 165

Figure A-5. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 250 ⁰C...... 166

Figure A-6. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 255 ⁰C...... 166

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Figure A-7. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 260 ⁰C...... 167

Figure A-8. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 265 ⁰C...... 167

Figure A-9. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 270 ⁰C...... 168

Figure A-10. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 235 ⁰C...... 168

Figure A-11. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 265 ⁰C...... 169

Figure A-12. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 270 ⁰C...... 169

Figure A-13. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % 90:10

CHMA:MMA grafted nanoparticles in a PS matrix. All samples were run at 265 ⁰C.

...... 170

Figure A-14. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % 90:10

CHMA:MMA grafted nanoparticles in a PS matrix. All samples were run at 270 ⁰C.

...... 170

Figure B-1. Kinetic plots for ARGET ATRP of P(tBMA-co-DEAEMA) at various comonomer ratios. The plots for A and B are for a copolymer targeted to be 70% tBMA and 30% DEAEMA; C and D are for a copolymer targeted to be 60% tBMA and 40% DEAEMA; E and F are for a copolymer targeted to be 40% tBMA and

60% DEAEMA; and G and H are for a copolymer targeted to be 30% tBMA and

70% DEAEMA. All reactions were run in anisole at T=35°C for 8 h ...... 175

Figure B-2. Swelling response of p(DEAEMA-co-MAA) brushes as a function of pH. The curves also display the refractive index from the “slab” model used in the analysis of the ellipsometric data ...... 176

xxi

Figure B-3. AFM images of p(tBMA-co-DEAEMA) brushes (“Protected”) and the p(MAA-co-DEAEMA) polyampholyte brushes resulting from deprotections (acid hydrolysis). The compositions, top to bottom, are 80:20, 50:50, and 20:80

[tBMA]:[DEAEMA] (left) and [MAA]:[DEAEMA] (right). RMS roughness values are inset in each image ...... 177

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List of Schemes

Scheme 2-1. Synthesis of P(MMA-r-CHMA) on silica nanoparticles using ARGET-

ATRP ...... 40

Scheme 4-1. Synthesis of PMMA grafted-nanoparticles from activated nanoparticles using grafting through process with a conventional free radical initiator ...... 78

Scheme 5-1. Synthesis of P(MAA-co-DEAEMA) on silicon substrates using

ARGET ATRP ...... 107

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Chapter 1: Introduction

1.1 Universal Importance of Polymer Thin Films

A polymer thin film is a layer of polymeric material ranging in thickness from a fraction of a nanometer (monolayer) to several micrometers. Due to thinness, polymer thin films have very different physical properties in comparison to bulk polymeric systems. This difference in physical properties is mainly due to the fact that confinement effects are present in thin films, and a high surface-to-volume ratio means that interfacial interactions play a significant role. Thus, as the thickness of the film decreases, deviations from bulk behavior generally widen.

While polymer thin films often are characterized by the thickness of the material, the characteristic geometry of the substrate, topology of the polymer, and its chemical nature are also important. Examples of the variety of polymer thin films can be seen in Figure 1-1. These include polymer brushes on a flat surface, block copolymer thin films, or polymer brushes coating a spherical nanoparticle, which creates a type of nanocomposite. The diversity of polymer thin film systems makes them universal or beneficial in most fields of research. For example, polymer thin films can be applied as antifouling coatings, cell scaffolding or biomaterials,1–4 biosensors,4–7 organic light-emitting or photovoltaic devices,8–11 optical storage or magnetic hybrid materials,12,13 high charge density batteries or functional membranes.14–16 They can be applied to provide corrosion resistance, stabilize colloids or construct controlled release nanocarriers.17,18

Polymer thin films have many potential applications due to the overall tailorability of polymeric materials and their integration into hybrid systems.

Polymers can be tuned based on the choice of monomers, composition, and structural design, which provides a source for both functionality and tailorability.

This gives rise to a wide array of molecular properties that have a significant impact on the macroscopic properties of thin films. In my dissertation research, I have used this tailorability to vary the overall miscibility of copolymer-grafted nanocomposites in a chemically dissimilar matrix and to change stimuli response behavior of random polyampholyte brushes. The capability to manipulate properties or responsiveness of polymer thin films requires an understanding how the macromolecular design of polymers and their confinement impacts interfacial structure and properties. Providing this type of insight is the main theme of this dissertation research.

1.2 Research Objectives

The goal of my dissertation research is to enhance the fundamental understanding of how structural design of confined polymers affects their response. Specifically, my work falls into two areas related to polymer thin films: charged polymer brushes and copolymer grafted nanoparticles. In the area of copolymer-grafted nanoparticles, I examine the miscibility of copolymer grafted nanoparticles in a chemically dissimilar matrix, which forms a copolymer grafted nanocomposite. In this work, the overall composition of the copolymers tethered to the silica nanoparticles is varied in order to tailor the dispersion of the copolymer grafted nanoparticles within the matrix polymer. These systems also are examined

4 in the context of 3D printing. My other area of research involving polymer thin films focuses on polyampholyte brushes. In this work, the swelling response of polyampholyte brushes, which contain repeating units having weak acid and weak base groups, is investigated as a function of copolymer composition.

My studies focused on the miscibility of copolymer-grafted nanoparticles in a dissimilar matrix was accomplished by synthesizing well-defined copolymer- grafted nanoparticles that consisted of a tethered layer of poly(methyl methacrylate-random-cyclohexyl methacrylate) (P(MMA-r-CHMA)) chains on silica nanoparticles. To do this, I used a variant of atom transfer radical polymerization

(ATRP) known as ARGET ATRP (activators regenerated by electron transfer

ATRP) to grow the chains from the nanoparticle surface. Then, the copolymer- grafted nanoparticles having different composition were dispersed in a chemically distinct polystyrene matrix, which creates a copolymer-grafted nanocomposite

(CPGN). My system was chosen so that the graft and matrix polymers exhibit attractive enthalpic interactions, and the strength of that interaction depends on the composition of the tethered copolymers

There are two factors that contribute to the miscibility of polymer blends mixing: entropy of mixing is always negative and, therefore, favors mixing. On the other hand, enthalpic contributions are key because most polymers are immiscible and their repulsive interactions promote phase separation. In order for a polymer blend to be miscible the enthalpic contribution of the polymer must be very small or attractive to create a homogeneous system. Much research has been done on the miscibility of polymer nanocomposites in which the polymer chains that are

5 grafted to the nanoparticles that are the same chemical type as the matrix polymer.19–29 This is considered a “symmetric” system, and because the Flory-

Huggin’s χ parameter is zero, the entropy of mixing dictates the miscibility of the polymer-grafted nanoparticles in the matrix polymers. In this case, the system is mixed (dispersed) when the molecular weight of the grafted polymer is higher than the molecular weight of the matrix chains, and phase separation occurs when graft chains molecular weight is lower than that of the matrix polymer.22,30,31

Not much exploration has been directed toward systems comprised of chemically dissimilar polymers. In my research, I used a system consisting of polystyrene matrix and random copolymers of poly(methyl methacrylate-r- cyclohexyl methacrylate) grafts. Incorporation of cyclohexyl methacrylate provides a weakly attractive interaction that drives compatibilization and miscibility. This type of phenomena has been observed when increasing amounts of suitable functional groups are integrated into to polymer chains, which transforms an immiscible polymer system into a miscible polymer blend.17,32–36 This ability to disperse particles due to graft−matrix miscibility allows for higher filler loadings than what can be achieved in chemically identical graft−matrix composites.37 By tuning the miscibility of a polymer nanocomposite by varying copolymer composition, the strength and sign of enthalpic interactions between the graft and the polymer matrix can be changed, and this should, in turn, lead to improvements in macroscopic properties and more robust materials.

Another theme in this dissertation is to investigate the thermomechanical properties of these copolymer grafted nanocomposites and apply them in a form

6 of 3D-printing known as fused deposition modelling (FDM). To accomplish this, I prepared filaments containing the poly(methyl methacrylate-r-cyclohexyl methacrylate)-grafted nanoparticles at various loading levels in polystyrene, printed samples by FDM, and investigated the impact that the hybrid additives have on thermomechanical properties of the system. An improvement in thermomechanical properties is expected due to the miscibility and dispersion of the copolymer-grafted nanoparticles in the polystyrene matrix. The change in the thermomechanical properties was primarily investigated through the use of dynamic mechanical analysis (DMA).

Associated with the theme of polymer-grafted nanoparticles, I investigated the scalability of synthesizing polymer-grafted nanoparticles using a “grafting through” approach. The grafting through approach uses a self-assembled monolayer that contains polymerizable vinyl groups. When growth of polymer chains is initiated in solution, a surface-bound monomer unit can be integrated into a propagating chain, which permanently anchors the polymer chain to the substrate. This particular approach uses conventional free radical polymerization, which makes it a scalable technique for synthesizing polymer-grafted nanoparticles on the gram scale. I characterized these poly(methyl methacrylate)- grafted nanoparticles using thermogravimetric analysis (TGA), Fourier transform infrared spectroscopy (FTIR), and made some preliminary investigations using dynamic light scattering (DLS).

The final theme of this work is to understand the impact of modifying the composition of surface-anchored polyampholytes on swelling behavior.

Polyampholytes are comprised of charge-positive and charge-negative repeat units, which directly contributes to trade-offs between charge, which is externally regulated by solution pH and added salt, and structure. A series of swelling studies were performed to examine how copolymer composition affects structural response of random polyampholyte brushes as pH is changed and as betaine, a zwitterion, is added. The novel elements of this work were the use of ARGET ATRP to control surface-initiated formation of the polyampholyte brushes and studies focused on how behaviors vary with polyampholyte composition. The swelling responses of the various polyampholyte brushes were examined through the use of multi-angle ellipsometry.

While the work described in this dissertation involves the investigation of several types of polymer thin films, the common theme is developing a clear picture how the structural design and composition affects the properties of polymer brushes, both as the copolymer-grafted nanocomposites and on planar surfaces.

In total, this research provides considerable insight into how polymer design, polymer structure and response are linked.

1.3 Thin Film Copolymer-grafted Nanocomposites

1.3.1 Miscibility of copolymer-grafted nanoparticles in a chemically dissimilar matrix

There are a variety of polymer nanocomposites (PNC) created by the incorporation of either organic (carbon nanotubes,38,39 clay,40,41 graphene42,43) or inorganic (silica,19,44 titania,45,46 metal nanowires47,48) nanomaterials into a polymer melt. These PNCs frequently improve thermomechanical properties compared to the pure polymeric material, and this improvement is mainly achieved when the

8 nanofillers are well-dispersed the polymer matrix.30,49–52 Due to the improvement of thermomechanical properties, PNCs show promise as structural components for the transportation industry, or functional hybrid materials, including systems that offer enhanced electrical properties or providing shielding for electronics.8,53,54

Many of these successes are based on the ability of the nanofillers to add functionality without substantially changing polymer properties or processability.8,20,55–57 Blends of polymers and nanofillers are usually immiscible due to strong van der Waals interactions between the nanofillers. As the filler aggregates and clusters form, there is a detrimental impact on the interfacial contact between the polymer and the filler, causing phase separation and a decrease in the resulting bulk properties. Therefore, finding strategies for efficient dispersion of nanofillers in polymer matrices remains an active area of research.

One successful method to control the dispersion of nanoparticles in a polymer matrix is to shield the particle from particle-particle interactions.31,21,58,59,22,60,61 Theoretical studies of polymer functionalized nanoparticles have shown that the extent of matrix polymer penetration into

(exclusion from) the grafted polymer layer, also termed as “wetting” (“dewetting”), will dictate the extent of dispersion (aggregation) of polymer-grafted particles in the polymer matrix.22,23,33,62,63 Experimental work shows that the chemistry of the grafted polymers, nanoparticles and the polymer matrix play crucial roles in determining the spatial organization of the nanoparticles.19,24–26,64 In symmetric systems, the grafting density of the attached polymer chains dictates how the polymer-grafted nanoparticle behaves once blended in the polymeric matrix. When

9 the polymer-grafted nanoparticles have a high grafting density, the polymer completely shields the surface, perhaps extending into a brush conformation.19,65,66

The polymer-grafted nanoparticles placed in the matrix will disperse (aggregate) if the molecular weight of the matrix polymer is lower (higher) than that of the grafted homopolymer. At low grafting density, portions of the nanoparticle are exposed, and there is a combination of matrix-surface, surface-surface, and graft-surface interactions that regulate dispersion/ aggregation.18,22,23,31,67 As a result, the behavior of the polymer-grafted nanoparticle is dictated by the effective interparticle interactions arising because of surface of the nanoparticle that is exposed versus the fraction that is covered by the grafted chains.23,68–70

Different from homopolymer functionalization, copolymer functionalization produces additional parameters, such as choice of chain design and monomer chemistry, that provide additional control over the assembly of the PGNCs. With the addition of copolymer, the system is now considered an “asymmetric” system.

Typical studies of copolymer grafted nanocomposites (CPGNCs) use a copolymer grafted onto a nanoparticle in a homopolymer matrix, and one comonomer of the grafted chain is usually chemically identical to the matrix in order to promote dispersion. Theory,71–74 simulations,75–77 and experiments have been used to study copolymers grafted onto nanoparticles. Jayaraman et al.71 studied the interactions of copolymer-grafted nanoparticles in a homopolymer matrix using self-consistent PRISM theory and Monte Carlo simulations. They studied the effects of monomer sequence in the grafted chain, using both alternating and diblock polymers, length of grafted and matrix chains, attraction strength between

10 pairs of like monomers, and the nanoparticle diameter. They found that alternating and diblock copolymer-grafted particles exhibit significantly different trends in their potentials of mean force (PMF). Monomer sequence dictates how the attractive monomers would aggregate, and the sizes of those monomer aggregates depends on the intermonomer attraction strength and graft chain length. Zhu et al.74 employed both self-consistent field theory (SCFT) and density functional theory

(DFT) to study a dense system of diblock copolymer-grafted nanoparticles with a single chain grafted onto each particle. In these studies, they did not include matrix chains, instead choosing to simulate the limit of loading of 100% CPGNC. When the particle surface was chemically neutral to the grafted chain, they saw typical block copolymer morphologies (i.e. cylinders and lamellae). These morphologies were determined by both the composition of the copolymer and the size of the particle. Ordered hierarchical structures were observed when the particle interactions were repulsive to both blocks of the copolymer. They observed structures including of cylinders with cylinders at the interfaces, lamellae with cylinders at the interfaces, and lamellae with cylinders inside a domain. These particular hierarchical structures are not typical of block copolymer melts;74 thus this study highlights the potential of copolymer grafts to change nanostructure observed in CPGNCs.

Another type of asymmetric system is one in which the matrix polymer is chemically dissimilar to the copolymers grafted on the nanoparticles. In this case, aggregation/dispersion of the nanoparticles is governed mostly by enthalpic forces, unlike symmetric systems discussed previously, which are largely

11 governed by the entropy of mixing. Theoretical studies of colloids or flat surfaces grafted with a dense polymer brush placed in a chemically dissimilar polymer matrix have shown how the choice of a chemically dissimilar graft and matrix polymers lead either to attractive or repulsive Flory−Huggins χ parameter, which impacts the conformations and the phase separation behavior.17,32–36 Through the use of controlled synthesis, the particle dispersion can be tuned by manipulating the enthalpic interactions between the grafted polymer and matrix chains.37,78,79

This was demonstrated by Ojha et al., who synthesized a system comprised of poly(styrene-r-acrylonitrile) (SAN) grafted nanoparticles and dispersed those copolymer-grafted nanoparticles in a chemically dissimilar poly(methyl methacrylate) (PMMA) matrix. They observed that even when the molecular weight of the grafted-SAN copolymer (Mn = 16,000 g/mol) was well below the matrix molecular weight of Mn = 93,000 g/mol, the nanoparticles remained dispersed within the matrix.37 In this case, there are attractive enthalpic interactions between the copolymer-grafted nanoparticles and matrix chains, driving the mixture to be dispersed. This behavior is distinctive -- a symmetric system having grafted chains with a lower molecular weight than the matrix chains, would aggregate.37 The ability to disperse particles due to graft-matrix miscibility allows for higher filler loadings than chemically identical graft-matrix composites.37

1.3.2 Overview of fused deposition modeling

Three-dimensional (3D) printing is an additive manufacturing (AM) technique that creates 3D objects with unique structure and diverse

12 properties. Currently, there are a variety of techniques used in AM manufacturing, such as fused deposition modeling (FDM),80 stereolithography apparatus (SLA),81 continuous liquid interface production (CLIP),82 digital light processing

(DLP),83 and selective laser sintering (SLS).84 These techniques have been developed to form stereoscopic objects with complex architecture. Presently, FDM has become the most widely used 3D printing method due to it being simple to use, cost effective and environmentally friendly.85–87 FDM printers work by using a controlled extrusion of thermoplastic filaments. When preparing to print with FDM, the thermoplastic filaments are fed into a nozzle, where they are heated to a semi- liquid state and then extruded onto a heated platform that is rastered to create a layer by layer build. As they cool after “printing”, the layers fuse together and solidify, creating the final part. A schematic of this preparation can be seen in

Figure1-2.

Figure 1-2. Schematic of a typical FDM setup.

FDM is also increasingly used in product development and prototyping, education88,89 and manufacturing processes which include a variety of areas, including the pharmaceuticals,90–92 design,93 biomedical engineering,94 electronics,95 robotics,96 sensing,87,97 aerospace,98 and automation.99

Even with the wide variety of applications, FDM has several limitations. A key disadvantage of FDM is that the mechanical strength of the printed parts are usually worse than those made by injection molding.100,101 This decline in mechanical strength is primarily due to the fact that once the polymer has been deposited, there are voids and inclusions that become present between the layers, causing deformations within the printed part.102 These defects cause the printed, solidified part to lose much of the mechanical strength. Another disadvantage of

FDM is that the printed polymer beads tend to have poor layer to layer adhesion, which cause the parts to fail at an accelerated rate.103 Cooling of the printed thermoplastic material can lead to shrinkage, causing warping of the printed component, which further reduces the strength of the printed material.102 Due to these disadvantages, finding strategies that can counteract and improve the thermomechanical properties for FDM printed materials remains an active area of research.

The use of PNCs in order to enhance the mechanical properties of the printed parts has been one active area of focus. The introduction of nanofillers at a low weight percentage can improve the performance of polymers, enhancing their mechanical strength, rigidity, toughness, dimensional stability,

14 thermostability, and aging resistance, for example. The inclusion of nanoparticles generally has little effect on the polymers melt fluidity and surface quality of the final printed products. Thus, development of nanocomposite materials would be an effective way to improve 3D printing processability and the overall performance of components produced by FDM. Several studies have even used nanoreinforcements in FDM 3D printed materials, such as graphene,104 vapor‐ grown carbon fibers,101 and single‐walled carbon nanotubes104 to improve the polymer properties. However, development of nanocomposites for 3D printing applications is still in a preliminary research stage.

1.3.3 Scaling of polymer grafted-nanoparticles via the grafting “through” approach

One of the major drawbacks in the adoption and use of polymer-grafted nanocomposites is the actual creation of the polymer-grafted nanoparticles. Many techniques have been developed for the covalent attachment of polymer to surfaces.21,28,31,105,106 These techniques are often divided into two main categories either -- “grafting to” or “grafting from” approaches. A schematic of these can be seen in Figure 1-3. In the “grafting to” approach (Figure 1-3a), premade functional chains are attached by reaction with groups located on the nanoparticles surface.18,107–109 The functional groups are most often located either at the chain ends, but they also can be present along the backbone of the polymer chain, or part of side chains. One of the major disadvantages of the grafting to approach is a low grafting density is achieved through this technique, mainly caused by the size of the polymer. Attached chains create a barrier to new chains diffusing to the

15 surface and attaching. In the case of the “grafting from” approach (Figure 1-3b), which also is known as “surface-initiated polymerization”, the strategy for attachment of the polymeric chains is taken from a bottom-up approach.110–114 The surface of the nanoparticle is modified with an initiating species and the polymer chain is grown from the surface. In the “grafting from” approach, higher grafting densities can be achieved; however, the amount of time for synthesis and purification needed to generate a large high amount of polymer-grafted nanoparticles is a drawback because this approach generally relies on controlled polymerization methods, which add complexity.110–114

A third approach, known as “grafting through” (shown in Figure 1-3c), is also used. In this method the surfaces have been modified with a self-assembled monolayer that contains polymerizable groups – often vinyl groups.115 The growth of the polymeric chains is initiated in solution and during propagation, a surface- bound monomer unit can be integrated into the growing chains.

Figure 1-3. A general depiction of different surface modification techniques of grafting polymers to nanoparticles: (A) “grafting to”, (B) “grafting from”, and (C)

“grafting through” methods.

The growth of the polymeric chains is initiated in solution and during propagation, a surface-bound monomer unit can be integrated into the growing chains. This results in a permanent anchoring of the polymer chains and the polymer can continue to propagate, adding more repeat groups to the polymer chain. The grafting through approach in essence is the attachment of a polymer chain by a method that contains elements of both the “grafting to” and “grafting from” approaches. A key difference is that in the grafting through approach, conventional free radical polymerization is typically used.115 Free radical polymerization creates a simple and facile way to scale the production of the polymer-grafted nanoparticles and thus will allow PNCs to be produced in a more optimized fashion.

1.4 Polyampholyte Brushes Synthesized via ARGET ATRP

Polymer brushes are a type of polymer thin film consisting of end-tethered chains anchored by one end to a surface at areal densities high enough to cause adjacent chains to overlap laterally. This overlap causes an extended conformation because stretching is how the polymer chains reduce repulsive inter-chain interactions. A specialized class of polymer brushes are polyelectrolyte brushes, which contain one type of ionizable groups along the polymer chain.

Polyelectrolyte brush systems have been comprehensively studied via experiments,116–121 theory,122–126 and simulation.127–130 These studies showed that in order for the brushes to compensate charge caused by dissociation of charged groups, there is an influx of counterions into the brush that causes osmotic repulsion. This charge-charge interaction affects the swelling, causing an inter-

17 relationship between conformation and charge. This relationship can be affected by pH, the type, and concentration of salts. These factors not only cause complex behaviors, but also prompts their use in myriad applications.

Additional complexity in polyelectrolyte phase behavior can be promoted by incorporating both cationic and anionic groups into the same polymer chain. This incorporation creates a “polyampholyte” brush. Polyampholytes can be made from pairs of monomers that are comprised of weak acids and weak bases,131–134 strong acids and strong bases,135 or a combination of strong and weak electrolytic groups.136–139 In systems in which the polyampholytes consist of repeating units that are weak electrolytes, the complete response of the polymer can be altered simply by varying the copolymer composition during synthesis. The composition of the polyampholyte brush changes the number of positive or negative repeat units and thus, the degree-of-dissociation of those repeat units can be altered by changing using solution pH or ionic strength. In other words, they affect in order to manipulate the equilibrium between charged and uncharged states.140,141 In addition, the presence of positively- and negatively-charged groups on a polyampholyte gives rise to a state in which the net charge on the chain is zero due to self-neutralization of the electrolytic groups. This self-neutralized state of the system is known as the isoelectric point.142 Self-neutralization between the oppositely charged groups ultimately leads to a collapsed conformation. At high charge asymmetry, polyampholytes will exhibit behavior characteristic of either a polyanionic or a polycationic polyelectrolyte due to an excess of charge-negative or charge-positive groups, respectively.132,143

In polymeric systems, sequence matters. This applies to both synthetic and natural polyampholytes.144–147 Studies of polyampholyte brushes, to-date, has been focused mainly on multiblock copolymers122,131,133,134,136,141,143 or strictly alternating arrangements of comonomers.148–150 For example Cao et al. and

Baratlo et al. used molecular dynamic simulations to examine how alternating arrangements of charge-positive and charge-negative comonomers affect the structure of polyampholyte brushes. They showed that for flexible polyampholytes

(PAs) having longer “portions” of alternating comonomers, the electrostatic interactions between the oppositely charged sequences near either planar or spherical surfaces tend to collapse the chains. In these alternating designs, they also showed that when the “run length” of similarly charged repeat units increases, there is a strong dependence between the interplay of overall chain stiffness and electrostatics, promoting the buckling of chains in order for the charge to be compensated.148–150 These factors lead to complex dependencies of chain conformation and brush structure on stiffness, added salt, and grafting density.

These realities directly motivate my studies of polyampholyte brushes.

1.5 Organization of Dissertation

The remainder of this dissertation is laid out as follows: the next chapter details the ARGET ATRP synthesis and characterization of poly(methyl methacrylate-random-cyclohexyl methacrylate)-grafted nanoparticles of various composition and examines the miscibility of these copolymer-grafted nanoparticles in the chemically dissimilar polystyrene matrix. Chapter 3 describes results of the

FDM studies that used the CPGNs that are presented in Chapter 2. In Chapter 4,

I expand the potential of CPGNs for FDM by presenting a scalable method to synthesize polymer-grafted nanoparticles is investigated. Grafting through is used to make poly(methyl methacrylate)-grafted nanoparticles as a simple and facile route for the creation of PNCs. Chapter 5 describes the ARGET-ATRP synthesis and swelling response of compositionally varied poly(methacrylic acid-co-N, N- diethylaminoethyl methacrylate) polyampholyte brushes. My dissertation concludes with perspectives drawn from the entire body of work (Chapter 6).

Additional supporting information is contained in the Appendices.

Chapter 2: Dispersion of Polymer-grafted Nanoparticles Poly(methyl methacrylate-r-cyclohexyl methacrylate) Synthesized via Surface-initiated

ARGET ATRP in a Chemically Dissimilar Matrix

2.1 Abstract

By incorporating enthalpic interactions between graft and matrix chains, the miscibility of a blend of copolymer-grafted nanoparticles and matrix polymers can be tuned, providing a way to control the dispersion and macroscopic properties of polymer nanocomposites. Structurally well-defined polymer-grafted nanoparticles consisting of a tethered layer of poly(methyl methacrylate-random-cyclohexyl methacrylate) (P(MMA-r-CHMA)) chains were prepared by modifying the surface of silica nanoparticles with initiators and polymerizing from the surface using activators regenerated by electron transfer atom transfer radical polymerization

(ARGET ATRP). The “pseudo-living” character of ARGET ATRP allowed P(MMA- r-CHMA) chains of low dispersity to be grown from silica nanoparticles with control over their composition and molecular weight. Copolymer-grafted nanoparticles having different P(MMA-r-CHMA) compositions were dispersed in a chemically distinct polystyrene (PS) matrix, and miscibility assessed using atomic force microscopy. These studies show that miscibility depends on graft copolymer composition due to attractive enthalpic interactions between CHMA and styrene repeat units. Specifically, incorporating at least 25% CHMA (molar basis) into the graft copolymer promotes dispersion of the copolymer-grafted nanoparticles in the

PS matrix. These studies show that the miscibility of a polymer nanocomposite can be tuned by manipulating the enthalpic interactions between the graft and the polymer matrix, which should, in turn, lead to improvements in macroscopic properties and more robust materials.

2.2 Introduction

Polymer nanocomposites (PNCs) created by adding nanoparticles (NPs) into a polymer melt frequently display significantly improved thermomechanical properties compared to the pure polymeric material.30,49–52 Experimental studies have shown that the properties of polymer-nanoparticle mixtures depend on the extent to which the nanoparticles are dispersed in the matrix polymer.151,152

Dispersion of nanoparticles is required within the matrix for PNCs to display enhanced macroscopic properties.8,20,55–57 However, many blends of polymers and nanofillers are immiscible due to strong van der Waals interactions between the fillers, which drives aggregation. The formation of clusters of nanoparticles and large aggregates has a detrimental impact upon the interfacial contact between the polymer and the filler, which leads to a substantial decrease in the bulk properties. As a result, finding strategies to overcome the tendency of nanofillers to aggregate by efficiently dispersing nanoparticles in polymer matrices remains an active area of research.

One of the most successful methods to control the dispersion of nanoparticles in a polymer matrix is to shield the particle from strong van der Waals forces and prevent attraction by depletion forces by grafting polymers onto the nanoparticles that are chemically identical to the matrix polymer.31,21,58,59,22,60,61

This leads to the formation of polymer-grafted nanocomposites (PGNCs). Many theoretical studies22,23,33,62,63 and experimental investigations19,24–26,64 of polymer- functionalized nanoparticles have shown that the size and chemical nature of the

23 grafted polymers, the geometry and size of the nanoparticles, and the chemical identity and size of the polymer matrix play crucial roles in determining the spatial organization of the nanoparticles. The extent of penetration of the matrix polymer into the grafted polymer layer and the interaction between those chains will dictate the extent of dispersion of polymer grafted particles in the free matrix polymer. If the graft and matrix polymers are chemically identical the system is considered athermal. This type of system is often referred to as a “symmetric” system. At high grafting density,19,65,66 the grafted polymers extend due to crowding, promoting a brush-like conformation of the tethered chains decorating the nanoparticle surface.

The homopolymer-grafted nanoparticles will disperse (aggregate) if the molecular weight of the matrix homopolymer is lower (higher) than that of the grafted homopolymer. This is purely an entropy effect related to mixing: Entropy increases if low molecular weight chains mix into the grafted layer, which lowers the free energy of the system. At low grafting density,18,22,23,31,67 the grafted chains may not cover the nanoparticle surface completely, leaving exposed patches of the surface.

In this situation, the fraction of surface area that is exposed will dictate the effective interparticle interactions.23,68–70 Homopolymer-grafted nanoparticles at low grafting densities have been shown to assemble into a variety of unusual arrangements in solution67,153–155 and in polymer matrices.31

Copolymer functionalization, as opposed to homopolymer functionalization, yields additional parameters that can be used to control the assembly of copolymer-grafted nanoparticles in a polymer matrix. In addition to molecular weight, comonomer type and composition affect dispersion of these “asymmetric”

24 systems of copolymer grafted nanocomposites (CPGNCs). In the last few years, theory,71–74 simulations,75–77 and experiments have been used to study copolymer- grafted nanocomposites that have a matrix that is chemically similar matrix to one of the comonomers being investigated. Vorselaars et al.73 used self-consistent field theory (SCFT) to study the dense layers of diblock copolymers grafted onto spherical nanoparticles. They found that domains of various shape and size could form on the particle surface depending on the composition of the copolymer. They examined the stability of various morphologies adopted by the grafted chains upon microphase segregation on the highly curved nanoparticle surfaces, and drew contrasts to those on surfaces having zero curvature. Zhu et al.74 used both SCFT and density functional theory (DFT) to study a dense system of nanoparticles with a single diblock copolymer chain grafted onto the particle. In these studies, they did not include matrix chains. When the particle surface was chemically neutral to the grafted chain, they saw typical block copolymer morphologies (i.e. cylinders and lamellae), which were set by both the composition of the copolymer and the size of the particle. When the particle surface was repulsive to both blocks of the copolymer, they observed ordered hierarchical structures, such as cylinders with cylinders at the interfaces, lamellae with cylinders at the interfaces, and lamellae with cylinders inside a domain. These are not typical of block copolymer melts, highlighting the potential of copolymer grafts to change nanostructure.

Not much exploration into placing copolymer-grafted nanoparticles in a chemically distinct matrix has been done this addition of copolymer-grafted nanoparticles causes the formation of another type of an asymmetric system.

While dispersion−aggregation (phase behavior) of chemically identical PGNCs with dense brushes is largely governed by entropic driving forces, there is a competition between the enthalpic and entropic driving forces in chemically dissimilar PGNCs. Theoretical studies of colloidal particles or flat surfaces grafted with a dense polymer brush placed in a chemically dissimilar polymer matrix have shown that the Flory-Huggins parameter χ between graft and matrix polymers governs the tendency of matrix chains to wet (dewet) the graft chains, which impacts phase segregation behavior.17,32–36 Experiments made with asymmetric systems also show that particle dispersion and aggregation are tuned by enthalpic interactions between grafted and matrix chains.37,78,79 Ojha et al. demonstrated this by synthesizing poly(styrene-r-acrylonitrile) (SAN) grafted nanoparticles and dispersing the copolymer-grafted nanoparticles into a chemically dissimilar poly(methyl methacrylate) (PMMA) matrix. They observed that even when the molecular weight of the grafted-SAN copolymer (Mn = 16,000) was well below the matrix molecular weight (Mn = 93,000), the nanoparticles were dispersed within the matrix.37 In this case, the attractive (enthalpic) interactions between the grafted polymers and the matrix chains, drives the system toward dispersion, even though the grafted chains are of lower molecular weight than the matrix chains. This is, of course, a case where chemically identical systems (symmetric systems) would exhibit dewetting and nanoparticle aggregation.37 This ability to disperse particles due to graft−matrix miscibility allows for higher filler loadings than chemically identical graft−matrix composites, which can lead to property improvements.37

Surface-initiated controlled free radical polymerization is often used to create a densely grafted layer of end-anchored chains on nanoparticles. Controlled polymerization methods provide a high degree of control of the size, composition, and uniformity of the grafted polymer chains.44,110,156–159 While atom transfer radical polymerization (ATRP) is often used, the form of ATRP known as activators regenerated by electron transfer atom transfer polymerization (ARGET ARTP) offers certain advantages. In ARGET ATRP reducing agents, such as ascorbic acid or tin(II) 2-ethylhexanoate, are used to generate the oxygen-sensitive Cu(I) catalyst species in situ from the higher oxidation state Cu(II) species. This feature leads to enhanced control in ARGET ATRP because of the constant regeneration of the Cu(I) activator by a reducing agent. In addition, catalyst levels are decreased to a level of parts-per-million (from parts-per-thousand in traditional ATRP), and

ARGET ATRP is tolerant to limited amounts of air. These last two conditions make this synthetic technique simpler in practice than traditional ATRP because no removal of copper is required after the polymerization and less stringent preparation conditions are used in ARGET ATRP.

In this chapter, I describe the ARGET ATRP synthesis, characterization and dispersion behavior of silica nanoparticles grafted with random copolymers of poly(methyl methacrylate-r-cyclohexyl methacrylate) (P(MMA-r-CHMA)) in a chemically dissimilar polystyrene (PS) matrix. Because polymers of CHMA are miscible with PS due to attractive interactions between CHMA and styrene,160 copolymer composition is varied to probe miscibility of copolymer-grafted nanoparticles in the chemically dissimilar PS matrix. I have discovered that the

27 dispersion of the copolymer-grafted nanocomposites depends strongly on the amount of CHMA incorporated into the grafted polymer.

2.3 Experimental Section

2.3.1 Materials and preparations

Methyl methacrylate (MMA, 99%, Aldrich) and cyclohexyl methacrylate

(CHMA) were passed through a basic alumina column to remove inhibitors.

Anisole (Aldrich, 99%) was dried by stirring over CaH2 over-night. Then, anisole was distilled and the middle fraction reserved for use. Copper (II) bromide, ethyl 2- bromoisobutyrate (EBIB, Aldrich, 97%), tris (2-pyridylmethyl) amine (TPMA,

Aldrich, 98%), tin (II) ethylhexanoate (Aldrich, 98%) and all other chemicals were purchased from Aldrich and used as received.

2.3.2 General characterization

Number-average molecular weights, Mn, and dispersities of P(MMA-r-

CHMA) copolymers were determined using gel permeation chromatography

(GPC) on a Tosoh EcoSEC GPC fitted with two Tosoh TSKgel

SuperMultiporeHZ-M columns 4μ (4.6 x 150 mm) and a TSKgel

SuperMultiporeHZ-M guard column. Measurements were made at 40 °C using tetrahydrofuran (THF) with 5% triethyl amine used as the mobile phase at a flowrate of 1 mL/min. Molecular weights were determined using the EcoSEC Data

Analysis package (version 1.04) and polystyrene and PMMA standards of narrow dispersity (Polymer Laboratories, Inc.) were used for calibration.

1H NMR spectra were obtained using a Varian Mercury 300 MHz NMR spectrometer. When following the kinetics of ARGET ATRP, aliquots of the

28 polymerization solution were taken at regular time intervals using a nitrogen- purged syringe and analyzed immediately by 1H NMR spectroscopy to determine conversion. Copolymer-grafted nanocomposite thin films were prepared on freshly-cleaned silicon wafers by spin coating at 1500 rpm for 30 seconds using a

Laurel Technologies Corporation spin coater (model WS-650MZ-23NPP).

Topographical images were acquired using an Asylum Research MFP3D atomic force microscope operating in the tapping mode using rectangular cantilevers with a spring constant of ~40 N m-1, which typically have resonance frequencies between 250-300 kHz.

2.3.3 Synthesis of poly(methyl methacrylate-r-cyclohexyl methacrylate) (P(MMA- r-CHMA)) via ARGET ATRP:

All syntheses were run with a monomer to initiator ratio, [M]:[I] = 1000:1 and fixing the total monomer concentration at 2.5 M. The copper(II) concentration was held constant at 100 ppm. As a representative example, a synthesis is detailed for a random copolymer targeted to be 90 mol% CHMA. To a dry 50 mL three-neck round bottom flask, CuBr2 (1.08 mg, 0.0045 mmol) and TPMA (5.22 mg, 0.018 mmol) were added and the head space purged with dry nitrogen. Then, using a nitrogen-purged syringe, anisole (2 mL) was added and the catalyst complex allowed to form for 30 minutes at room temperature. Once the complex had formed, which was indicated by the solution turning a yellow color, more dry anisole (6 mL) was added to the flask using a nitrogen-purged syringe. A previously prepared solution of CHMA (6.00 g, 36.6 mmol), MMA (0.370 g, 3.66 mmol), anisole (1 mL), and EBIB (8.58 mg, 0.044 mmol) was then added to the

29 flask slowly via a nitrogen-purged syringe. Then, a previously prepared solution of tin (II) ethylhexanoate (0.020 g, 0.05 mmol) dissolved in dry anisole (1 mL) was immediately added to the flask via a nitrogen-purged syringe. Once the tin (II) ethylhexanoate was added to the flask, four freeze-pump-thaw cycles were used to remove dissolved O2. The flask was then lowered into an oil bath thermostatted at 50 °C to start the polymerization. The kinetics of the polymerization was monitored by taking aliquots every hour. The polymerization was quenched after

480 min by lowering the flask into liquid nitrogen and opening the flask to air. The polymer was then purified by precipitation into cold methanol (approximately 100 mL). After isolation and purification, the recovered P(MMA-r-CHMA) polymer was analyzed by SEC and 1H NMR spectroscopy.

2.3.4 Synthesis of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane

Preparation of the initiator which can be anchored to inorganic surfaces followed protocols described by Tugulu et al.161 The reaction was performed under a nitrogen atmosphere. In a dry 50 mL round bottom flask a solution of 3- aminopropyltrimethoxysilane (1.8 mL, 0.8 mmol) and triethylamine (1.2 mL, 0.8 mmol) in 15 mL of anhydrous tetrahydrofuran (THF) was prepared. Then 2- bromoisobutyryl bromide (1 mL, 0.8 mmol) was added dropwise. The solution was stirred for 3 hr at 0 ⁰C and then for an additional 10 hr at room temperature. The byproduct triethylammonium bromide(solid) was removed by gravity filtration. The solvent was slowly removed from the filtrate by evaporation from the filtrate under reduced pressure. The purified product, 3-(2-bromoisobutyramido) propyl(trimethoxy)silane, was analyzed by 1H NMR spectroscopy (300 MHz): δ

6.87 (s, 1H), 3.53 (s, 9H), 3.22 (m, 2H), 1.91 (s, 6H), 1.62 (m, 2H), 0.62 (t, J= 8.12

Hz, 2H). These are consistent with results from Tugulu et al.161

2.3.5 Activation of nanoparticles using 3-(2-bromoisobutyramido)- propyl(trimethoxy)silane

A solution of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane (1 mL) and silica nanoparticles (4 ml, 14 nm; Nissan) in 20 mL of anhydrous THF was prepared and then added to a dry, 50 mL round bottom flask equipped with a condenser.

The solution was then placed in an oil bath that was thermostatted at 80 ⁰C and allowed to react for 24 hr. After 24 hours, the solution was removed from the heat and diluted with THF. Once diluted, the solution was separated into a minimum of

2 conical centrifuge vials. The centrifuge vials were placed in a centrifuge

(Eppendorf, Centrifuge 5702) and centrifuged at 4000 rpm for 10 min. A pellet of activated nanoparticles formed in the bottom of conical vials after 10 min. The supernatant was then removed by decanting the solvent and the remaining pellet of nanoparticles was resuspended in THF. The conical vials were again centrifuged at 4000 rpm for 10 minutes. This process was repeated a total of 5 times to ensure that the activated nanoparticles were clean and free of any excess initiator. The activated nanoparticles were then dried in a vacuum oven overnight.

Thermogravimetric analysis was used to determine the grafting density of initiator on the activated nanoparticles.

2.3.6 Synthesis of polymer grafted nanoparticles via surface-initiated ARGET

ATRP

ARGET ATRP was used to grow random copolymer chains from the initiator-functionalized nanoparticles. A typical synthesis to create a copolymer- grafted nanoparticle with a copolymer targeted to be 90 mol % CHMA (nominally) is described. A dry, 50mL three-neck round bottom flask was used for this preparation. To this round bottom flask, CuBr2 (1.05 mg, 0.0045 mmol), TPMA

(5.20 mg, 0.018 mmol), activated nanoparticles (120 mg) and anisole (2 mL) were added and the headspace purged with dry nitrogen. The solution was stirred for

30 minutes in order to allow the catalyst complex to form as well as to fully suspend the nanoparticles in the solution. After this time, CHMA (6.00 g, 36.6 mmol), MMA

(0.370 g, 3.66 mmol), EBIB (8.58 mg, 0.044 mmol) and dry anisole (1 mL) were added to the flask via a nitrogen-purged syringe. Then, a previously prepared solution of tin (II) ethylhexanoate (20 mg, 0.44 mmol) dissolved in anisole (1 mL) was added to the round bottom flask using a nitrogen-purged syringe. Once the addition was complete, the flask was lowered immediately into liquid nitrogen. The reaction mixture was degassed by four freeze-pump-thaw cycles, and after the final cycle it was placed in an oil bath set at 50 ⁰C to initiate the polymerization.

The polymerization solution was stirred at 500 rpm for 480 min. After 480 min, the polymerization was quenched by lowering the flask into liquid nitrogen and opening the flask to air. Once thawed, the solution was diluted with THF and divided into conical centrifuge vials. Centrifugation was used to isolate and collect the copolymer-grafted nanoparticles. After centrifugation, the supernatant was

32 removed by decanting. The decanted solution contains free chains because the procedure includes sacrificial EBIB initiator. These free chains were precipitated from solution in cold methanol and collected for analysis. The nanoparticles were then rinsed by adding more THF and concentrated by centrifugation. This process of using centrifugation to concentrate CPGNPs and separate and isolate free chains by precipitation of the supernatant was repeated until no polymer precipitated out of the supernatant. Typically, this required 5 cycles, which ensures that no free polymer was present with the copolymer-grafted nanoparticles

(CPGNPs). The CPGNs were then dried overnight and analyzed via thermogravimetric analysis.

2.3.7 Preparation of copolymer nanocomposites

Copolymer nanocomposites were prepared using a simple solution mixing procedure. Initially, polystyrene (PS) (Mw= 50,000 g/mol, referred to as PS50k) was dried under vacuum overnight prior to use. A 5 wt % blend of CPGNs in matrix

PS was prepared by co-dissolving the CPGNs (silica NPs grafted with P(MMA-r-

CHMA)) and PS 50k for a total of 100 mg in toluene and stirring the solution at room temperature for at least 24 hours. A nanocomposite thin film was created by spin coating onto a clean silicon wafer at 1500 rpm for 30 seconds. As cast and annealed thin films were analyzed by atomic force microscopy (AFM).

2.4 Results and Discussion

2.4.1 Kinetics of ARGET-ATRP of P(MMA-r-CHMA) in solution

Surface-initaited polymerizations from nanoparticles often suffer from a loss of control due to termination reactions, which is exacerbated by there being a low

33 number of active chains growing from small nanoparticles. In order to ensure there was no loss of control due to side reactions or termination events, solution based kinetic studies were performed. These were used to confirm that the polymerization was pseudo-living, following first order kinetics, and that comonomer incorporation tracked solution composistion. ARGET ATRP of

P(MMA-r-CHMA) in anisole was conducted either at 50 °C or 70 °C, depending on the composistion being targeted.

ATRP of MMA typically is run at 70 °C.162 When random copolymers having higher MMA content were targeted, a temperature of 70 °C was used in order to maintain control. As the target MMA content was decreased, ARGET ATRP was run at a lower temperature, 50 °C, to ensure that control was maintained. Three different copolymer compositions were created by varying the [CHMA]/[MMA] in the feed (See Table 2-1) in order to cover the full miscibility window of this random copolymer with polystyrene. The copolymers synthezied were 10/90, 75/25 and

90/10 MMA/CHMA. Previously, Paul et al. showed that p(MMA-r-CHMA) random

χ 163 copolymers were miscible with PS when fCHMA > 21.6 % because CHMA-S < 0.

Polymers of CHMA are miscible with polystrene due to attractive enthalpic

χ 160 interactions, and Paul et al. report CHMA-S= -0.03. The resulting copolymers were analyzed by 1H NMR and SEC chromatography. The kinetic plots generated were based on aliquots taken at regular intervals during the polymerization and are expressed as conversion determined from 1H NMR as a function of time. (See

Figure 2-1.)

Table 2-1. Macromolecular characteristics of P(MMA-r-CHMA) polymerized in solution.

Copolymer Molar Molar Compositional a a Composition Mn (kDa) ᴆ Ratio (Feed) MMA:CHMA MMA:CHMAb

90:10 91:9 80.4 1.10

75:25 74:26 80.1 1.10

10:90 11:89 70.6 1.11 aMolecular weight and polydispersity (ᴆ) obtained of copolymer. bAnalyzed by 1H

NMR.

Figure 2-1. Kinetic plots for ARGET ATRP of P(MMA-r-CHMA) at various comonomer ratios. For plots A and B, the copolymer was targeted to be 10% MMA

(90% CHMA); for C and D the copolymer was targeted to be 75% MMA (25%

CHMA); and E and F the copolymer was targeted to be 90% MMA (10% CHMA).

All reactions performed in anisole with A and B at 50 ⁰C while the others were done at 70 ⁰C.

The linear relationship reflected in each of the kinetic plots (Figure 2-1) indicates that the polymerizations are first order with respect to monomer conversion for all comonomer compositions studied. The slope extracted from each yields information about the effective rate of the ARGET ATRP reaction, keff.

The slope deduced from the kinetic plot for the 10/90 MMA/CHMA copolymerization is 0.00285 s-1, which is similar to but smaller than the slopes determined from polymerizations at comonomer ratios of 75/25 and 90/10, which

-1 both have keff = 0.00345 s . The overall polymerized rate is slower for 10/90 because this copolymerization was run at a lower temperature (50 ⁰C) than the other two polymerizations (which were run at 70 ⁰C). Not only is the keff for each polymerization is similar for each of the polymerizations, 1H NMR results show that copolymer composition follows solution composition. Taken together these results suggest a lack of preference of the incorporation of one monomer over the other.

The consumption of comonomers and formation of the random copolymer also is reflected in the data presented in Figure 2-2. In this stacked NMR spectra, the peak 4.85 ppm (1H) is exclusive to CHMA. The signal from this proton shifts to the right (4.65 ppm) and increases and becomes broader as the monomer is incorporated in the polymer. For MMA, the relevant peak attributed to the monomer is present at 3.76 ppm (3H), and this signal shifts to 3.55 ppm during the course of the polymerization. These peaks allow the copolymer composition and monomer concentration to be determined during the kinetic studies.

Figure 2-2. 1H NMR spectra for ARGET ATRP of P(MMA-r-CHMA) at a composition of 75/25 MMA/CHMA. An aliquot was taken every 1 hr and a total of

7 aliquots were acquired. The scheme above the spectra indicates the signature peaks for the monomers and the copolymer: A 4.85 ppm (monomer peak for

CHMA); B 3.76 ppm (monomer peak for MMA); C 4.65 ppm (CHMA in polymer), and D 3.55 (MMA in polymer).

2.4.2 Synthesis of P(MMA-r-CHMA)-grafted nanoparticles

The synthesis of tethered P(MMA-r-CHMA) chains on silica nanoparticles via a “grafting from” approach is illustrated in Scheme 2-1. The p(MMA-r-CHMA) grafted nanoparticles, ranging in composition from comonomer 90 mol % to 10 mol

% were synthesized based on conditions identified through the solution polymerization p(MMA-r-CHMA) by ARGET-ATRP. The CuBr/TPMA catalyst system, and free initiator ethyl-2-bromoisobutyrate were added to the polymerization at a monomer-to-initiator ratio, [M]:[I] = 1000:1. It has been shown that free initiator is required to provide enough Cu(I)/L activator to establish the equilibrium between active and dormant chains during the controlled (free) radical polymerization of chains from low area substrates. The addition of the sacrificial initiator allowed the free copolymer to be used to determine the relative mole percent of each monomer incorporated in the copolymer. The free copolymer is used as a representation of the surface tethered chains. The analysis confirms smooth comonomer incorporation when initiator-functionalized silica nanoparticles are present, similar to solution-only copolymerizations. An example 1H NMR spectrum is presented in Figure 2-3, where integration of peaks at 4.65 ppm

(CHMA) and 3.55 ppm (MMA) yield a composition of fCHMA = 25%, which agrees with the comonomer feed ratio of 25% CHMA. Specifically, the peaks 3.7 ppm (3H) which is distinct for MMA and 4.8 ppm (1H) is exclusive to CHMA. Characterization by GPC shows that copolymer chains recovered from solution have narrow dispersities (Đ < 1.2). These values are shown in Table 2-2 along with Mn values

(relative to PS standards) and compositional information obtained by 1H NMR

Scheme 2-1. Synthesis of P(MMA-r-CHMA) on silica nanoparticles using ARGET-

ATRP.

Table 2-2. Characteristics of copolymer-grafted nanoparticles of P(MMA-r-CHMA).

σ a a MMA:CHMA Mn (kDa) ᴆ chains/np (chains/nm2)b

0.04 90:10 30.2 1.09 21

0.04 80:20 40.2 1.07 25

0.04 75:25 40.3 1.08 28

0.04 10:90 40.5 1.08 23

aBased on measurements of recovered from the free polymer. bGrafting density determined via thermogravimetric analysis.

42 spectroscopy; furthermore, the grafting densities were calculated from the weight loss measured by thermogravimetric analysis. The dispersities and molecular weights are similar to those obtained from kinetic studies, and analysis by 1H NMR spectroscopy again shows that comonomer incorporation is smooth and follows the comonomer feed ratio.

Copolymer compositions of 10/90, 75/25, 80/20 and 90/10 MMA/CHMA were chosen in order to distinguish how the incorporation of CHMA affects the miscibility of the copolymer-grafted nanoparticles once incorporated into a polystyrene matrix. Due to the variations in the copolymer compositions the effective χ parameter will be different for each random copolymer, and CPGNPs are likely to become more miscible with polystyrene as fCHMA increases, leading to dispersion of CPGNs.

The grafting density of chains was determined from the mass loss measured by thermogravimetric analysis (TGA). Figure 2-4 shows an example, and the grafting density was calculated using equation 1:

 WnpNA  (W pgnp) (W sil)  σ   (1)  2   Mn4π r (100  Wpgnp) (100  Wsil)    in this expression Wnp is the weight loss determined for bare nanoparticles, NA is

Avogadro’s number, Mn is the number-average molecular weight of the copolymer,

Wpgnp is the weight loss measured for the polymer grafted nanoparticle, Wsil is the weight loss of the silane functionalized nanoparticles, and r is the radius of the nanoparticles, which are assumed to be spherical. The bare nanoparticles had a weight loss of 2.3% (Wnp), the silane activated nanoparticle had a weight loss of

Figure 2-4. Thermogravimetric analysis data for bare nanoparticles, nanoparticles activated with 3-(2-bromoisobutyramido)propyl(trimethoxy)silane and 10/90

MMA/CHMA copolymer-grafted nanoparticles.

9.35% (Wsil), and the 10/90 MMA/CHMA copolymer-grafted nanoparticles had a weight loss of 36.3% (Wpgnp). When comparing the weight loss of the 3-(2- bromoisobutyramido)propyl(trimethoxy)silane activated nanoparticles to that of the

10/90 MMA/CHMA copolymer there is an increase in the weight loss due to the presence of tethered chains. This indicates that surface-initiated ARGET-ATRP was successfully used to graft chains to the nanoparticles surface, creating a copolymer-grafted nanoparticle. Using equation 1 and the weight losses determined by TGA, the grafting density for the silane-activated nanoparticles was

0.82 initiators/nm2, while the tethered copolymer chains had a grafting density of

0.04 chains/ nm2. A similar result was found for all copolymer compositions as shown in Table 2-2.

2.4.3 Miscibility of copolymer-grafted nanoparticles in a chemically distinct matrix

The miscibility of the copolymer grafted nanoparticles in a chemically dissimilar polystyrene matrix was examined by imaging drop-cast thin films with

AFM. The miscibility could be tuned by varying the composition of copolymers grafts on the silica nanoparticles. While, PMMA is immiscible with PS with a Flory-

χ = 3 160,164,165 Huggins interaction parameter, S-MMA 0.18 cal/cm indicating repulsive,

χ the interactions between styrene and CHMA is weakly attractive: CHMA-S = -0.03 cal/cm3.160 The attractive interaction between CHMA and styrene can overcome the repulsion between MMA and styrene, causing P(MMA-r-CHMA) random copolymers to be miscible in PS matrices.160,163,166 Representative AFM height images of PS films containing P(MMA-r-CHMA)-grafted silica nanoparticles of

45 various compositions are shown in Figure 2-5. (The only digital modification of the images was done by applying first-order flattening using the Digital Instruments software.) In addition to imaging the as-cast films, they were annealed for various lengths of time and reimaged as indicated across the column headings of Figure

2-5. Root mean square (RMS) roughness values, β, were calculated as a way to quantify the variation in surface topography. Values of β were determined directly from the height data using equation 2:

2 ∑ (Zi-Zavg) (2) β =√ N where β is the RMS roughness, Zi is the feature height relative to the average height of the sample Zavg, and N is the total number of samples. RMS roughnesses were calculated for each sample by imaging 4 to 7 different areas of the sample.

The copolymer containing 90% CHMA had an initial value, βo of 55 nm, and as the film was annealed, the surface topography became smoother, as seen in Figure

2-5. As inset in the image β120 = 12 nm. As this film was annealed for longer times, the roughness decreased as reflected in the row of images and associated RMS roughness values: β24 = 21 nm, β48 = 17 nm, and as noted symbol β120 = 12 nm.

This pattern of behavior is not surprising because the copolymer is mainly comprised of CHMA, which is miscible with polystyrene. The system containing random copolymer grafted nanoparticles of 25% CHMA, which is close to the miscibility limit (at least fCHMA ≥ 21.6% is required for the copolymer to be miscible with the polystyrene) also shows a relative smooth as-cast topography.

Each composition was imaged as cast and then as a function of annealing time.

Samples were annealed at 150 ⁰C under vacuum for up to 120 hr.

The initial RMS roughness β0 = 44 nm, and as the film is annealed the surface became smoother, ending with a β120 = 25 nm. The CPGNs made with copolymer grafts of 90% CHMA and 25% CHMA do not show any signs of a macrophase separation such as large aggregates. The films became smoother with increasing annealing time, which confirms that these random copolymer-grafted nanoparticles were miscible with the chemically dissimilar PS matrix. The fact that these systems maintain dispersion even when the matrix molecular weight (MW) is greater than the graft molecular weight ratio is significant. In athermal systems where the graft and matrix chains are of the same chemical type, aggregation is observed when the matrix molecular weight is greater than the molecular weight.

Even though matrix molecular weight exceeds that of the grafted copolymer by at least 10 kDa, the reason I see dispersion in these two CPGNs systems (90%

CHMA and 25% CHMA grafts) is because of attractive enthalpic interactions

(negative χ) drive graft−matrix contacts. This favorable enthalpy of mixing overcomes the tendency for these systems to demix, resulting in well-dispersed nanocomposites.

Nanocomposites made with CPGNs containing 20% CHMA display significant roughness after they are annealed for a long period of time. This behavior suggests that these copolymer-grafted nanoparticles are not miscible with polystyrene. The RMS roughness for the as-cast 20% CHMA composition β0

= 200 nm, which is significantly rougher than the previous copolymer-grafted nanocomposite films containing CPGNs of either 90% CHMA or 20% CHMA. With annealing, the film topology becomes slightly smoother, as seen by β120 =180 nm.

The nanocomposite having 20% CHMA copolymer-grafted nanoparticles did not show a steady decrease in the RMS roughness as β24 ≈ β48 ≈ β120 ≈ 180 nm, suggesting that the copolymer-grafted nanoparticles were slightly immiscible with the PS matrix. This result was expected, as Paul et al. showed that copolymers containing ≥ 21.6 mol % CHMA are miscible with polystyrene,160 however, in this case the phase segregation is not as drastic. The copolymer-grafted nanocomposite containing nanoparticles with grafts that are 10% CHMA also showed phase separation and the formation of large aggregates within the PS matrix. This was exemplified by the fact that the CPGNs had an as-cast RMS roughness of 1870 nm and β120 = 2140 nm. When comparing the 20% and 10%

CHMA nanocomposites to the 90% and 25% CHMA nanocomposites, large aggregates, which are indicative of phase separation can be seen. In addition, the

90 % and 25 % CHMA nanocomposite films fully cover the substrate, but the 20% and 10 % CHMA nanocomposites dewet the surface, providing additional support to the idea that the copolymer-grafted nanoparticles are not miscible with the PS matrix. As the relative amount of CHMA is reduced, the enthalpic forces become repulsive (positive effective χ) which along with the entropic driving forces leads to demixing of graft and matrix chains, and resulting in aggregation in the nanocomposites.

2.5 Conclusions

I describe the use of ARGET ATRP to control the random copolymerization of MMA and CHMA and synthesize copolymer-grafted nanoparticles using a

“grafting from” approach. It is found that the kinetics of polymerization for all the

49 polymer compositions are similar, following first-order behavior with smooth comonomer incorporation. The composition of the copolymer graft p(MMA-r-

CHMA) was varied, and those copolymer-grafted nanoparticles were dispersed in a chemically dissimilar polystyrene matrix. AFM imaging of the nanocomposites showed that the composition of the random copolymer strongly influences dispersion in the polystyrene matrix. Systems that are miscible cover the substrate and the films tend to become smoother as annealing time increases. Systems that are immiscible show strong evidence of phase separation that result in large aggregates having irregular features and dewet the substrates. My work shows that the miscibility of a polymer nanocomposite can be tuned by using the enthalpic interactions between the graft and the polymer matrix to alter dispersion. This should, in turn, lead to improvements in macroscopic properties and more robust materials.

Chapter 3: Fused Deposition Modeling of Poly(methyl methacrylate-r- cyclohexyl methacrylate) Copolymer-grafted Nanoparticles in a Chemically

Dissimilar Matrix

3.1 Abstract

The addition of a nanofiller into a polymer matrix has been known to create more robust materials, by improving thermomechanical properties of the polymer nanocomposite, but also integrating properties of the nanofiller. By tailoring the strength of enthalpic interactions between graft and matrix chains, the miscibility copolymer-grafted nanoparticles in the polymer matrix can be tuned, providing a way to control the dispersion and macroscopic properties of the resulting polymer nanocomposite. In this chapter, surface tethered chains of poly(methyl methacrylate-random-cyclohexyl methacrylate) on silica nanoparticles, which disperse in a polystyrene matrix, are used to create nanocomposite filaments.

These are used in fused deposition modeling (FDM), and the resulting thermomechanical properties of FDM-printed parts are evaluated.

3.2 Introduction

Three-dimensional (3D) printing is an additive manufacturing (AM) technique used to create objects with complex geometry or unique structure. There are a variety of techniques used in AM manufacturing, including fused deposition modeling (FDM),80 stereolithography (SLA),81 continuous liquid interface production (CLIP),82 digital light processing (DLP),83 and selective laser sintering

(SLS)84. These techniques have been developed to form stereoscopic objects with complex architecture. FDM has become one of the most widely used 3D printing methods due to its simplicity, cost-effectiveness and accessibility.85–87 FDM printers function by using a controlled extrusion of thermoplastic filaments. FDM printing requires that the thermoplastic filaments are heated to a semi-liquid state

52 in the print head and extruded through a nozzle in a layer by layer fashion onto a heated platform. On this platform, the layers fuse together as they cool and solidify into final form. An image of the FDM printer used in this study is shown in Figure

3-1. FDM is increasingly used in product development, such as prototyping, and gaining use in manufacturing processes. FDM is being applied in a variety of areas, including education,88,89 the pharmaceutical industry,90–92 biomedical engineering,94 electronics,95 robotics,96 sensing,87,97 aerospace,98 and automation.99

Although there are a wide variety of potential uses and applications, FDM has several disadvantages that limit its application. One of the main disadvantages of FDM is the mechanical strength of the printed polymer parts is usually lower than parts made by a traditional method such as injection molding.100,101

Figure 3-1. Luzbot mini FDM printer.

In short, incomplete fusion of extruded “beads” leave voids, and give rise to weak adhesion of layers. These “defects” in the printed parts lead to a loss of the mechanical strength. Poor layer-to-layer adhesion causes the parts to fail at an accelerated rate, and uneven cooling of the printed thermoplastic materials also leads to shrinking, which causes warping of the printed component. This also reduces the strength of the printed material. Due to these disadvantages, it is important to find strategies that can overcome these issues and improve the thermomechanical properties of FDM-printed materials.

One active area of focus has been to use of polymer nanocomposites

(PNCs) to enhance the mechanical properties of the printed parts. It is well known that the introduction of nanofillers at a low weight percentage can improve the mechanical strength, rigidity, toughness, dimensional stability, thermostability, and aging resistance of polymers. The inclusion of nanoparticles improves performance, but they also tend to have a negligible effect on the melt fluidity of the polymer and surface quality of the final printed products. Thus, application of nanocomposite materials would be an effective way to improve 3D printing processability and the overall performance of components produced by

FDM. Several studies have used nanoreinforcements in FDM 3D printed materials, such as graphene, vapor‐grown carbon fibers, and single‐walled carbon nanotubes.167 However, the development of nanocomposites in 3D printing applications is still in the preliminary research stage.

3.3 Experimental Section

3.3.1 Materials and preparations

Inhibitors were removed from methyl methacrylate (MMA, 99%, Aldrich) and cyclohexyl methacrylate (CHMA) by passing the monomers through a basic alumina column. The solvent anisole (Aldrich, 99%) was dried overnight using

CaH2 and distilled. Copper (II) bromide (CuBr2), ethyl 2-bromoisobutyrate (EBIB,

Aldrich, 97%), tris(2-pyridylmethyl) amine (TPMA, Aldrich, 98%), tin (II) ethylhexanoate (Aldrich, 98%), polystyrene pellets (molecular weight 260 kDa,

Scientific Polymer Products Inc.) and all other chemicals were used as received, unless otherwise specified.

3.3.2 General characterization

Gel permeation chromatography (GPC) was used to determine number- average molecular weight, Mn, and polydispersity of poly(methyl methacrylate-r- cyclohexyl methacrylate) (P(MMA-r-CHMA)) copolymers with a Tosoh EcoSEC

GPC fitted with two Tosoh TSKgel SuperMultiporeHZ-M columns 4μ (4.6 x 150 mm) and a TSKgel SuperMultiporeHZ-M guard column. Measurements were made at 40 °C using tetrahydrofuran (THF) with 5% triethyl amine as the mobile phase at a flowrate of 1 mL/min. Molecular weights were determined using the

EcoSEC Data Analysis package (version 1.04) and polystyrene standards of narrow dispersity (Polymer Laboratories, Inc.) were used for calibration. 1H NMR spectra were obtained using a Varian Mercury 300 MHz NMR spectrometer.

Copolymer nanocomposites were extruded using a Filabot FOEX2-110 Original

EX2 Extruder. 3D printing of the filament was done using a LulzBot Mini v1.0. A

TA instruments Q-800 DMA was used to perform dynamic mechanical analysis

(DMA) of the printed copolymer nanocomposites using the dual cantilever setup.

3.3.3 Synthesis of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane

The silane initiator was prepared using protocols described by Tugulu et al.161 Air free protocols were used such that all reactions were carried out under a nitrogen atmosphere. A solution of 3-aminopropyltrimethoxysilane (1.8 mL, 0.8 mmol) and triethylamine (1.2 mL, 0.8 mmol) in 15 mL of anhydrous tetrahydrofuran

(THF) was prepared and dispensed in a dry 50 mL round bottom flask. Then, 2- bromoisobutyryl bromide (1 mL, 0.8 mmol) was added dropwise. The solution was stirred for 3 hr at 0 ⁰C and then for an additional 10 hr at room temperature.

Triethylammonium bromide was removed by gravity filtration to recover the product. Then, the solvent was slowly removed from the filtrate by evaporation under reduced pressure. The product, 3-(2-bromoisobutyramido)- propyl(trimethoxy)silane, was analyzed by 1H NMR spectroscopy (300 MHz): δ

6.87 (s, 1H), 3.53 (s,9H), 3.22 (m, 2H), 1.91 (s, 6H), 1.62 (m, 2H), 0.62 (t, J= 8.12

Hz, 2H) which are consistent with results from Tugulu et al.161

3.3.4 Immobilization of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane onto silica nanoparticles:

A dry 50 mL round bottom flask equipped with a condenser was used to prepare a solution of 3-(2-bromoisobutyramido)propyl(trimethoxy)silane (1 mL) and silica nanoparticles (4 ml, 14 nm; Nissan Chemical). 20 mL of anhydrous THF was used. The flask was then placed in an oil bath at 80 ⁰C and allowed to react for 24 hr. The solution was removed from the heat and diluted with THF. The

56 solution then was separated into a minimum of 2 conical vials. The conical vials were placed in a centrifuge (Eppendorf, Centrifuge 5702) and centrifuged at 4000 rpm for 10 min, which causes a pellet of activated nanoparticles to collect at the bottom of the tube. The supernatant was removed by decanting and the pellet of nanoparticles was resuspended in THF. The centrifugation process was repeated a total of 5 times to ensure that the activated nanoparticles were clean and free of any excess initiator. The activated nanoparticles were dried in a vacuum oven overnight. Thermogravimetric analysis was used to determine the grafting density of initiator on the activated nanoparticles.

3.3.5 Synthesis of polymer grafted nanoparticles via surface initiated ARGET

ATRP

ARGET ATRP was used to grow random copolymer chains from the initiator-functionalized nanoparticle surface. A typical synthesis to create a copolymer grafted nanoparticle with a copolymer targeted to be 90 mol % CHMA

(nominally) is described. To a dry 50 mL round bottom flask, CuBr2 (1.05 mg,

0.0045 mmol), TPMA (5.20 mg, 0.018 mmol), activated nanoparticles (120 mg) and anisole (2 mL) were added, and the headspace purged with dry nitrogen. To ensure that the catalyst complex formed and that the activated nanoparticles were fully suspended, the solution was allowed to stir at room temperature for 30 minutes. After 30 minutes, nitrogen-purged syringes were used to add, CHMA

(6.00 g, 36.6 mmol), MMA (0.370 g, 3.66 mmol), EBIB (8.58 mg, 0.044 mmol) and dry anisole (1 mL) to the flask. Then, a previously prepared solution of tin (II) ethylhexanoate (20 mg, 0.44 mmol) dissolved in anisole (1 mL) was added to the

57 round bottom flask using a nitrogen-purged syringe. Once the tin (II) ethylhexanoate solution was added, the flask was immediately lowered into liquid nitrogen. The reaction mixture was degassed by four freeze-pump-thaw cycles, and after the final cycle (thawed) it was placed in an oil bath set at 50 ⁰C to initiate the polymerization. The polymerization solution was stirred at 500 rpm for 480 min. After 480 min, the polymerization was quenched by lowering the flask into liquid nitrogen and opening the flask to air. Once thawed, the solution was diluted with THF and divided into conical vials. Centrifugation was used to isolate and collect the copolymer-grafted nanoparticles. After centrifugation, the supernatant was removed by decanting. Because the procedure includes sacrificial EBIB initiator, the decanted solution contains free chains. These free chains were precipitated from solution in cold methanol, recovered by gravity filtration and dried under vacuum. The copolymer-grafted nanoparticles were then rinsed by adding more THF and concentrated by centrifugation. This process of using centrifugation to concentrate CPGNPs and separate free chains by precipitation of the supernatant was repeated until no polymer precipitated out of the supernatant.

Typically, this required 5 rinses to ensure that no free polymer was present with the copolymer-grafted nanoparticles (CPGNPs). The CPGNs were then dried overnight and analyzed via thermogravimetric analysis.

3.3.6 Preparation of the copolymer nanocomposite filament and FDM printing

The copolymer grafted nanoparticles and PS pellets were dried in a vacuum oven overnight. Preparation of the copolymer nanocomposite used a simple mixing procedure. A 1 wt % blend of CPGNs in matrix PS was prepared by combining the

58 proper amounts of polystyrene pellets and copolymer-grafted nanoparticles in a grinder and blending. Once blended, the polymer mixture was added to a Filabot

FOEX2-110 Original EX2 Extruder and extruded at 115 ⁰C to create the nanocomposite filament. The filament was printed with a LulzBot Mini v1.0 using a printing profile (parameters) for high impact polystyrene, which is a standard profile found in the CURA software. The specimens were printed in the X Y plane with a raster angle of +45 ⁰/-45 ⁰, and the printing temperatures from 230 ⁰C to 270

⁰C were examined to elucidate optimum printing conditions. The bed temperature was set to 110 ⁰C and the inset fill density was set to 100%. The thermomechanical properties of the printed specimens were analyzed using dynamic mechanical analysis.

3.4 Results and Discussion

3.4.1 Optimizing printing conditions for pure polystyrene and copolymer-grafted nanocomposites

Polystyrene is typically used in injection molding, and because of this the printing of polystyrene using FDM has not been studied as comprehensively. Since

PS is not generally used in FDM, a variety of printing conditions were examined in order to examine the impact on thermomechanical properties and, as a result, optimize the printing process. A range of optimal printing temperatures were determined by using thermogravimetric analysis to examine degradation of PS pellets. This ensures that the temperature is below that at which degradation occurs during the printing process. As seen in the thermogram (Figure 3-2), the onset of degradation for the PS pellets begins ~300 ⁰C with full decomposition by

Figure 3-2. Thermogravimetric analysis of pure PS pellets indicates onset of degradation is at ~300 ⁰C.

450 ⁰C. Thus, printing temperatures below 300 ⁰C were deemed to be appropriate.

To ensure that the PS filament was in a molten state, the temperature window that was chosen for these studies was 230 ⁰C to 270 ⁰C, and increments of 5 ⁰ were used. Multiple DMA specimens were printed at each of these temperatures. A series of stress-strain curves generated by DMA can be seen in Figure 3-3. The

Young’s modulus for each sample at each temperature are obtained from the initial slope, and values can be found in Appendix A. Evaluating multiple samples allowed the average Young’s modulus (and standard deviation) to be determined at each temperature across the range, and those values are shown in Table 3-1.

When examining the values of Young’s modulus measured at the various printing temperatures, there were three which exhibited the highest thermomechanical properties and the greatest sample-to-sample consistency. The

PS samples printed at 270 ⁰C had a Young’s modulus of 1775 ± 282 MPa, PS at

265 ⁰C had a Young’s modulus of 1949 ± 124 MPa, and PS at 235 ⁰C had a

Young’s modulus of 1772 ± 74 MPa. The higher printing temperatures (265 and

270) appear to have higher thermomechanical properties. One possible reason for this is the higher temperature helps facilitate fusion and bonding of the printed for this is the higher temperature helps facilitate fusion and bonding of the printed layers causing a more structurally sound part.168

Table 3-1. Average Young’s Modulus of pure PS printed at various temperatures.

Temperature (⁰C) Young’s Modulus (MPa)

230 1385 ± 230

235 1772 ± 74

240 1560 ± 195

245 1298 ± 321

250 1614 ± 250

255 1381 ± 387

260 1750 ± 348

265 1949 ± 124

270 1775 ± 282

aAll samples printed were pure PS and each temperature was run in triplicate.

There were two main reasons why temperature 235 ⁰C was chosen. First, samples printed at 235 ⁰C had the most consistency (lowest standard of deviation) of all of the printed samples. The second reason was that those samples will provide insight into how the lower temperature affects the Young’s modulus of the printed polymer nanocomposites. These three temperatures produced specimens with the highest (270 ⁰C and 265 ⁰C) and most consistent (235 ⁰C) Young’s modulus and so they were chosen to print copolymer nanocomposites containing either bare silica nanoparticles or copolymer-grafted nanoparticles dispersed within the

PS matrix.

Using the conditions determined from the previous study, polymer nanocomposites consisting of 1 wt % bare nanoparticles in PS matrix were printed at 235 ⁰C, 265 ⁰C, and 270 ⁰C. The average thermomechanical properties of the nanocomposites containing bare silica nanoparticles are presented in Figure 3-4 and the average Young’s modulus are listed in Table 3-2. The remaining data for the 1 wt % bare silica nanocomposites can be found in Appendix A. The Young’s modulus for nanocomposites containing bare silica nanoparticles at 1 wt % that were printed at 235 ⁰C, 265 ⁰C, and 270 ⁰C are1422 ± 330 MPa, 1449 ± 256 MPa, and 1667 ± 207 MPa respectively. The decrease observed in the Young’s modulus between the nanocomposites containing 1 wt % bare silica nanoparticles and the printed pure polystyrene samples (see Table 3-2) could be due to dewetting of the bare silica nanoparticles by the polystyrene matrix. This dewetting causes the bare silica nanoparticles to aggregate within the PS matrix, causing a decrease in the thermomechanical properties. Although the Young’s modulus of the

Figure 3-4. Representative stress-strain curves measured by DMA used to determine the Young’s modulus of the printed nanocomposites comprised of 1 wt

% bare silica nanoparticles in a PS matrix at 235 ⁰C, 265 ⁰C, and 270 ⁰C.

Table 3-2. Average Young’s modulus of 1wt % bare nanocomposites printed at various temperatures.

Temperature (⁰C) Young’s modulus (MPa)

235 1422 ± 330

265 1449 ± 256

270 1667 ± 207 aAll samples were 1 wt % bare silica nanoparticle dispersed in a PS matrix and each temperature was run in triplicate.

65 nanocomposites containing bare silica nanoparticles appears to be lower than that of the pure PS, the Young’s moduli of nanocomposites containing bare silica nanoparticles are, within the range (statistical variation) of the printed pure PS samples.

Copolymer grafted nanocomposites consisting of 1 wt % 90:10 CHMA:MMA copolymer-grafted nanoparticles were printed at 265 ⁰C and 270 ⁰C. The copolymer-grafted nanocomposites were not printed at 235 ⁰C since printing at that temperature was problematic due to a beading effect that caused the printed layers to not be uniform. The molecular weight of the 90:10 CHMA:MMA copolymer that was grafted from the surface of the activator modified silica nanoparticle were comparable to the copolymers synthesized in the previous chapter. Specifically,

2 the Mn was 40 kDa and the grafting density was 0.04 chains/nm . A filament comprised of a 1 wt % 90:10 CHMA:MMA grafted nanoparticles in a PS matrix extruded at 215 ⁰C and samples were printed at both 265 ⁰C and 270 ⁰C. The thermomechanical properties were analyzed by DMA, and are shown in Figure 3-

5 and the average Young’s modulus are presented in Table 3-3. The Young’s modulus was 1706 ± 173 MPa at 265 ⁰C and 1633 ±111 MPa at 270 ⁰C. These values of Young’s modulus are higher than the average Young’s modulus of the nanocomposites containing 1 wt % bare silica. I hypothesize that this improvement is mainly due to the 90:10 CHMA:MMA grafted nanoparticles being more dispersed within the polystyrene matrix, which is a result of miscibility of p(MMA-r-CHMA) grafted nanoparticles in the PS matrix.

Figure 3-5. Representative stress-strain curves for the printed nanocomposites comprised of 1 wt % 90:10 CHMA:MMA grafted nanoparticles in a PS matrix at

265 ⁰C and 270 ⁰C.

Table 3-3. Average values of Young’s modulus of 1 wt % 90:10 CHMA:MMA copolymer nanocomposites printed at various temperatures.

Temperature (⁰C) Young’s modulus (MPa)

265 1706 ± 173

270 1633 ± 111 aAll samples were 1 wt % 90:10 CHMA:MMA grafted nanoparticles dispersed in a

PS matrix and multiple samples were run at the temperatures chosen.

However, when comparing the performance of the 1 wt % copolymer-grafted nanocomposites with those made of pure PS, the thermomechanical properties are similar, if not slightly worse.

3.5 Conclusions

This work focused on reinforcing PS 3D printed materials by adding bare silica nanoparticles or copolymer-grafted nanoparticles forming polymer nanocomposites. PS is not a widely used material in FDM 3D printing, so optimal printing conditions were determined in my initial studies. Overall, the study showed that the parts printed at higher temperatures exhibited higher thermomechanical properties overall. The higher temperatures help facilitate adhesion of the layers of the printed parts. The results of these studies need more work to clearly identify behaviors to examine dispersion in filaments, since the thermomechanical properties of the pure PS, the 1 wt % bare nanocomposites and the 1 wt % 90:10

CHMA:MMA were within standard deviation of one another. Further investigation will need to be performed to verify if tuning the miscibility of copolymer grafted nanocomposites can in fact lead to better macroscopic properties and more robust materials.

Chapter 4: Scalable Preparation of Poly(methyl methacrylate)-grafted

Nanoparticles via the Grafting “Through” Approach

4.1 Abstract

The scalable preparation of polymer grafted nanoparticles by using the

“grafting through” surface technique involves less stringent preparation that results in a facile way to synthesize polymer-grafted nanoparticles. In this chapter, well- defined polymer-grafted nanoparticles having a tethered layer of poly(methyl methacrylate) (PMMA) chains were prepared by modifying the surface of silica nanoparticles with two different vinyl silanes. These modified nanoparticles were then used and poly(methyl methacrylate) was polymerized using conventional free radical polymerization to produce polymer grafted nanoparticles on the gram scale, rather than on the milligram scale. Grafting densities of the polymer grafted nanoparticles produced using this grafting through approach were comparable to those made by the synthetic method known as “grafting to”. FTIR-ATR and DLS were used to confirm that PMMA was grafted onto the activated silica nanoparticles. These studies show that polymer grafted nanoparticles can be produced on a larger scale, which may facilitate their application as polymer nanocomposites on an industrial scale.

4.2 Introduction

The inclusion of polymer-grafted nanoparticles in a polymer matrix has been shown to improve the thermomechanical properties of the resulting polymer- grafted nanocomposites (PGNCs) relative to the unmodified, pure polymer system without nanofillers. One of the major drawbacks in the preparation of polymer- grafted nanocomposites is intensity required in the synthesis of the polymer- grafted nanoparticles. Though many techniques have been developed for the

70 covalent attachment of polymers to surfaces, none have been used to synthesize polymer-grafted nanoparticles on a larger scale. These techniques are often separated into two main categories--either “grafting to” or “grafting from” approaches. A schematic of the methods can be seen in Figure 4-1. In the “grafting to” approach (Figure 4-1a), presynthesized polymer chains having desired molecular weight, composition, and functionality are reactively attached to the nanoparticle surface.18,107–109 These functional groups that anchor the chain to the surface can be located either at the chain ends of the polymer, along the backbone of the polymer chain, or as a part of the side chains of the polymer. One of the major disadvantages to the “grafting to” approach is the low grafting density resulting from this technique. The low grafting density is mainly due to fact that polymer chains already anchored to the surface cause difficulties with incoming polymer chains diffusing to the surface and attaching. In the case of “grafting from”

(Figure 4-1b), which is also known as “surface-initiated polymerization”, the strategy for attachment of the polymer chains is taken from a bottom-up approach.110–114 In this case, the substrate or surface of the nanoparticle is modified with an initiating species. Then, the polymer chains are grown directly from these surface-modified substrates or nanoparticles. Higher grafting densities can be achieved using this method, but a major drawback remains the amount of time and rigorous purifications needed to acquire large amounts of polymer-grafted nanoparticles. These approaches almost always rely on controlled polymerization methods, which are required either to prevent termination of the growing chains

71 from the surface or, perhaps more importantly, to provide the chains with proper end groups for subsequent surface grafting.110–114

A third approach known as “grafting through” (Figure 4-1c) can be used. In this approach, the surfaces are modified to carry a self-assembled monolayer that contains polymerizable groups.115 When the growth of the polymeric chains is initiated in solution, during propagation, a surface-bound monomer unit can be integrated into the growing chains, which anchors the polymer chain to the surface.

The polymer can continue to propagate through the surface to enchain other free monomers, which adds repeat units to the tethered polymer chain. In essence, the grafting through approach attaches a polymer chain by using elements of both the “grafting to” and “grafting from” techniques. The key difference is that in the grafting through approach, conventional free radical polymerization is usually employed. This simpler polymerization method creates a facile way to scale-up the preparation of the polymer-grafted nanoparticles.

Figure 4-1. A general depiction of different surface modification techniques used to graft polymers to nanoparticles: (A) “grafting to”, (B) “grafting from”, and (C)

“grafting through” may be used for the attachment of polymer chains to surfaces

This may lead to the efficient production of PGNCs, which may widen their utility in industrial applications.

In this chapter, I describe the preparation of silica nanoparticles grafted with polymers of poly(methyl methacrylate) (PMMA) using the grafting through method, as well as characterization of the polymer grafted nanoparticles. I have discovered that the PMMA-grafted nanoparticles can be produced easily on the gram scale in a much shorter period of time (compared to the grafting from technique).

4.3 Experimental Section

4.3.1 Materials

The monomer, methyl methacrylate (MMA, 99%, Aldrich), was passed through a basic alumina column to remove any inhibitors present. Tetrahydrofuran

(THF, 99%, Aldrich) was dried by passing through an activated alumina column.169

The thermal initiator azobisisobutyronitrile (AIBN, 98%, Aldrich) was recrystallized from methanol to ensure that impurities are removed. 3-(Trimethoxysilyl)propyl methacrylate (SPM, 98%, ACROS), vinyltrimethoxysilane (VTS, 98%, ACROS) and all other chemicals were purchased from Aldrich and used as received.

4.3.2 General characterization techniques

Number-average molecular weights, Mn, and dispersities of PMMA polymers were determined using gel permeation chromatography (GPC) on a

Tosoh EcoSEC GPC fitted with two Tosoh TSKgel SuperMultiporeHZ-M columns

4μ (4.6 x 150 mm) and a TSKgel SuperMultiporeHZ-M guard column.

Measurements were made at 40 °C using tetrahydrofuran (THF) with 5% triethyl amine as the mobile phase at a flowrate of 1 mL/min. Molecular weights were

73 determined using the EcoSEC Data Analysis package (Version 1.04) and polystyrene standards of narrow dispersity (Polymer Laboratories, Inc.) were used for calibration. 1H NMR spectra were obtained using a Varian Mercury 300 MHz

NMR spectrometer. Thermogravimetric analysis (TGA) was performed using a TA

Instruments Q50, which determined mass loss during a temperature ramp from 0

⁰C to 800 ⁰C at 20 ⁰C per minute. Grafting density of chains was calculated based on the total weight loss and surface area of nanoparticles. Fourier-transform infrared attenuated total reflection (FTIR-ATR) spectroscopy was performed using a Nicolet iS50 FT-IR spectrometer which measures the absorption spectrum in the mid-IR region (5000-400 cm-1). Dynamic light scattering (DLS) measurements were performed as described previously.170–172 These measurements were taken using an ALV goniometer system operating at a wavelength of 632.8 nm with a linearly polarized 22 mW HeNe laser. At each angle a counting time of 60 s was used at 16 angles ranging from 20° to 146° this was done to obtain reliable statistics for the light intensity autocorrelation function. The hydrodynamic radius

(Rh) was then calculated giving an apparent hydrodynamic radius that was determined at a concentration of 0.5 mg/mL; though, for convenience I will refer to them simply as the hydrodynamic radius.

4.3.3 Immobilization and activation of nanoparticles using 3-(trimethoxysilyl)propyl methacrylate

A solution of 3-(trimethoxysilyl)propyl methacrylate (MPS) (0.346 g, 8.4 mmol), silica nanoparticles (8 ml, 14 nm; Nissan), and triethylamine (1.6 mL, 75 mmol) was prepared in 30 mL of anhydrous THF. This solution was then added to

74 a dry, 100 mL round bottom flask equipped with a condenser. The flask was then placed in an oil bath that was thermostatted at 80 ⁰C and allowed to react for 3 hr.

After 3 hrs, the flask was removed from the bath and the solution cooled to room temperature. Once cooled it was diluted with THF (approximately 50 mL). After dilution, the solution was separated into a minimum of 2 conical vials. The conical vials were then placed in a centrifuge (Eppendorf, Centrifuge 5702) and centrifuged at 4000 rpm for 10 min. After 10 min, a pellet of activated nanoparticles formed in the bottom of the conical vials. The supernatant of the solution was decanted and the remaining pellet of activated nanoparticles was resuspended in

THF. The conical vials were then centrifuged a second time at 4000 rpm for 10 minutes. This process was repeated a total of 5 times to ensure that the activated nanoparticles were free of any excess activator. The activated nanoparticles were then dried in a vacuum oven overnight. Once dried, the grafting density of the activated nanoparticles was determined through the use of TGA. This procedure was also used to make VTS modified nanoparticles.

4.3.4 Synthesis of poly(methyl methacrylate)-grafted nanoparticles via conventional free radical polymerization

Conventional free radical polymerization was used to grow poly(methyl methacrylate) (PMMA) chains from the SPM-functionalized nanoparticle surface.

A typical synthesis to create polymer-grafted nanoparticles is described. A dry 250 mL round bottom flask was used for this preparation. To this round bottom flask,

AIBN (32.5 mg, 0.2 mmol), activated nanoparticles (1.0 g), either VTS or SPM,

MMA (28.0 g, 0.58 mol) and dry THF (65 mL) were added. The solution was then

75 sparged with argon for 30 min while stirring to ensure that there was no oxygen remaining in the solution and to fully suspend the nanoparticles in the solution.

After this, the solution was lowered into an oil bath set to 80 ⁰C to initiate the polymerization. The solution was stirred at 500 rpm for 480 min. Then after 480 min, the polymerization was quenched by submerging the flask into liquid nitrogen after 480 min and the flask was opened to air. When thawed, the solution was diluted with THF and separated into several conical vials. Centrifugation was used to isolate and collect the polymer-grafted nanoparticles in a similar fashion to the process described in section 4.3.2 for the activated nanoparticles. The main difference is that the solution contains free polymer chains because AIBN is used as the initiation source. The free chains were separated from the pellet of the polymer-grafted nanoparticles by decanting and recovered by precipitating from the decanted solution into cold methanol. The polymer-grafted nanoparticles were then rinsed and resuspended by adding more THF, and then concentrated by centrifugation. This process of using centrifugation to concentrate PGNPs and separate free chains by precipitation of the supernatant was repeated until no polymer precipitated out of the supernatant. Typically, this required 5 cycles, which ensures that no free polymer was present with the polymer-grafted nanoparticles

(PGNPs). The PGNPs were then dried overnight and analyzed via TGA, FTIR-

ATR, and DLS.

4.4 Results and Discussion

4.4.1 Synthesis of PMMA-grafted nanoparticles by using SPM and VTS as activators

The synthesis of tethered PMMA chains on silica nanoparticles via a

“grafting through” approach is illustrated in Scheme 4-1. The PMMA-grafted nanoparticles were synthesized using two different vinyl silane-activated nanoparticles, VTS and SPM nanoparticles. The main difference between the two activators is the size of overall molecules and the stability of the radical that is formed during propagation. Although, both activators contain vinyl groups, VTS is significantly smaller than SPM, as seen in Scheme 4-1, and will produce a less stable secondary radical due to the lack of any other functionality. Whereas SPM contains both a vinyl group and a methacrylate group, so once a radical is formed a more stable tertiary radical will be present during propagation. Thus the comparison between the two silanes provides insight into how the design of the anchoring group affects the propagation of polymer through the groups on surface of the silica nanoparticles. As noted previously AIBN is added to the solution to initiate polymerization. Characterization by GPC shows that the free polymer chains have typical dispersities (Đ >1.8). These values along with Mn values

(relative to PS standards) are shown in Table 4-1. The grafting densities were calculated using thermogravimetric analysis.

Scheme 4-1. Synthesis of PMMA grafted-nanoparticles from activated nanoparticles using grafting through process with a conventional free radical initiator.

Table 4-1. Characteristics of PMMA-grafted nanoparticles.

σ a a Activator Mn (g/mol) Đ # chains/NP (chains/nm2)b

90,000 1.8 21 0.031 SPM 93,000 1.5 19 0.030

90,000 1.8 3 0.010 VTS 92,000 1.5 4 0.011

aBased on measurements of the recovered free polymer. bGrafting density calculated using weight loss measured by thermogravimetric analysis.

Figure 4-2 shows a stacked thermogram comparing both SPM and VTS activated nanoparticles and the resulting PMMA-grafted nanoparticles, the grafting density was calculated using equation 1:

 WnpNA  (W pgnp) (W sgnp)  σ   (1)  2   Mn4π r (100  Wpgnp) (100  Wsgnp)   

In this expression, Wnp is the weight loss of a bare nanoparticle, NA is Avogadro’s number, Mn is the number-average molecular weight of the free polymer chains,

Wpgnp is the weight loss of the polymer grafted nanoparticle, Wsgnp is the weight loss of the silane-modified nanoparticle, and r is the radius of the nanoparticle, which is assumed to be spherical. Based on the thermograms, the bare nanoparticles had a weight loss of 2.3%, the SPM-activated nanoparticle showed a weight loss of 5.20%, and the VTS-activated nanoparticles had a weight loss of

2.93%. Using equation 1 and the weight loss, the grafting density for the SPM activated nanoparticles was calculated to be 0.37 silanes/nm2 and 0.10 silanes/nm2 for the VTS activated nanoparticles. The PMMA-grafted nanoparticles created by grafting through using SPM-activated nanoparticles had a weight loss of 43.7%, and the PMMA-grafted nanoparticles produced by grafting through VTS- activated nanoparticles had a weight loss of 15.2%. When comparing the SPM- activated nanoparticles to the PMMA-grafted nanoparticles made using those

SPM-activated nanoparticles, there is an increase in the weight loss by over 35%

(due to the loss of PMMA chains). On the other hand, comparing the VTS-activated nanoparticles to the PMMA-grafted nanoparticles that were produced by VTS- activated nanoparticles shows that there was an increase in weight loss as well,

Figure 4-2. Thermogravimetric analysis of bare nanoparticles, activated nanoparticles decorated with (A) SPM and (B) VTS, and the resulting PMMA- grafted nanoparticles made by grafting through.

but the weight loss of the PMMA-grafted nanoparticles was only 12% more. While weight losses indicate that I grafted polymer onto the nanoparticles, there are significantly fewer chains grafted to the VTS-modified nanoparticles, assuming that the molecular weight of grafted chains are the same as the free polymer. The

PMMA-grafted nanoparticles synthesized using VTS have a grafting density of

0.01 chains/nm2, while the PMMA-grafted nanoparticles synthesized with SPM have grafting densities of 0.03 chains/nm2, these low grafting densities are comparable to grafting densities achieved via grafting to.173 The difference between the grafting densities illustrates that SPM was more effective during propagation, which is directly related to the stability of the radical forming during propagation. In which the formation of the VTS radical is significantly less stable and therefore more difficult to form versus the formation of the more stable SPM tertiary radical.

Another way to verify the presence of chains grafted to the nanoparticles was through analysis using FTIR-ATR. Figure 4-3 shows spectra acquired for bare nanoparticle, PMMA-grafted nanoparticles, and pure PMMA made using the SPM- modified nanoparticles. The vibrational stretching at 1050 cm-1 is indicative of Si-

O bonds in the bare silica spectrum.174,175 This vibrational mode is the only prominent peak in the bare silica nanoparticle spectrum illustrating that primarily

Si-O bonds are present and that the nanoparticles have not been modified. The spectrum of pure PMMA displays FTIR peaks at 3000 cm-1, 2950 cm-1, 1725 cm-1, and 1100 cm-1, which are assigned to C-H stretches, C=O saturated aliphatic ester

Figure 4-3. FTIR-ATR spectra of bare nanoparticles (blue), PMMA grafted nanoparticles (red), and pure PMMA (black). The spectra are offset vertically

(stacked) for clarity.

stretching, and C-O-C bond vibrational stretching modes, respectively.174,175 The peak at 1100 cm-1 in pure PMMA represents C-O-C stretching and is located in a similar location to the Si-O stretching present in the bare silica nanoparticle spectrum. The FTIR-ATR spectrum acquired for PMMA- grafted nanoparticles displays very weak peaks around 3000 cm-1 indicating that C-H stretching is present. There also is a prominent aliphatic ester stretching mode at 1725 cm-1, and a peak at 1075 cm-1 that is assigned to both Si-O and C-O and shifted due to a combination of both types of stretching being present in the PMMA-grafted nanoparticles. The appearance of these peaks at 1725 cm-1 and 1075 cm-1 suggests that PMMA was successfully grafted onto the SPM-activated nanoparticles is made.

DLS measurements were made using THF solutions of the bare nanoparticles, SPM activated nanoparticles, and PMMA-grafted nanoparticles at a concentration of 0.5 mg/mL. The distributions of hydrodynamic size, Rh, can be seen in Figure 4-4. The hydrodynamic radius of bare nanoparticles ~45 nm. The

SPM-activated nanoparticles had and average size of ~48 nm, and the PMMA grafted nanoparticles had an average size of ~233 nm. The hydrodynamic radius for the bare nanoparticles suggests that bare silica nanoparticles, which are nominally 14 nm in diameter, have formed small aggregates. Another indication of aggregation is the width of the distribution, which is large and ranges from ~7 nm to ~90 nm for the bare silica nanoparticles. The hydrodynamic radius increases by

~3 nm once the bare nanoparticles were activated with SPM, though it again seems likely that aggregates are present. The size and the distributions are similar

A) B)

Figure 4-4. Distribution of hydrodynamic radius, Rh, measured at four different angles for (A) bare nanoparticles, (B) SPM-activated nanoparticles, and (C)

PMMA-grafted nanoparticles in THF at a concentration of 0.5 mg/mL.

to those of the bare nanoparticles. This issue of aggregation could be examined by lowering the concentration by an order of magnitude (to 0.05 mg/mL). If aggregation is still present, the solution could be heated to help disperse the nanoparticles. The hydrodynamic radius increased significantly from the activated nanoparticles to the PMMA-grafted nanoparticles, though once again aggregation seems may be present since the range of hydrodynamic sizes is from 30 to 250 nm. The idea that the hydrodynamic radius for the PMMA-grafted nanoparticles being ~230 nm is far too large for 90 kDa PMMA to be grafted onto 14 nm silica.

Nevertheless, the increase in the Rh qualitatively confirms that PMMA chains are grafted on the nanoparticles.

4.5 Conclusions

I describe the use of the “grafting through” technique to successfully scale the synthesis of PMMA-grafted nanoparticles made using conventional free radical polymerization methods. The PMMA grafted nanoparticles were successfully synthesized on the gram scale, rather than on the milligram scale. The grafting densities for the polymer grafted nanoparticles were comparable to those made using the grafting to technique. PMMA was present on the nanoparticles surface, as evidenced by FTIR-ATR and DLS measurements. The hydrodynamic radius of the nanoparticles also increased from activated to polymer-grafted, and the size change also confirms that the silica nanoparticles contained a layer of grafted chains. My work shows that polymer-grafted nanoparticles can be produced on a much larger scale, facilitating production of polymer nanocomposites.

Chapter 5: ARGET-ATRP Synthesis and Swelling Response of

Compositionally Varied Poly(methacrylic acid-co-N, N-diethylaminoethyl

methacrylate) Polyampholyte Brushes

This chapter describes work submitted to Soft Matter 2018. I synthesized and characterized all of the random copolymers discussed in this work. Other coauthors include Dr. Jeremiah Woodcock, who provided the basis for the work, and Prof. S. Michael Kilbey II, advised this work. This work was supported by the

National Science Foundation (Award Numbers CBET 1133320 and 1512221).

5.1 Abstract

Modifying the composition of polyampholytes, which are comprised of charge-positive and charge-negative repeat units, directly contributes to trade-offs between charge and structure, which are externally regulated by solution pH and added salt. Here, the relative ratio of anionic and cationic comonomers is varied to tailor the stimuli-responsiveness of poly(methacrylic acid-co-N,N- diethylaminoethyl methacrylate) (P(MAA-co-DEAEMA)) brushes to changes in solution pH and an added zwitterion. These systems display a strong dependence on excess repeating units of either type and the random incorporation appears to facilitate self-neutralization of charges. Pseudo-living growth with smooth comonomer incorporation is achieved using activators regenerated by electron transfer atom transfer radical polymerization (ARGET ATRP), creating well- defined brushes. In situ ellipsometry measurements of solvated brush thickness indicate that at low and high pH, the brushes display polyelectrolyte behavior with a strong compositional dependence, but at intermediate pH values, the brushes show the characteristic collapse attributed to self-neutralization of polyampholytes.

The polyampholyte brushes maintain these patterns of behavior across all compositions and in the presence of an added zwitterion, which contributes

88 additional hydrophobic character as evidenced by decreases in the swollen layer thicknesses. The response of the P(MAA-co-DEAEMA) brushes to the organic osmolyte betaine is consistent with its tendency to stabilize proteins and peptides in a chaotropic fashion. These studies add perspective to efforts to manipulate sequence in polyampholytic polymers, support broader efforts to tailor interfacial soft films for applications in biotechnology and sensing, and understand aggregation and stability of biological polymers.

5.2 Introduction

Polymer brushes are assemblies of end-tethered chains that are anchored by one end to a surface or interface at densities high enough that adjacent chains overlap laterally. This crowding causes chains to adopt a stretched conformation in order to reduce repulsive inter-chain interactions. Polyelectrolyte brushes are an important class of polymer brushes containing ionizable groups along the polymer chain, and they have been extensively studied due to their application in colloidal stabilization,176 lubrication,177,178 and surface protection,179 as well as in biotechnology180 and the development of “smart” interfacial layers that respond to one or more stimuli181 or have utility as biomedical constructs.180 Polyelectrolyte brush systems have been extensively examined by theory,122–126 simulation,127–130 and experiments,116–121 and from these studies it is appreciated that the dissociation of charged groups and influx of counterions into the brush in order to compensate charge causes intra- and inter-chain repulsion that modulates swelling. This inter-relationship between conformation and charge, which is affected by pH and the type and concentration of salts or other electrolytes, not

89 only leads to complex behaviors but also engenders their use in a variety of applications.

Additional complexity in polyelectrolyte phase behavior can be created by incorporating both cationic and anionic groups along the same polymer chain, creating what are termed polyampholytes. Polyampholytes are made from pairs of monomers that are weak acids and weak bases,131–134 strong acids and strong bases,135 or a combination of strong and weak electrolytic groups. 136–139 In the case of polyampholytes consisting of repeating units that are weak electrolytes, the overall response can be altered by varying the copolymer composition during synthesis and by changing the degree-of-dissociation of the repeat units by using solution pH or ionic strength to manipulate the equilibrium between charged and uncharged states.140,141 The ability to alter the number of positively- and negatively- charged groups on a polyampholyte gives rise to the situation where the net charge on the chain is zero, which is known as the isoelectric point.142 At and around the isoelectric point, pairing between the oppositely charged groups in the same or neighboring chains leads to a collapse. At high charge asymmetry, either far above or below the isoelectric point, polyampholytes will behave like a polyanion or a polycation polyelectrolyte due to an excess of charge-negative or charge-positive groups, respectively.132,143

Sequence matters in polymer systems, and this is also true in both synthetic and natural polyampholytes.144–147 In the area of polyampholyte brushes, most of the focus to-date has been on multiblock122,131,133,134,136,141,143 or strictly alternating arrangements of comonomers.148–150 For example, through molecular dynamics

90 simulations, Cao et al.148 and Baratlo et al.149,150 examined how alternating arrangements of charge-positive and charge-negative comonomers affect the structure of polyampholyte brushes. Their anionic and cationic groups were weak electrolytes that the equilibrium between neutral and charged states is regulated by pH. Specifically in their work they varied the number of cationic and anionic repeat units arranged in alternating fashion from one to a few monomers in length.

They showed that for flexible polyampholytes (PAs) having longer “runs” of comonomers arranged in an alternating sequence, electrostatic interactions between oppositely charged sequences near either planar or spherical surfaces tend to collapse the chains. In these alternating designs, there also is a strong dependence on the interplay between chain stiffness and electrostatics that promotes buckling of chains in order to compensate charge when the “run length” of similarly charged repeat units increases.148–150 These factors lead to complex dependencies of chain conformation and brush structure on stiffness, added salt, and grafting density.

The responsive behavior and conformational properties of weak polyampholyte diblock copolymer brushes have been theoretically and experimentally studied by Genzer et al.134 and by Linse and coworkers.122 Genzer and co-workers synthesized poly((dimethyl aminoethyl methacrylate)-block- poly(acrylic acid)) (DMAEMA-b-PAA) brushes and used ellipsometry to study their swelling response to changes in pH. Trends in ellipsometric angles Ψ and Δ were used to interpret the total thickness of the wet brush and to produce a description of the structural changes. At high and low pH, the brushes were swollen while at

91 the intermediate pH values, they observed an increase in Δ (or Ψ), which indicates a decrease in the thickness of the brush. They reasoned that the collapse was due to pairing of charge-positive and charge-negative repeat units either within a chain

(self-collapsing) or between neighboring chains.134 This behavior is consistent with earlier results by Linse and co-workers, who used mean-field lattice theory to study the structure of diblock polyampholyte brushes grafted on a planar substrate (as well as on a spherical substrate).122 They showed that chain conformation depended on the relative charge of the blocks, which is controlled by solution pH.

When the polyampholyte brush is at a fairly low charge ratio, two types of chain conformations exist simultaneously: blocks are either stretched or coiled

(collapsed). At a relatively high charge ratio they found that a hairpin bend at the junction between the two blocks exists, driven by electrostatic interactions between the oppositely charged blocks.122 This allows the diblock chains to self-neutralize charge, which collapses the layer due to the hydrophobic effect. This characteristic pattering of swelling behavior of polyampholyte multiblock brushes was also observed by Brittain and coworkers,131,136 who used a surface-initiated atom transfer radical polymerization (SI-ATRP) (typically referred to as a ‘‘grafting from’’ approach) to synthesize poly(acrylic acid-b-2-vinyl pyridine) and poly(acrylic acid- b-4-vinyl pyridine) diblock copolymer brushes. Layer thicknesses determined by ellipsometry showed that the chains of the polyampholyte brush were stretched at low and high pH values but collapsed at intermediate pH values. Yu et al. used

AFM measurements to deduce that the collapses of polyampholytic chains around the isoelectric point causes a roughening of the surface.133

Beyond these examples of multiblock polyampholyte brushes, Tran and coworkers examined the behavior of random polyampholyte brushes comprised of a stoichiometric amount of 2-(dimethylamino)ethylmethacrylate (DMAEMA) and methacrylic acid (MAA).137,138 These polyampholytes were synthesized via surface-initiated ATRP and used to study how variations in the grafting density of chains affected the swelling behavior of the polyampholyte brushes. Neutron reflectometry was used to determine monomer volume fraction profile as a function of pH. From these measurements, they determined at pH = 2, for the two different grafting densities studied (0.193 and 0.132 chains nm-2), the chains are stretched and the polyampholyte brush acts as a charge-positive polyelectrolyte brush. On the other hand, over the pH range of low net charge, which is near the isoelectric point, the chains are collapsed, with the less-dense brush being significantly more collapsed, likely due to a greater ability for ion-pairing between chains.

Within these works, most efforts to synthesize polyampholyte brushes have relied on surface-initiated ATRP with free chains polymerized in solution (through the addition of sacrificial initiator) being recovered and used as proxies to describe the surface-tethered chains of the brush. Though standard ATRP is well known for its ability to synthesize well-defined polymers and copolymers of specific composition and low polydispersity,182–184 the variant known as ARGET-ATRP presents certain advantages for controlled polymerization.

Activators regenerated by electron transfer (ARGET) ARTP is a derivative form of ATRP in which reducing agents, such as ascorbic acid or tin(II) 2- ethylhexanoate is used to generate the oxygen-sensitive Cu(I) catalyst species in

93 situ from a higher oxidation state Cu(II) species. Cu(I) activates the polymerization while Cu(II) reversibly deactivates growing polymer chains and through these reactions, an equilibrium that favors “dormant” chains and minimizes irreversible termination reactions is established.111,185,186 The use of a reducing agent allows for less stringent conditions to initiate the polymerization, improves tolerance to oxygen, and in comparison to traditional ATRP, ARGET ATRP uses significantly lower concentrations of copper catalyst (down to ppm levels). ARGET ATRP also offers the advantage (compared to traditional ATRP) of significantly reducing side reactions of the metal/ligand catalyst complex. This feature allows ARGET ATRP to reach higher levels of monomer conversion and improve crossover by retaining chain end functionality, thereby enabling copolymers of higher molecular weight to be prepared.187–189 As the concentration of the copper catalyst is lower, the likelihood of side reactions and termination events is reduced, which enhances control over the polymerization.111,185,190,191 Together, these features make ARGET

ATRP an attractive method for preparing polyampholyte brushes, especially brushes grown from surfaces where there is a low number of active chain ends due to surface densities of chains that are typically of the order of 1015 or 10 16 ch m−2. ARGET ATRP also provides control in polymerizations involving amine- containing monomers, which has advantages in copolymerizations, allowing for control over the polymers overall compositions and smooth incorporation of comonomers.192–194

With these as motivation, we demonstrate the utility of ARGET ATRP for the synthesis of poly(methacrylic acid-co-diethylaminoethyl methacrylate) P(MAA-

94 co-DEAEMA) across a range of copolymer compositions with excellent control.

Moreover, kinetic studies verify the solution polymerizations follow first-order kinetics. This approach is then translated to produce the corresponding polyampholytic brushes having tailored ratios of charge-positive and charge- negative weak electrolytic repeat units. ARGET ATRP was chosen for its tolerance to amine containing monomers, and charge-negative MAA repeating units were obtained by hydrolysis of the tert-butyl ester of tert-butyl methacrylate. Finally, a series of swelling studies were performed to examine how copolymer composition affects structural response of random polyampholyte brushes of various composition as pH is changed and betaine a zwitterion is added.

5.3 Experimental Section

5.3.1 Materials

tert-Butyl methacrylate (tBMA, 98%, Aldrich) and N,N-Diethylaminoethyl methacrylate (DEAEMA, 99%, Aldrich) were passed through a basic alumina column to remove inhibitor. Dry anisole (Aldrich, 99%) was obtained by distillation after stirring over CaH2 overnight. All other chemicals, including copper (II) bromide, ethyl 2-bromoisobutyrate (EBIB, 97%), tris (2-pyridylmethyl) amine

(TPMA, 98%), tin (II) ethylhexanoate (Sn(II)Oct, 98%), anhydrous tetrahydrofuran

(THF, 99%), and betaine (98%)) were purchased from Aldrich and used as received.

5.3.2 General characterization methods

Number-average molecular weights, Mn, and dispersities of poly(tert-butyl methacrylate-co-diethylaminoethyl methacrylate) P(tBMA-co-DEAEMA)

95 copolymers were determined by gel permeation chromatography (GPC) on a

Tosoh EcoSEC GPC fitted with two Tosoh TSKgel SuperMultiporeHZ-M columns

4μ (4.6 x 150 mm) and a TSKgel SuperMultiporeHZ-M guard column.

Measurements were made at 40 °C using tetrahydrofuran (THF) with 5% triethyl amine as the mobile phase at a flowrate of 1 mL/min. The EcoSEC Data Analysis package (version 1.04) and conventional calibrations using polystyrene and poly(methyl methacrylate) standards of narrow dispersity (Polymer Laboratories,

Inc.) were used to analyze macromolecular characteristics of the P(tBMA-co-

DEAEMA) copolymers. 1H NMR spectra were acquired using a Varian Mercury

300 MHz NMR spectrometer. When the kinetics of ARGET ATRP were being followed, aliquots of the polymerization solution were taken at regular time intervals using a nitrogen-purged syringe and analyzed immediately by 1H NMR spectroscopy to determine conversion. Polymer film thicknesses were measured using a Beaglehole Instruments phase-modulated Picometer ellipsometer, which uses a 632.8 nm HeNe laser as the incident light source. Ellipsometric angles were measured at multiple angles ranging from 60° to 80° in 1° increments. The angle-dependent ellipsometric data were fit using a “slab-like” model to determine layer thicknesses of the dry brushes in air as well as in aqueous buffer solutions.

Dry layer thicknesses are reported as the average of multiple measurements acquired from at least 3 replicate samples. Measurements were made in a series of buffers having pH = 3, 4, 5, 6, 7, 8 and 9 (± 0.1) and an ionic strength of 30 mM.

Recipes for the buffer solutions can be found in Appendix B.

5.3.3 Synthesis of poly(tert-butyl methacrylate-co-diethylaminoethyl methacrylate)

(P(tBMA-co-DEAEMA)) via ARGET ATRP

All copolymer syntheses were run using a total monomer to initiator ratio,

[M]:[I] = 1000:1, while keeping both the total amount of monomer constant at 2.5

M and the copper(II) concentration at 100 ppm. As a representative example, a synthesis is detailed for a copolymer targeted to be 20% DEAEMA (nominally) on a molar basis. CuBr2 (1.01 mg, 0.0045 mmol) and TPMA (5.20 mg, 0.018 mmol) were added to a dry, 50 mL three-neck round bottom flask. The flask was sealed with rubber septa and purged with dry nitrogen. Then, anisole (2 mL) was added using a nitrogen-purged syringe. This solution was maintained at room temperature for 30 minutes to allow the catalyst complex to form. Once the complex had formed, as indicated by the solution turning a bright yellow color, more anisole (6 mL) was added to the flask via a nitrogen-purged syringe. A previously prepared solution of tert-butyl methacrylate (5.70 mL, 35.2 mmol),

DEAEMA (1.80 mL, 8.80 mmol) and EBIB (8.50 mg, 0.044mmol) in 1 mL of dry anisole was then added slowly to the flask via a nitrogen-purged syringe. Finally, tin (II) ethylhexanoate (20.1 mg, 0.05 mmol), which had been dissolved in dry anisole (1mL), was added immediately to the flask. Once the tin (II) ethylhexanoate was added, four freeze-pump-thaw cycles were used to remove any dissolved O2.

After the final thawing of the contents of the flask, it was lowered into an oil bath thermostatted at 35 °C to start the polymerization. The kinetics of polymerization were monitored by taking aliquots every 60 min. The polymerization was quenched after 480 min by submerging the flask into liquid nitrogen. After thawing, the

97 polymer was purified by precipitation into approximately 100 mL of an 80/20 (v/v) mixture of hexanes and ethyl ether that was maintained at −78 oC using a dry ice/acetone bath. After isolation and purification, the recovered P(tBMA-co-

DEAEMA) copolymer was analyzed by GPC and 1H NMR spectroscopy.

5.3.4 Synthesis of 3-(2-Bromoisobutyramido)propyl(trimethoxy)silane

Preparation of an initiator that could be tethered to inorganic surfaces followed protocols described by Tugulu et al.161 Rigorously dried glassware and an inert argon atmosphere were used and all transfers were done using dried, argon- purged syringes. A solution consisting of 3-aminopropyltrimethoxysilane (0.18 mL,

0.8 mmol) and triethylamine (0.12 mL, 0.8 mmol) in 10 mL of anhydrous THF was prepared in a 50 mL round bottom flask. Then, 2-bromoisobutyryl bromide (0.1 mL,

0.8 mmol) was added dropwise. The solution was then placed in an ice bath and stirred for 3 h; it was then removed from the ice bath and stirred for another 10 h at room temperature. To purify the product, triethylammonium bromide was removed by gravity filtration and the solvent was slowly removed from the filtrate by evaporation under reduced pressure. The product, 3-(2- bromoisobutyramido)propyl(trimethoxy)silane, was analyzed by 1H NMR spectroscopy (300 MHz): δ 6.87 (s, 1H), 3.53 (s,9H), 3.22 (m, 2H), 1.91 (s, 6H),

1.62 (m, 2H), 0.62 (t, J= 8.12 Hz, 2H). These findings are consistent with results from Tugulu et al.161

5.3.5 Immobilization of the ARGET ATRP silane initiator onto silicon surfaces

Silicon wafers (1 × 1.2 cm; Silicon Quest) were cleaned by immersion in a

“piranha acid” solution [7:3 (H2SO4/H2O2)] for 45 min at 100 ᵒC. (Caution: Piranha

98 acid is a strong acid and strong oxidizer.) The surfaces were removed, rinsed with copious amounts of water, and dried under a nitrogen stream. The surfaces were then irradiated via ozonolysis for another 45 min using a NovaScan PSD Series ozone generator to ensure complete oxidation of groups on the surface. The cleaned wafers were then immersed for 6 h in a 40 mM solution of 3-(2- bromoisobutyramido)propyl(trimethoxy)silane in anhydrous THF. After allowing for self-assembly the initiator-functionalized wafers were rinsed with anhydrous THF, dried with a stream of nitrogen, and then transferred to the Schlenk tube used for growing brushes by surface initiated ARGET ATRP. The surface functionalization procedure consistently yielded initiator layers having a thickness of 3 nm, as measured by multi-angle ellipsometry.

5.3.6 Synthesis of polymer brushes via surface initiated ARGET ATRP

As an example, the synthesis of a copolymer brush containing 20%

DEAEMA (nominal composition) on a molar basis is described. A home-built perforated glass platform was inserted into a Schlenk tube equipped with a magnetic stir bar. The platform holds several initiator-decorated silicon substrates level above the stirrer, which allows replicate samples to be produced during a single polymerization. The initiator-functionalized wafers were placed face-up on the platform and the Schlenk tube was purged with nitrogen to ensure the contents were in an inert atmosphere. The remainder of the synthesis was completed with the contents under a positive pressure of nitrogen. CuBr2 (1.05 mg, 0.0045 mmol),

TPMA (5.20 mg, 0.018 mmol) and dry anisole (2 mL) were added to the Schlenk tube via syringe and the solution was then stirred for 30 minutes. A solution

99 containing the comonomers in the appropriate ratio was prepared separately in a clean, dry vial the appropriate molar ratio – for this example, tBMA (5.70 mL, 35.2 mmol), and DEAEMA (1.80 mL, 8.80 mmol) were mixed. Next, the sacrificial initiator, EBIB, (8.85 mg, 0.044 mmol) was added along with dry anisole (2 mL).

This solution of comonomers and sacrificial initiator was then added to the Schlenk tube via a nitrogen-purged syringe. In all cases, the comonomer solution covered the initiator-decorated wafer. Finally, a separately-prepared solution of tin (II) ethylhexanoate (20 mg, 0.05 mmol) dissolved in anisole (1 mL) was added to the tube via a nitrogen- purged syringe. The Schlenk tube then was immediately sealed and lowered into liquid nitrogen to begin a sequence of four freeze-pump- thaw cycles. After the final thaw, the Schlenk tube was lowered into an oil bath set at 35 oC to initiate the polymerization. The mixture was stirred at 500 rpm. After

600 min, the polymerization was quenched by opening the Schlenk tube to the atmosphere and freezing the contents in LN2. After thawing, the surfaces were removed and cleaned by rinsing with dichloromethane (Fisher, reagent grade) followed by sonication for 10 min in dichloromethane to ensure that any physisorbed polymer generated from sacrificial initiator (EBIB) present in solution was removed. The dry film thicknesses of P(tBMA-co-DEAEMA) brushes were determined via ellipsometry. Free polymer present in the solution was recovered by precipitation into approximately 100 mL of an 80/20 (v/v) mixture of hexanes and ethyl ether that was maintained at −78 oC using a dry ice/acetone bath.

Following isolation, purification, and drying, these “free” polymer chains were analyzed via GPC and 1H NMR spectroscopy.

100

5.3.7 Generation of P(MAA-co-DEAEMA) brushes via post-polymerization modification

The tert-butyl groups of tBMA repeat units were cleaved to yield the polyampholyte, P(MAA-co-DEAEMA), by immersing P(tBMA-co-DEAEMA) brush- modified wafers in a solution (~3.7 mL) of dichloromethane containing 30% trifluoroacetic acid (TFA) for 30 minutes.195 The wafers were removed from the

TFA solution, rinsed with copious amounts of THF and then water. Finally, they were dried using a stream of dry nitrogen gas. The thicknesses of the deprotected brushes were measured at a few different locations on each surface by multi-angle ellipsometry

5.3.8 Ellipsometric swelling studies of P(MAA-co-DEAEMA) brushes

Brush-modified wafers were secured to a custom-built Teflon holder that fits in a cylindrical glass cell used for swelling or adsorption measurements.196–199

Buffer solutions at pH = 9, 8, 7, 6, 5, 4 and 3 (each ± 0.1 pH unit) were prepared at an ionic strength of 30 mM and introduced to the glass cell individually. To mitigate cross-contamination when changing pH, after measurements in a given buffer solution, the electrolyte solution was withdrawn and the cell was flushed with copious amounts of deionized water (refilling and draining). The wafer and cell were flushed three times using deionized water before a new buffer solution was introduced into the cell. The polymer brushes were allowed to equilibrate in the buffer solution for 20 min before multi-angle ellipsometry measurements were made. While the brush was equilibrating, the brush-modified surface was re-

101 aligned in the filled cell to ensure fluid changes did not alter the measurement geometry.

5.4 Results and Discussion

5.4.1 Kinetic studies of solution synthesis P(tBMA-co-DEAEMA) copolymers

While ATRP is tolerant to many functional groups, in practice, acrylic acid and methacrylic acid require additional considerations because these charge- negative monomers interact or react with the basic metal/halide complexes. Side reactions, such as formation of metal carboxylates, result in inefficient deactivators that lead to loss of control. Furthermore, the metal carboxylates cannot be reduced to reactivate chain growth in ATRP.200 A standard approach, therefore, is to conduct the ATRP using the “protected” monomers tert-butyl acrylate (tBA) or tert- butyl methacrylate (tBMA), followed by hydrolysis or pyrolysis of the tert-butyl ester

“protecting” groups to produce the corresponding weak polyacid, either poly(acrylic acid) or poly(methacrylic acid).131,136,201 Also, it is known that tertiary amines are capable of reducing the higher oxidation state metal/ligand catalyst complex to the active form, resulting in an increase in polymerization and potential loss of control.194 Therefore, ARGET ATRP was used in order to reduce the possibility of side reactions and termination events, and ensure control over chain growth.

Furthermore, because polymerizations intitated from flat surfaces are difficult in part because of there being a low number of active chains (usually on the order of

1015 – 1016 chains m−2), kinetic studies of solution copolymerization were performed in order to elucidate conditions that were pseudo-living/controled, as evidenced by the polymerization adhereing to first order kinetics with a persistant

102 radical population. To this end, P(tBMA-co-DEAEMA) copolymers were synthesized in anisole at 35 °C via ARGET ATRP using EBIB initiator with a total monomer concentration of 2.5 M. Random copolymers were produced across a wide range of comonomer feed ratios, as seen in Table 5-1.

Seven copolymerizations were performed in total, varying the

[tBMA]/[DEAEMA] comonomer feed ratio to examine the efficacy of ARGET ATRP over a wide range of comonomer compositions. Aliquots were taken every hour throughout the polymerizations and the random copolymers recovered after 8 h of polymerization time were analyzed by 1H NMR spectroscopy and GPC. Kinetic plots based on first-order kinetics using monomer conversion determined from 1H

NMR spectroscopy are shown in Figure 5-1 for copolymers that were synthesized by ARGET ATRP using comonomer feed ratios of 80/20, 50/50, and 20/80 tBMA/DEAEMA. The corresponding figures for all other compositions can be found in Appendix B. The relationship reflected in each of the plots (including those presented in Appendix B) indicates that the polymerization is first order with respect to monomer conversion at all comonomer compositions. The slope extracted from the kinetic profile of each of the three copolymerizations yields information about the effective rate of the ARGET ATRP reaction, keff: The slope deduced from the kinetic plot for the 80/20 tBMA/DEAEMA copolymerization is

0.0013 s-1, which is similar to but smaller than the slopes determined from polymerizations at comonomer ratios of 50/50 and 20/80, which have keff = 0.0017

-1 -1 s and keff = 0.0018 s , respectively.

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Table 5-1. Characteristics of P(tBMA-co-DEAEMA) copolymers.

Molar Compositional Actual Copolymer Mn Đ b Ratio tBMA:DEAEMA Compositiona (kDa)b

80:20 79:21 45.4 1.14

70:30 72:28 45.8 1.12

60:40 61:39 48.1 1.16

50:50 51:49 48.3 1.15

40:60 39:61 46.3 1.16

30:70 28:72 45.6 1.17

20:80 19:81 47.7 1.14 a Determined via 1H NMR spectroscopy.

bRelative to polystyrene standards.

104

T=35 °C for 8 h.

105

These values suggest that increasing the relative abundance of DEAEMA increases the rate of polymerization, which is consistent with the acceleration of

ATRP due to tertiary amines seen by Radosz et al.194 Nevertheless, excellent control is maintained in the ARGET ATRP copolymerization across the range of comonomer feed ratios, which used

[monomer]:[EBIB]:[Cu(II)Br2]:[TMPA]:[Sn(II)Oct] = 1000:1:0.1:0.4:1.

This impact of DEAEMA monomer acting as a reducing agent and affecting the rate of polymerization, Rp is expected based on the rate expression developed by Matyjaszewski et al.

[Cu(I)] R =K K [In] ⌊M⌋ (1) p p eq [Cu(II)X]

In this expression, Kp is the polymerization constant, Keq is the equilibrium constant for propagation, and [M] is the monomer concentration. Assuming steady-state kinetics, as the concentration of the activator, the Cu-metal/ligand complex Cu(I)/L, increases due to reduction of the deactivator, Cu(II)X/L, either by the Sn(II)Oct or

DEAEMA acting as reductant, the rate of polymerization increases.

Although DEAEMA exerts a slight influence on the rate of polymerization, the polymerizations are well controlled. In addition and as shown in Table 5-1, the overall composition of the polymer closely follows the comonomer feed ratio.

Copolymer composition determined by 1H NMR spectroscopy agrees with the feed ratio within 1-2 %. This suggests a lack of preference for incorporation of the two methacrylate-based comonomers. Finally, ARGET ATRP resulted in low dispersities that range from 1.12-1.17, which is lower than 50:50 random

106 polyampholyte brushes synthesized by ATRP, which had a dispersity of 1.4.138

These show that ARGET ATRP provides excellent control over copolymerization of amine containing copolymers across the entire range of comonomer feed ratios.

5.4.2 Synthesis of P(MAA-co-DEAEMA) Polymer Brushes

With useful conditions for ARGET ATRP resolved through kinetic studies of solution polymerizations, P(MAA-co-DEAEMA) brushes on flat silicon substrates were created via the “grafting from” approach illustrated in Scheme 5-1. As noted previously, the protected form of MAA was used, and t-butyl groups were subsequently removed by acid-catalyzed hydrolysis. In these polymerizations, free

EBIB initiator was added to the solution to set a monomer-to-initiator ratio of 1000.

It is well-accepted that free initiator is required to achieve controlled growth chains of surface-attached chains.202–204 P(MAA-co-DEAEMA) brushes were synthesized across comonomer feed ratios ranging from 80 mol % to 20 mol % DEAEMA and the recovered free chains used as proxies for the surface-tethered chains.

Scheme 5-1. Synthesis of P(MAA-co-DEAEMA) on silicon substrates using

ARGET ATRP.

107

In these polymerizations, free EBIB initiator was added to the solution to set a monomer-to-initiator ratio of 1000. It is well-accepted that free initiator is required to achieve controlled growth chains of surface-attached chains.202–204 P(MAA-co-

DEAEMA) brushes were synthesized across comonomer feed ratios ranging from

80 mol % to 20 mol % DEAEMA and the recovered free chains used as proxies for the surface-tethered chains. Characterization by GPC shows that the free copolymer chains have narrow dispersities (Đ < 1.2), and these are shown in Table

5-2 along with Mn values (relative to PS standards) and compositional information obtained by 1H NMR spectroscopy. The dispersities and molecular weights are similar to those obtained from kinetic studies and analysis by 1H NMR spectroscopy again shows that comonomer incorporation is smooth and follows the comonomer feed ratio. Furthermore, the ellipsometric thicknesses of the deprotected polyampholyte brushes, HPA, are large and generally correlate with the number-average molecular weight, which together suggest that the surface is uniformly covered with a dense polymer brush. This contention is supported by

AFM images showing homogenous coverage and low values of root mean square roughnesses. (See Appendix B)

The film thicknesses of the p(tBMA-co-DEAEMA) brushes, Helli, obtained from ellipsometry show an increase in thickness as composition increases up to

50 mol % DEAEMA but then decrease as incorporation of DEAEMA continues to increase. This pattern of behavior seems to not follow molecular weight of the copolymers, which suggests that sample-to-sample variations in grafting density are impacting thickness, even though each surface underwent the same treatment.

108

Table 5-2. Characterization data for P(MAA-co-DEAEMA) and dry film thicknesses.

Mol % Mn Helli HPA σ tBMA:DEAEMA Đa DEAEMAa (kDa)a (nm)b (nm)c (Chains/nm2)d

80:20 19 55.8 1.13 45±3 28±2 0.50

70:30 32 55.2 1.11 53±5 - 0.59

60:40 42 65.1 1.15 61±4 - 0.58

50:50 48 68.4 1.15 78±4 59±3 0.71

40:60 63 44.3 1.17 63±3 - 0.88

30:70 72 45.6 1.17 53±4 - 0.72

20:80 83 55.3 1.12 42±3 38±3 0.50

a Based on measurement of free chains recovered from solution. b Average thickness of protected films obtained from several measurements of multiple samples via ellipsometry. c Average ellipsometric thickness of deprotected polyampholyte brushes taken from multiple measurements of multiple samples. d

Based on dry film thickness of P(tBMA-co-DEAEMA) brushes.

109

This general sensitivity of surface-initiated polymerizations, coupled with the fact that grafting density of chains, σ, is calculated from layer thickness by the dimensional expression σ = HelliρNA⁄Mn, provides context as to why the brush characteristics vary somewhat substantially across the range of comonomer composition. Another cause for the variations in the grafting density may arise through the use of DEAEMA as a comonomer: it has been reported that aliphatic tertiary amines can reduce Cu(II) to Cu(I) and produce cation radicals.194,205,206

Thus, if DEAEMA functioned as an internal reducing agent, it would affect the equilibrium between dormant and active states or lead to additional side reactions and termination events may have a significant impact on the growth of the brush.

This may be especially influential in the case of propagation of chains from the surface, where there is a low areal density of chains. (Typically on the order of 1015

– 1016 chains m-2.)

There is additional information within the set of data that buttresses the contention that variations are due to the sensitivity of these surface initiated polymerizations: The ratios of brush heights of the parent and deprotected, polyampholyte (PA) brushes, HPA/Helli, are compared to the corresponding ratios of molecular weights, Mn,PA/Mn, and presented in Table 5-3. As outlined in the

Appendix B, molecular weights and compositions of the 80/20, 50/50, and 20/80 random copolymers allow the theoretical molecular weight of the polyampholyte brushes to be calculated.

110

Table 5-3. Comparison of “parent” and polyampholyte brush thickness and molecular weight ratios

Mol %

tBMA:DEAEMA Mn, PA/Mn HPA /Helli DEAEMAa

80:20 19 0.68 0.62

50:50 48 0.79 0.76

20:80 83 0.93 0.91

a Based on measurement of free chains recovered from solution

In light of the fact that assumptions of constant grafting density, complete transformation of P(tBMA-co-DEAEMA) to P(MAA-co-DEAEMA) by acid-catalyzed hydrolysis, and constant mass density are invoked, the correspondence between

HPA/Helli and Mn,PA/Mn as a function of composition is excellent. (See Table 5-3.)

Moreover, the best-fit line describing the relationship between copolymer composition and HPA/Helli extrapolates to 0.54 at fDEAEMA = 0.0 and to 0.98 at fDEAEMA

= 1.0. This self-consistency strongly supports for the notion that variations in brush characteristics are simply inherent to these surface-initiated polymerizations.

5.4.3 Swelling studies of polyampholyte brushes

The response of P(MAA-co-DEAEMA) brushes to changes in pH was investigated by measuring the brush thickness as solution pH was incrementally lowered from pH 9 to 3 at a constant ionic strength of 30 mM. The swelling ratio,

T/Tdry, is used to facilitate comparison of the response of different brushes as a function of pH, as shown in Figure 5-2, which presents the swelling responses of

111

P(MAA-co-DEAEMA) random copolymer brushes synthesized with comonomer compositions of 20, 50, and 80 mol % DEAEMA, respectively.

The behaviors depicted in Figure 5-2 reflect, in general, polyelectrolyte behavior at low and high pH and polyampholytic character at intermediate pH values. Moreover, they are consistent both with changes in relative amount of the two comonomers and behaviors of confined systems of charged polymers. These aspects are discussed in turn, beginning with the response at low and high pH where the brushes behave like polyelectrolytes. At low pH values, both the

DEAEMA segments and MAA are in their protonated form, rendering DEAEMA repeat units positively charged while MAA repeat units exist in their neutral carboxylic acid form. Previous neutron reflectivity studies by Deodhar and coworkers195 show that MAA is strongly hydrophobic when in its neutral

(protonated or non-dissociated) form. As a result, the polyampholyte brush containing 20% DEAEMA does not show much swelling due to the dominance of the excess hydrophobic MAA repeat units. At low pH values and as the relative amount of MAA decreases, the DEAEMA exerts greater influence over the swelling behavior: the p(DEAEMA-co-MAA) brush that is 50 mol % DEAEMA is more extended than the 20% DEAEMA brush, but less swollen than the 80% DEAEMA brush. In addition, careful inspection of the three swelling curves suggests

112

Figure 5-2. Swelling response of p(DEAEMA-co-MAA) brushes as a function of pH. Amine content varies, with (A) representing 20 mol % DEAEMA, (B) 50 mol %

DEAEMA, and (C) 80 mol % DEAEMA. In each case, the swollen thickness measured by in situ ellipsometry is normalized by the thickness of the brush measured in ambient conditions.

113 that the onset at which polybase character manifests at low pH (swelling when

DEAEMA is protonated)207 shifts to higher pH as DEAEMA content increase.

Analogous behavior is observed in the swelling response of the brushes at high pH, where the chains take on poly(acid) character due to deprotonation of

MAA repeat units. At a pH = 9, which eclipses the pKa of homopolymers of

208 DEAEMA (pKa of the conjugate acid = 7.5) and its monomer (pKa = 8.8), the

20/80 and 50/50 p(DEAEMA-co-MAA) brushes show significant swelling, with

T/Tdry ~ 6; the 80/20 p(DEAEMA-co-MAA) brush is not as strongly stretched due to the fact that the uncharged DEAEMA repeat units are only slightly hydrophilic.207,209,210 The behaviors observed at both low and high pH are consistent with the notion that electrolytic group present in abundance exerts greater influence over the swelling behavior.

There are two aspects related to ellipsometric measurements of dry and swollen brushes that merit discussion. First, while the measured swollen thicknesses are normalized by the dry layer thickness measured for each of the brushes, the dry layer thickness must be interpreted with caution. As shown by

Deodhar et al.,195 PMAA brushes and copolymer brushes incorporating electrolytic

MAA repeat units and neutral, water-soluble hydroxyethyl methacrylate monomers absorb water from humid environments. Thus, even in the state referred to as “dry”, the brushes may contain a significant fraction of water: Deodhar et al. showed that the volume fraction of water in the dry brush approached 30%, with the apparent trend that higher grafting density of chains led to a lower volume fraction of absorbed water.195 In the case of these p(DEAEMA-co-MAA) brushes, there is no

114 reason to expect that the amount of latent hydration is consistent across the various compositions studied. Second, and as discussed in our recent studies of polypeptide brushes,211 because stretching of brush chains in a good solvent produces a smoothly-decaying segment density profile and the analysis of ellipsometry data used herein assumes a “step-like” profile, ellipsometric thicknesses underrepresent the height of solvent-swollen brushes, with increasing departure as chain extension increases. For these reasons, the pattern of behavior displayed by a given brush and comparison between the different brushes are more informative than specific values of the swelling ratio.

Another point worth discussing is the apparent change in pKa exhibited by the p(DEAEMA-co-MAA) polyampholyte brushes. As noted by Deodhar et al., the

195 pKa of PMAA brushes shifts to higher values compared to free chains. This change, which is consistent with behaviors observed by Genzer and coworkers,212 can be attributed to the local dielectric environment of a brush, which differs from that of solvent water because of the high concentration of repeat units confined near the solid/fluid interface.212–214 The local dielectric function is affected by the dipole moment of the ion-pairs, the concentration of counterions and co-ions, and the local composition. In addition, because the dielectric function follows the segment density profile of an electrolytic brush,75,76 the inhomogeneous environment affects the acidity and basicity of the groups (the charge dissociation equilibrium) due to the varying intermolecular interactions and electrostatic potential.213 This also has been experimentally observed through titration studies of weak poly(acid) brushes, specifically PAA and PMAA. Ober et al. also showed

115

surf that the effective pKa at the film/air interface pKa , is about two units smaller than

bulk the pKa of the bulk material, pKa . This implies that carboxylic acid groups on repeat units further away from the surface are more readily ionized than acid

215 surf groups near the surface. They suggested that at pH < pKa , all of the acid groups are protonated making the poly(acid) brushes relatively hydrophobic. In this

surf case, the chains of the brush are collapsed. When the pH approaches pKa , the acid groups in the uppermost layer begin to ionize, making the outer region of the polymer brush hydrophilic. Even though a major portion of the chains remain collapsed, the ionized segments at the periphery stretch into solution. They also showed that a small fraction of acid groups located near the solid substrate remain uncharged, even at very high pH. While these behaviors are specific to weak poly(acid)s, it is likely that there are analogous effects in which the local pKb differs

bulk from pKb that occur with a weak poly(base) brush, such as pDEAEMA. Moreover, we would anticipate these effects would add complexity to the pH-mediated swelling behavior of polyampholytic brushes due to the presence of both types of weak electrolytic groups.

Sequence effects will also play a role in how the polyampholyte brushes behave in response to changes in pH. Using Monte-Carlo simulations, Ulrich and coworkers examined how the overall arrangement of repeating units in diblock, alternating, and random copolymer designs influence titration curves of weak polyampholytes.216 For A-B polyampholytes in which A is a weak acid and B is a

A B weak base, when pKa < pH < pKa , both A and B repeating units are fully charged,

116

A B but when pH < pKa or when pH > pKa , only B or A repeats are charged, respectively. With both random and alternating polyampholytes, Ulrich and coworkers saw that titration curves of repeating units of type A (type B) are shifted to smaller (higher) pH values as the ionization of acidic (basic) groups is facilitated, particularly at low salt concentrations where screening is lessened.216 Their studies show that these effects are more pronounced in an alternating polyampholyte designs compared to random copolyampholytes because the regular sequential arrangement of charge-positive and charge-negative repeat units facilitates ion- pairing through short-range electrostatic interactions. In addition to ionization of repeat units being promoted, chain flexibility plays an important role, with flexible polyampholytes adopting more compact conformations upon self-neutralization.216

Finally, in addition to changes in the effective pKa (pKb) due to either local dielectric effects or due to self-neutralization brought about by ion-pairing through short- range interactions along the chain, their simulations show that the charging mechanism of a weak polyampholyte also depends on the comonomer stoichiometry. Increases in compositional asymmetry results in chain extension due to an excess of charges.77

The patterns of behavior displayed in the titration curves of the polyampholyte brushes are consistent with these traits. For all of the p(MAA-co-

DEAEMA) polyampholytes, when the pH is in an intermediate range, the charge- negative carboxylates of MAA repeat units are neutralized by ion pairing with oppositely charged, quaternized amines of DEAEMA repeating units, leading to a reduction in brush height. Presumably, this is facilitated by the close proximity of

117 charge-negative and charge-positive repeat units, either through short-range interactions within the chain or interactions with neighboring chains in the crowded, confined environment of the brush. In this case, the chains adopt a collapsed conformation and exclude water and counterions from the interfacial layer. This self-neutralization at an intermediate pH, known as the isoelectric point, is associated with classical polyampholyte behavior and has been observed by

Brittain and coworkers in their studies of poly(acrylic acid-b-2-vinylpyridine) and poly(acrylic acid-b-4vinyl pyridine).136

In blocky polymer systems, the isoelectric point is sharp and generally well defined; however, in random polymer systems the isoelectric point appears to be less well defined due to the facility with which self-neutralization is achieved through short range interactions along the chain. All of the swelling curves measured for the polyampholytic brushes studied here go through a minimum between 4.5 < pH < 7.5. In the case of the compositionally symmetric brush (50%

DEAEMA; Figure 5-2b) there is a distinct minimum in swollen thickness at pH = 5, which is assigned as an apparent isoelectric point. Whether conformational entropy prohibits the brush from self-neutralization cannot be inferred from these measurements. Given that the layer does not collapse more strongly, it seems reasonable to posit that some electrolytic groups are neutralized by counterions present in the buffer, which leads to osmotic swelling of the layer.

The swelling behavior of the 50 mol % DEAEMA brush is similar to the behavior observed by Tran and coworkers in their studies of the effect of grafting density on poly(methacrylic acid-r-dimethylaminoethyl methacrylate) random

118 polyampholyte brushes made with a 1:1 comonomer ratio.138 For a brush having a grafting density of 0.132 ch nm-2, they observed a stronger collapse over the range

5 ≤ pH ≤ 7 with an isoelectric point between pH = 6.0 and 7.6. Swelling of the layer was evident at pH = 8 and pH = 3 which are characteristic of a polyacid and a polybase brush, respectively. While a symmetric polyampholyte brush having a grafting density of 0.193 ch nm-2 remained swollen across the range of pH values studied, in both cases and consistent with the behaviors reported here, the random polyampholyte brushes remain swollen with T/Tdry ≈ 3 when self-neutralized, rather than collapsing toward their dry layer thicknesses.

The behavior of the polyampholyte brushes containing 20 mol% DEAEMA and 80 mol % DEAEMA are also consistent with a tendency for self-neutralization at intermediate pH values, but the presence of excess weak acid and weak base and inherent solubility differences seem to dominate swelling behaviors. (Recall, uncharged MAA is more hydrophobic than uncharged DEAEMA, and charged

MAA is more hydrophilic than charged DEAEMA.) Specifically, while there is a minimum in the swelling response of brushes containing 20 mol% DEAEMA, the hydrophobicity of uncharged PMAA dominates the behavior, as reflected in the nearly-constant and low values of swelling ratio, T/Tdry, throughout the window where self-neutralization is observed to occur for the compositionally-symmetric poly(MAA-co-DEAEMA) brush. The drastic reduction in MAA content in the brush containing 80% DEAEMA allows a comparatively larger, yet modest swelling response, especially across the range of pH values where self-neutralization occurs. In addition, the minimum that may be ascribed to self-neutralization is

119 shifted to higher pH (relative to that of the compositionally symmetric brush system). As the percentage of DEAEMA is increased within the chain, the location of the minima shifts to higher pH values as the amine/acid ratio is increased.217

However, it is unclear whether this shift is a consequence of the onset of poly(base) character coupled with diminished influence of methacrylic acid repeat units, or if the relative abundance of DEAEMA repeating units and their quaternization is facilitating ionization of MAA repeating units and resulting in self-neutralization.

Finally, because the local dielectric environment influences the chemical equilibrium between charged and neutral states,74-76 it seems reasonable to conclude that self-neutralization will be similarly affected by the local composition within the polyampholyte brush as well as by the tendency of random polyampholytes to facilitate ionization and self-neutralization.

5.4.4 Swelling response in the presence of betaine zwitterion

The impact of an added zwitterion on the swelling behavior of the polyampholyte brushes was examined by adding the zwitterion, betaine (30 mM), to the buffer solutions. (This maintains the ionic strength at 30 mM.218) Because organic osmolytes like betaine are known to stabilize the folded conformation of proteins, they are classified as chaotropic because of their ability to disrupt the hydrogen bonding of water at protein surfaces. In keeping with this expectation and as seen in Figure 5-3, the total swelling of the polyampholyte brushes is diminished in the presence of betaine. However, the characteristic shape of the swelling curve of is not altered: At low and high pH, the brushes display polyelectrolyte behavior but at intermediate pH values they exhibit polyampholyte

120

121 behavior. In addition, the minima that were attributed to self-neutralization of the polyampholyte brushes containing 50 mol % DEAEMA and 80 mol % DEAEMA are also unchanged, while swelling of the polyampholyte brush containing 20 mol

% DEAEMA is strongly restricted by the excess MAA repeat units.

Given the consistency between the patterns of behaviors observed in buffered solution and in buffers containing the zwitterion, it appears that the main effect of added zwitterion swelling is to render the polyampholyte brushes more hydrophobic. The decrease in swelling exhibited at all pH values may be a direct result of betaine adding hydrophobicity to the chains upon pairing with charged groups along the chain, or it may be an indirect effect that is caused by depletion of the small molecule osmolyte from the polyampholyte interface. This latter speculation is consistent with a growing body of evidence suggesting that depletion forces are the general mechanism by which proteins and peptides are stabilized by organic osmolytes like betaine.219–223

5.5 Conclusions

We describe the use of ARGET ATRP to synthesize random copolymers of

DEAEMA and tBMA, and use conditions identified through solution polymerization to create polyampholyte brushes via “grafting from” method followed by chemical conversion to p(MAA-co-DEAEMA) polyampholytes. Although an excess of either comonomer influences the swelling behavior, multiangle ellipsometry measurements show that the random copolyampholyte brushes follow polyelectrolyte behavior at low and high pH, but polyampholytic character marked by a self-neutralization is observed at intermediate pH values. While the extent of

122 swelling is decreased when betaine a zwitterion is added to the solution, the mode of action appears to share commonality with how organic osmolytes stabilize intrinsically disordered proteins and peptides. Thus, the extension of ARGET

ARTP to create compositionally-tuned polyampholyte brushes and the use of those brushes as models to understand swelling response in different electrolytic environments may provide a basis to address structural responses and aggregation of biological polymers and the action of cryoprotectants.

123

Chapter 6: Summary, Conclusions, and Future Work

124

6.1 Summary and Conclusions

Understanding how the design of polymer thin films affects their structure and response is crucial for developing polymer thin film systems with desirable attributes that can be used in a variety of applications. The work I described in this dissertation provides fundamental insight into two categories of polymer thin films: polymer brushes and copolymer-grafted nanoparticles. The work on copolymer- grafted nanoparticles illustrates that the miscibility of copolymer nanocomposites can be tuned by changing the strength of the enthalpic interactions between the graft chains and the polymer matrix. This is primarily accomplished by changing the overall composition of the copolymer synthesized. Because I changed copolymer composition by changing the ratio of comonomers, I examined the kinetics of the solution polymerization. ARGET ATRP was used to synthesize poly(methyl methacrylate-ran-cyclohexyl methacrylate) while varying the overall composition of cyclohexyl methacrylate within the polymer. Then, copolymer grafted nanoparticles were synthesized and dispersed in a chemically dissimilar polystyrene matrix. AFM imaging of the copolymer nanocomposites showed that the composition of the random copolymer tethered to the silica nanoparticle surface influences dispersion into the chemically dissimilar polystyrene matrix.

Since the copolymer-grafted nanoparticles were successfully dispersed in the chemically dissimilar polystyrene matrix, fused deposition modeling was used to examine whether the mechanical properties of nanocomposites comprised of

PGNPs dispersed in a polystyrene matrix improved. This was done by first optimizing the temperature at which the polystyrene nanocomposite was printed.

125

Unfortunately, there was a large variation across replicate samples, which seemed to indicate that there was no appreciable change in thermomechanical properties.

Though the addition of copolymer-grafted nanoparticles did not show any reinforcement in the thermomechanical properties of the polystyrene matrix, several parameters that would help facilitate the printing of polystyrene were elucidated.

One of the major hurdles facing implementation of polymer-grafted nanocomposites is difficulty in producing large quantities of polymer-grafted materials. I have addressed this scaling problem by using the more facile “grafting through” approach, which utilizes conventional free radical polymerization.

Specifically, poly(methyl methacrylate) grafted nanoparticles were synthesized by modifying the surface of silica nanoparticle with small molecule silanes that contain a vinyl group. The presence of grafted poly(methyl methacrylate) chains was confirmed by thermogravimetric analysis, Fourier-transform infrared attenuated total reflection spectroscopy, and dynamic light scattering. This grafting through approach was able to produce polymer grafted nanoparticles on the gram scale rather than on the milligram scale, which is advantageous.

Finally, my research on polymer thin films included studies of polyampholyte brushes. A novel element of this work was the use of ARGET ATRP to control the surface initiated polymerization of the polyampholyte brushes. Specifically,

ARGET ATRP was used to synthesize random copolymers of diethylaminoethyl methacrylate and tert-butyl methacrylate in a controlled fashion and conditions identified through solution polymerization were used to create polyampholyte

126 brushes via “grafting from” method followed by chemical conversion to poly(methacrylic acid-co-diethylaminoethyl methacrylate) polyampholytes.

Multiangle ellipsometry measurements showed that the random copolyampholyte brushes follow polyelectrolyte behavior at low and high pH, but polyampholytic character marked by self-neutralization was observed at intermediate pH values.

While the extent of swelling was decreased when betaine, a zwitterion, was added to the solution. In total, my dissertation research provides considerable insight into how polymer design impacts polymer thin film structure and behavior.

6.2 Future Work

By generating new insights into the influence of structural design of polymer thin films and how the overall design affects their response, this work has shown that are many challenges that need to be addressed. To address these challenges and help drive the application of these soft materials, the following set of studies are recommended:

I. Further optimization needs to be done on the FDM printing of

PS. The temperature range studied to this point consistent storage

modulus, but the overall incorporation of nanofillers needs to be

investigated further. By incorporating the nanoparticles by coextrusion of

the pellets with the polymer-grafted nanoparticles, many uncertainties are

introduced. One uncertainty concerns whether the extruded filament

contains the polymer-grafted nanoparticles at the correct loading.

Aggregation or phase separation maybe occurring if the nanoparticles are

not smoothly incorporated within the system. One way to address this issue

127 is to look at other blending methods. One method that shows some promise is compounding. In compounding the polymer-grafted nanoparticles are melt blended with the matrix polymer for some period of time. After this melt blending the polymer nanocomposite is pelletized and then extruded in a more controlled way. This method will allow for better control over the loading level of the polymer grafted nanoparticles and potentially create better incorporation of the nanoparticles. After using this method, the thermomechanical properties of FDM printed samples can then be investigated, providing a clearer understanding of how the nanofiller is affecting the matrix polymer and properties.

II. To verify that the grafting through method is a viable method for not only the scaling of polymer grafted nanoparticles, but also results in reinforcement of polymer nanocomposites, the thermomechanical properties will need to be investigated and compared with a similar system.

This can be completed by using FDM with a symmetric polymer nanocomposite system, such as poly(methyl methacrylate)-grafted nanoparticles dispersed in a poly(methyl methacrylate) system. The grafting densities created using the grafting through method are comparable to the grafting densities of polymer grafted nanoparticles achieved from a grafting to method. Thermomechanical properties of FDM printed specimens

(polymer-grafted nanocomposites) made each of these ways should be tested to determine if the grafting-through method provides sufficient enhancement of properties. If so, then the grafting through method would

128

be a more facile and robust approach to creating polymer nanocomposites

on a larger scale.

III. With the polyampholyte brush system, it may be valuable to

explore how the composition of the polyampholyte affects the uptake and

release of small molecules loaded into the brush. Such studies would be

relevant to drug-release from polymeric delivery vehicles and also would

add to the current understanding of how the polyampholyte brush structure

affects response.

In total, these themes of study would further elaborate links between polymer design, thin film structure and the overall response and application of these thin films. This body of work would help to advance the development of polymer nanocomposites and anti-fouling or drug delivery systems leading to more robust materials for a variety of applications.

129

List of References

130

1. M. Jain, R. Rajan, S.-H. Hyon and K. Matsumura, Hydrogelation of Dextran-

based Polyampholytes with Cryoprotective Properties via Click Chemistry,

Biomater Sci, 2014, 2, 308–317.

2. A. Akthakul, R. F. Salinaro and A. M. Mayes, Antifouling Polymer

Membranes with Subnanometer Size Selectivity, Macromolecules, 2004,

37, 7663–7668.

3. Lee Seung Yun, Kim Hee Jin, Patel Rajkumar, Im Se Joon, Kim Jong Hak

and Min Byoung Ryul, Silver Nanoparticles Immobilized on Tthin Film

Composite Polyamide Membrane: Characterization, Nanofiltration,

Antifouling Properties, Polym. Adv. Technol., 2007, 18, 562–568.

4. C. Zhao, L.-Y. Li, M.-M. Guo and J. Zheng, Functional Polymer Thin Films

Designed for Antifouling Materials and Biosensors, Chem. Pap., 2012, 66,

323–339.

5. Y.-C. Tsai, S.-C. Li and J.-M. Chen, Cast Thin Film Biosensor Design Based

on a Nafion Backbone, a Multiwalled Carbon Nanotube Conduit, and a

Glucose Oxidase Function, Langmuir, 2005, 21, 3653–3658.

6. G. Urban, G. Jobst, F. Kohl, A. Jachimowicz, F. Olcaytug, O. Tilado, P.

Goiser, G. Nauer, F. Pittner, T. Schalkhammer and E. Mann-Buxbaum,

Miniaturized Thin-film Biosensors using Covalently Immobilized Glucose

Oxidase, Biosens. Bioelectron., 1991, 6, 555–562.

131

7. X. Wang, Y.-G. Kim, C. Drew, B.-C. Ku, J. Kumar and L. A. Samuelson,

Electrostatic Assembly of Conjugated Polymer Thin Layers on Electrospun

Nanofibrous Membranes for Biosensors, Nano Lett., 2004, 4, 331–334.

8. K. M. Coakley and M. D. McGehee, Conjugated Polymer Photovoltaic Cells,

Chem. Mater., 2004, 16, 4533–4542.

9. L. A. A. Pettersson, L. S. Roman and O. Inganäs, Modeling Photocurrent

Action Spectra of Photovoltaic Devices based on Organic Thin Films, J.

Appl. Phys., 1999, 86, 487–496.

10. P. Peumans, S. Uchida and S. R. Forrest, Efficient Bulk Heterojunction

Photovoltaic Cells using Small-molecular-weight Organic Thin Films,

Nature, 2003, 425, 158–162.

11. S. Swaraj, C. Wang, H. Yan, B. Watts, J. Lüning, C. R. McNeill and H. Ade,

Nanomorphology of Bulk Heterojunction Photovoltaic Thin Films Probed

with Resonant Soft X-ray Scattering, Nano Lett., 2010, 10, 2863–2869.

12. P. S. Ramanujam, S. Hvilsted, F. Ujhelyi, P. Koppa, E. Lörincz, G. Erdei

and G. Szarvas, Physics and Technology of Optical Storage in Polymer

Thin Films, Synth. Met., 2001, 124, 145–150.

13. P. Rochon, E. Batalla and A. Natansohn, Optically Induced Surface

Gratings on Azoaromatic Polymer Films, Appl. Phys. Lett., 1995, 66, 136–

138.

14. M. Ulbricht, Advanced Functional Polymer Membranes, Polymer, 2006, 47,

2217–2262.

132

15. B.-H. Jeong, E. M. V. Hoek, Y. Yan, A. Subramani, X. Huang, G. Hurwitz,

A. K. Ghosh and A. Jawor, Interfacial Polymerization of Thin Film

Nanocomposites: A New Concept for Reverse Osmosis Membranes, J.

Membr. Sci., 2007, 294, 1–7.

16. M. S. Wilson and S. Gottesfeld, Thin-film Catalyst Layers for Polymer

Electrolyte Fuel Cell Electrodes, J. Appl. Electrochem., 1992, 22, 1–7.

17. I. Borukhov and L. Leibler, Stabilizing Grafted Colloids in a Polymer Melt:

Favorable Enthalpic Interactions, Phys. Rev. E, 2000, 62, 41–45.

18. A. Striolo, Controlled Assembly of Spherical Nanoparticles: Nanowires and

Spherulites, Small, 2007, 3, 628–635.

19. Q. Lan, L. F. Francis and F. S. Bates, Silica Nanoparticle Dispersions in

Homopolymer Versus Block Copolymer, J. Polym. Sci. Part B Polym. Phys.,

2007, 45, 2284–2299.

20. M. A. Holmes, M. E. Mackay and R. K. Giunta, Nanoparticles for Dewetting

Suppression of Thin Polymer Films used in Chemical Sensors, J.

Nanoparticle Res., 2007, 9, 753–763.

21. S. Chandran, N. Begam and J. K. Basu, Dispersion of Polymer Grafted

Nanoparticles in Polymer Nanocomposite Films: Insights from Surface X-

ray Scattering and Microscopy, J. Appl. Phys., 2014, 116, 222203–222217.

22. G. D. Smith and D. Bedrov, Dispersing Nanoparticles in a Polymer Matrix:

Are Long, Dense Polymer Tethers Really Necessary?, Langmuir, 2009, 25,

11239–11243.

133

23. A. Jayaraman and K. S. Schweizer, Effective Interactions and Self-

Assembly of Hybrid Polymer Grafted Nanoparticles in a Homopolymer

Matrix, Macromolecules, 2009, 42, 8423–8434.

24. M. W. Matsen and J. M. Gardiner, Autophobic Dewetting of Homopolymer

on a Brush and Entropic Attraction between Opposing Brushes in a

Homopolymer Matrix, J. Chem. Phys., 2001, 115, 2794–2804.

25. D. Sunday, J. Ilavsky and D. L. Green, A Phase Diagram for Polymer-

Grafted Nanoparticles in Homopolymer Matrices, Macromolecules, 2012,

45, 4007–4011.

26. C. Xu, K. Ohno, V. Ladmiral and R. J. Composto, Dispersion of Polymer-

grafted Magnetic Nanoparticles in Homopolymers and Block Copolymers,

Polymer, 2008, 49, 3568–3577.

27. J. Liu, Y. Gao, D. Cao, L. Zhang and Z. Guo, Nanoparticle Dispersion and

Aggregation in Polymer Nanocomposites: Insights from Molecular

Dynamics Simulation, Langmuir, 2011, 27, 7926–7933.

28. K. K. Sriramoju and V. Padmanabhan, Controlling the Location of Bare

Nanoparticles in Polymer-Nanoparticle Blend Films by Adding Polymer-

Grafted Nanoparticles, Phys. Rev. Lett., ,

DOI:10.1103/PhysRevLett.114.258301.

29. D. Meng, S. K. Kumar, J. M. D. Lane and G. S. Grest, Effective Interactions

Between Grafted Nanoparticles in a Polymer Matrix, Soft Matter, 2012, 8,

5002.

134

30. S. E. Harton and S. K. Kumar, Mean-field Theoretical Analysis of Brush-

coated Nanoparticle Dispersion in Polymer Matrices, J. Polym. Sci. Part B

Polym. Phys., 2008, 46, 351–358.

31. P. Akcora, H. Liu, S. K. Kumar, J. Moll, Y. Li, B. C. Benicewicz, L. S.

Schadler, D. Acehan, A. Z. Panagiotopoulos, V. Pryamitsyn, V. Ganesan,

J. Ilavsky, P. Thiyagarajan, R. H. Colby and J. F. Douglas, Anisotropic Self-

assembly of Spherical Polymer-grafted Nanoparticles, Nat. Mater., 2009, 8,

354–359.

32. I. Borukhov and L. Leibler, Enthalpic Stabilization of Brush-Coated Particles

in a Polymer Melt, Macromolecules, 2002, 35, 5171–5182.

33. C. F. Laub and J. T. Koberstein, Effect of Brush Polydispersity on the

Interphase between End-Grafted Brushes and Polymeric Matrixes,

Macromolecules, 1994, 27, 5016–5023.

34. M. Müller and L. G. MacDowell, Wetting of a Short Chain Liquid on a Brush:

First-order and Critical Wetting Transitions, EPL Europhys. Lett., 2001, 55,

221–227.

35. B. Kim, D. Y. Ryu, V. Pryamitsyn and V. Ganesan, Dewetting of PMMA on

PS−Brush Substrates, Macromolecules, 2009, 42, 7919–7923.

36. T. B. Martin, K. I. S. Mongcopa, R. Ashkar, P. Butler, R. Krishnamoorti and

A. Jayaraman, Wetting–Dewetting and Dispersion–Aggregation Transitions

Are Distinct for Polymer Grafted Nanoparticles in Chemically Dissimilar

Polymer Matrix, J. Am. Chem. Soc., 2015, 10624–10631.

135

37. S. Ojha, A. Dang, C. M. Hui, C. Mahoney, K. Matyjaszewski and M. R.

Bockstaller, Strategies for the Synthesis of Thermoplastic Polymer

Nanocomposite Materials with High Inorganic Filling Fraction, Langmuir,

2013, 29, 8989–8996.

38. M. Moniruzzaman and K. I. Winey, Polymer Nanocomposites Containing

Carbon Nanotubes, Macromolecules, 2006, 39, 5194–5205.

39. Liu, I. Y. Phang, L. Shen, S. Y. Chow and W.-D. Zhang, Morphology and

Mechanical Properties of Multiwalled Carbon Nanotubes Reinforced Nylon-

6 Composites, Macromolecules, 2004, 37, 7214–7222.

40. S. Sinha Ray, K. Okamoto and M. Okamoto, Structure−Property

Relationship in Biodegradable Poly(butylene succinate)/Layered Silicate

Nanocomposites, Macromolecules, 2003, 36, 2355–2367.

41. A. C. Balazs, C. Singh and E. Zhulina, Modeling the Interactions between

Polymers and Clay Surfaces through Self-Consistent Field Theory,

Macromolecules, 1998, 31, 8370–8381.

42. X. Zhao, Q. Zhang, D. Chen and P. Lu, Enhanced Mechanical Properties of

Graphene-Based Poly(vinyl alcohol) Composites, Macromolecules, 2010,

43, 2357–2363.

43. H. Kim, A. A. Abdala and C. W. Macosko, Graphene/Polymer

Nanocomposites, Macromolecules, 2010, 43, 6515–6530.

44. J. Moraes, K. Ohno, T. Maschmeyer and S. Perrier, Synthesis of silica–

polymer core–shell Nanoparticles by Reversible Addition–Fragmentation

Chain Transfer Polymerization, Chem. Commun., 2013, 49, 9077–9088.

136

45. T. Ahmad, M. Shahazad, M. Ubaidullah, J. Ahmed, A. Khan and A. M. El-

Toni, Structural Characterization and Dielectric Properties of Ceria–titania

Nanocomposites in Low Ceria Region, Mater. Res. Express, 2017, 4,

125016.

46. T. E. Motaung, A. S. Luyt, F. Bondioli, M. Messori, M. L. Saladino, A.

Spinella, G. Nasillo and E. Caponetti, PMMA–titania Nanocomposites:

Properties and Thermal Degradation Behaviour, Polym. Degrad. Stab.,

2012, 97, 1325–1333.

47. A. S. Zeraati, M. Arjmand and U. Sundararaj, Silver Nanowire/MnO 2

Nanowire Hybrid Polymer Nanocomposites: Materials with High Dielectric

Permittivity and Low Dielectric Loss, ACS Appl. Mater. Interfaces, 2017, 9,

14328–14336.

48. A. Shayesteh Zeraati, S. A. Mirkhani and U. Sundararaj, Enhanced

Dielectric Performance of Polymer Nanocomposites Based on CNT/MnO 2

Nanowire Hybrid Nanostructure, J. Phys. Chem. C, 2017, 121, 8327–8334.

49. B. J. Ash, R. W. Siegel and L. S. Schadler, Mechanical Behavior of

Alumina/Poly(methyl methacrylate) Nanocomposites, Macromolecules,

2004, 37, 1358–1369.

50. K. Yoshimoto, T. S. Jain, K. V. Workum, P. F. Nealey and J. J. de Pablo,

Mechanical Heterogeneities in Model Polymer Glasses at Small Length

Scales, Phys. Rev. Lett., 2004, 93, 175501–175505.

51. M. E. Mackay, General Strategies for Nanoparticle Dispersion, Science,

2006, 311, 1740–1743.

137

52. P. H. T. Vollenberg and D. Heikens, Particle Size Dependence of the

Young’s Modulus of Filled Polymers: 1. Preliminary Experiments, Polymer,

1989, 30, 1656–1662.

53. M. H. Al-Saleh, W. H. Saadeh and U. Sundararaj, EMI shielding

effectiveness of carbon based nanostructured polymeric materials: A

comparative study, Carbon, 2013, 60, 146–156.

54. E. Tang, G. Cheng, X. Pang, X. Ma and F. Xing, Synthesis of Nano-

ZnO/poly(methyl methacrylate) Composite Microsphere through Emulsion

Polymerization and its UV-shielding Property, Colloid Polym. Sci., 2006,

284, 422–428.

55. V. Padmanabhan, Percolation of High-density Polymer Regions in

Nanocomposites: The Underlying Property for Mechanical Reinforcement,

J. Chem. Phys., 2013, 139, 144904–144912.

56. S. Gupta, Q. Zhang, T. Emrick, A. C. Balazs and T. P. Russell, Entropy-

driven Segregation of Nanoparticles to Cracks in Multilayered Composite

Polymer Structures, Nat. Mater., 2006, 5, 229–233.

57. A. C. Balazs, T. Emrick and T. P. Russell, Nanoparticle Polymer

Composites: Where Two Small Worlds Meet, Science, 2006, 314, 1107–

1110.

58. D. M. Trombly and V. Ganesan, Curvature Effects Upon Interactions of

Polymer-grafted Nanoparticles in Chemically Identical Polymer Matrices, J.

Chem. Phys., 2010, 133, 154904–154911.

138

59. A. Jayaraman, Polymer Grafted Nanoparticles: Effect of Chemical and

Physical Heterogeneity in Polymer Grafts on Particle Assembly and

Dispersion, J. Polym. Sci. Part B Polym. Phys., 2013, 51, 524–534.

60. S. Chandran, S. Saw, A. K. Kandar, C. Dasgupta, M. Sprung and J. K. Basu,

Suspensions of Polymer-grafted Nanoparticles with Added Polymers—

Structure and Effective Pair-interactions, J. Chem. Phys., 2015, 143,

084902–084914.

61. V. Goel, J. Pietrasik, H. Dong, J. Sharma, K. Matyjaszewski and R.

Krishnamoorti, Structure of Polymer Tethered Highly Grafted Nanoparticles,

Macromolecules, 2011, 44, 8129–8135.

62. T. A. Witten and P. A. Pincus, Colloid stabilization by Long Grafted

Polymers, Macromolecules, 1986, 19, 2509–2513.

63. J. U. Kim and M. W. Matsen, Interaction between Polymer-Grafted

Particles, Macromolecules, 2008, 41, 4435–4443.

64. P. Akcora, S. K. Kumar, V. García Sakai, Y. Li, B. C. Benicewicz and L. S.

Schadler, Segmental Dynamics in PMMA-Grafted Nanoparticle

Composites, Macromolecules, 2010, 43, 8275–8281.

65. M. K. Corbierre, N. S. Cameron, M. Sutton, K. Laaziri and R. B. Lennox,

Gold Nanoparticle/Polymer Nanocomposites: Dispersion of Nanoparticles

as a Function of Capping Agent Molecular Weight and Grafting Density,

Langmuir, 2005, 21, 6063–6072.

139

66. B. J. Kim, J. Bang, C. J. Hawker and E. J. Kramer, Effect of Areal Chain

Density on the Location of Polymer-Modified Gold Nanoparticles in a Block

Copolymer Template, Macromolecules, 2006, 39, 4108–4114.

67. J. Y. Lee, A. C. Balazs, R. B. Thompson and R. M. Hill, Self-Assembly of

Amphiphilic Nanoparticle−Coil “Tadpole” Macromolecules,

Macromolecules, 2004, 37, 3536–3539.

68. A. Jayaraman and K. S. Schweizer, Effect of the Number and Placement of

Polymer Tethers on the Structure of Concentrated Solutions and Melts of

Hybrid Nanoparticles, Langmuir, 2008, 24, 11119–11130.

69. A. Chremos, A. Z. Panagiotopoulos, H.-Y. Yu and D. L. Koch, Structure of

Solvent-free Grafted Nanoparticles: Molecular Dynamics and Density-

functional Theory, J. Chem. Phys., 2011, 135, 114901–114913.

70. A. Jayaraman and K. S. Schweizer, Effective Interactions, Structure, and

Phase Behavior of Lightly Tethered Nanoparticles in Polymer Melts,

Macromolecules, 2008, 41, 9430–9438.

71. N. Nair and A. Jayaraman, Self-Consistent PRISM Theory−Monte Carlo

Simulation Studies of Copolymer Grafted Nanoparticles in a Homopolymer

Matrix, Macromolecules, 2010, 43, 8251–8263.

72. T. B. Martin and A. Jayaraman, Tuning the Wetting–dewetting and

Dispersion–aggregation Transitions in Polymer Nanocomposites using

Composition of Graft and Matrix Polymers, Mater. Res. Express, 2016, 3,

034001–034014.

140

73. B. Vorselaars, J. U. Kim, T. L. Chantawansri, G. H. Fredrickson and M. W.

Matsen, Self-consistent Field Theory for Diblock Copolymers Grafted to a

Sphere, Soft Matter, 2011, 7, 5128–5137.

74. X. Zhu, L. Wang, J. Lin and L. Zhang, Ordered Nanostructures Self-

Assembled from Block Copolymer Tethered Nanoparticles, ACS Nano,

2010, 4, 4979–4988.

75. T. B. Martin, A. Seifpour and A. Jayaraman, Assembly of Copolymer

Functionalized Nanoparticles: A Monte Carlo Simulation Study, Soft Matter,

2011, 7, 5952–5964.

76. E. R. Chan, L. C. Ho and S. C. Glotzer, Computer Simulations of Block

Copolymer Tethered Nanoparticle Self-assembly, J. Chem. Phys., 2006,

125, 064905–064913.

77. D. M. Trombly, V. Pryamitsyn and V. Ganesan, Interfacial Properties of

Statistical Copolymer Brushes in Contact with Homopolymer Melts, J.

Chem. Phys., 2011, 134, 154903–154915.

78. T. Hashimoto, H. Tanaka and H. Hasegawa, Ordered Structure in Mixtures

of a Block Copolymer and Homopolymers. 2. Effects of Molecular Weights

of Homopolymers, Macromolecules, 1990, 23, 4378–4386.

79. K. Heo, C. Miesch, T. Emrick and R. C. Hayward, Thermally Reversible

Aggregation of Gold Nanoparticles in Polymer Nanocomposites through

Hydrogen Bonding, Nano Lett., 2013, 13, 5297–5302.

80. S. H. Masood, Intelligent rapid prototyping with fused deposition modelling,

Rapid Prototyp. J., 1996, 2, 24–33.

141

81. M. N. Cooke, J. P. Fisher, D. Dean, C. Rimnac and A. G. Mikos, Use of

Stereolithography to Manufacture Critical-sized 3D Biodegradable

Scaffolds for Bone Ingrowth, J. Biomed. Mater. Res. B Appl. Biomater.,

2003, 64B, 65–69.

82. J. R. Tumbleston, D. Shirvanyants, N. Ermoshkin, R. Janusziewicz, A. R.

Johnson, D. Kelly, K. Chen, R. Pinschmidt, J. P. Rolland, A. Ermoshkin, E.

T. Samulski and J. M. DeSimone, Continuous Liquid Interface Production

of 3D Objects, Science, 2015, 347, 1349–1352.

83. I. Cooperstein, M. Layani and S. Magdassi, 3D Printing of Porous Structures

by UV-curable O/W Emulsion for Fabrication of Conductive Objects, J.

Mater. Chem. C, 2015, 3, 2040–2044.

84. M. Agarwala, D. Bourell, J. Beaman, H. Marcus and J. Barlow, Direct

Selective Laser Sintering of Metals, Rapid Prototyp. J., 1995, 1, 26–36.

85. B. C. Gross, J. L. Erkal, S. Y. Lockwood, C. Chen and D. M. Spence,

Evaluation of 3D Printing and Its Potential Impact on Biotechnology and the

Chemical Sciences, Anal. Chem., 2014, 86, 3240–3253.

86. D. Bak, Rapid Prototyping or Rapid Production? 3D Printing Processes

Move Industry Towards the Latter, Assem. Autom., 2003, 23, 340–345.

87. S. J. Leigh, R. J. Bradley, C. P. Purssell, D. R. Billson and D. A. Hutchins,

A Simple, Low-Cost Conductive Composite Material for 3D Printing of

Electronic Sensors, PLOS ONE, 2012, 7, e49365.

142

88. V. F. Scalfani, C. H. Turner, P. A. Rupar, A. H. Jenkins and J. E. Bara, 3D

Printed Block Copolymer Nanostructures, J. Chem. Educ., 2015, 92, 1866–

1870.

89. D. B. Short, Use of 3D Printing by Museums: Educational Exhibits, Artifact

Education, and Artifact Restoration, 3D Print. Addit. Manuf., 2015, 2, 209–

215.

90. A. Goyanes, J. Wang, A. Buanz, R. Martínez-Pacheco, R. Telford, S.

Gaisford and A. W. Basit, 3D Printing of Medicines: Engineering Novel Oral

Devices with Unique Design and Drug Release Characteristics, Mol.

Pharm., 2015, 12, 4077–4084.

91. A. Goyanes, A. B. M. Buanz, G. B. Hatton, S. Gaisford and A. W. Basit, 3D

Printing of Modified-release Aminosalicylate (4-ASA and 5-ASA) Tablets,

Eur. J. Pharm. Biopharm., 2015, 89, 157–162.

92. A. Goyanes, A. B. M. Buanz, A. W. Basit and S. Gaisford, Fused-filament

3D Printing (3DP) for Fabrication of Tablets, Int. J. Pharm., 2014, 476, 88–

92.

93. Y. L. Yap and W. Y. Yeong, Additive Manufacture of Fashion and Jewellery

Products: A Mini Review, Virtual Phys. Prototyp., 2014, 9, 195–201.

94. S. V. Murphy and A. Atala, 3D Bioprinting of Tissues and Organs, Nat.

Biotechnol., 2014, 32, 773–785.

95. Y. L. Kong, I. A. Tamargo, H. Kim, B. N. Johnson, M. K. Gupta, T.-W. Koh,

H.-A. Chin, D. A. Steingart, B. P. Rand and M. C. McAlpine, 3D Printed

Quantum Dot Light-Emitting Diodes, Nano Lett., 2014, 14, 7017–7023.

143

96. F. Ilievski, A. D. Mazzeo, R. F. Shepherd, X. Chen and G. M. Whitesides,

Soft Robotics for Chemists, Angew. Chem. Int. Ed., 2011, 50, 1890–1895.

97. J. T. Muth, D. M. Vogt, R. L. Truby, Y. Mengüç, D. B. Kolesky, R. J. Wood

and J. A. Lewis, Embedded 3D Printing of Strain Sensors within Highly

Stretchable Elastomers, Adv. Mater., 2014, 26, 6307–6312.

98. E. Kroll and D. Artzi, Enhancing Aerospace Engineering Students’ Learning

with 3D Printing Wind‐tunnel Models, Rapid Prototyp. J., 2011, 17, 393–

402.

99. B. P. Conner, G. P. Manogharan, A. N. Martof, L. M. Rodomsky, C. M.

Rodomsky, D. C. Jordan and J. W. Limperos, Making Sense of 3-D Printing:

Creating a Map of Additive Manufacturing Products and Services, Addit.

Manuf., 2014, 1–4, 64–76.

100. A. R. Torrado Perez, D. A. Roberson and R. B. Wicker, Fracture Surface

Analysis of 3D-Printed Tensile Specimens of Novel ABS-Based Materials,

J. Fail. Anal. Prev., 2014, 14, 343–353.

101. L. J. Love, V. Kunc, O. Rios, C. E. Duty, A. M. Elliott, B. K. Post, R. J. Smith

and C. A. Blue, The Importance of Carbon Fiber to Polymer Additive

Manufacturing, J. Mater. Res., 2014, 29, 1893–1898.

102. T.-M. Wang, J.-T. Xi and Y. Jin, A model research for prototype warp

deformation in the FDM process, Int. J. Adv. Manuf. Technol., 2007, 33,

1087–1096.

144

103. S. C. Ligon, R. Liska, J. Stampfl, M. Gurr and R. Mülhaupt, Polymers for 3D

Printing and Customized Additive Manufacturing, Chem. Rev., 2017, 117,

10212–10290.

104. S. F. A. Acquah, B. E. Leonhardt, M. S. Nowotarski, J. M. Magi, K. A.

Chambliss, T. E. S. Venzel, S. D. Delekar and L. A. Al-Hariri, in Carbon

Nanotubes - Current Progress of their Polymer Composites, eds. M. R.

Berber and I. H. Hafez, InTech, 2016.

105. D. Maillard, S. K. Kumar, B. Fragneaud, J. W. Kysar, A. Rungta, B. C.

Benicewicz, H. Deng, L. C. Brinson and J. F. Douglas, Mechanical

Properties of Thin Glassy Polymer Films Filled with Spherical Polymer-

Grafted Nanoparticles, Nano Lett., 2012, 12, 3909–3914.

106. S. K. Kumar, N. Jouault, B. Benicewicz and T. Neely, Nanocomposites with

Polymer Grafted Nanoparticles, Macromolecules, 2013, 46, 3199–3214.

107. N. Tsubokawa, K. Kobayashi and Y. Sone, Grafting Onto Carbon Black by

the Reaction of Reactive Carbon Black having Masked Isocyanate or Acyl

Azide Group with Functional Polymers, J. Polym. Sci. Part Polym. Chem.,

1988, 26, 223–233.

108. H. Ben Ouada, H. Hommel, A. . Legrand, H. Balard and E. Papirer,

Organization of the Layers of Polyethylene Oxide Grafted with Different

Densities on Silica, J. Colloid Interface Sci., 1988, 122, 441–449.

109. K. Bridger and B. Vincent, The Terminal Grafting of Poly(ethylene oxide)

Chains to Silica Surfaces, Eur. Polym. J., 1980, 16, 1017–1021.

145

110. C. M. Hui, J. Pietrasik, M. Schmitt, C. Mahoney, J. Choi, M. R. Bockstaller

and K. Matyjaszewski, Surface-Initiated Polymerization as an Enabling Tool

for Multifunctional (Nano-)Engineered Hybrid Materials, Chem. Mater.,

2014, 26, 745–762.

111. K. Matyjaszewski, H. Dong, W. Jakubowski, J. Pietrasik and A. Kusumo,

Grafting from Surfaces for “Everyone”: ARGET ATRP in the Presence of

Air, Langmuir, 2007, 23, 4528–4531.

112. R. Barbey, L. Lavanant, D. Paripovic, N. Schüwer, C. Sugnaux, S. Tugulu

and H.-A. Klok, Polymer Brushes via Surface-Initiated Controlled Radical

Polymerization: Synthesis, Characterization, Properties, and Applications,

Chem. Rev., 2009, 109, 5437–5527.

113. L. Wu, U. Glebe and A. Böker, Synthesis of Hybrid Silica Nanoparticles

Densely Grafted with Thermo and pH Dual-Responsive Brushes via

Surface-Initiated ATRP, Macromolecules, 2016, 49, 9586–9596.

114. J.-R. Wu, G.-Q. Lai, H.-J. Yu and Z.-H. Luo, Synthesis and Characterization

of Poly(γ-methacryloxypropyltrimethoxysilane)-grafted Silica Hybrid

Nanoparticles Prepared by Surface-initiated Atom Transfer Radical

Polymerization, J. Appl. Polym. Sci., 2012, 124, 3821–3830.

115. M. Henze, D. Mädge, O. Prucker and J. Rühe, “Grafting Through”:

Mechanistic Aspects of Radical Polymerization Reactions with Surface-

Attached Monomers, Macromolecules, 2014, 47, 2929–2937.

146

116. P. Guenoun, A. Schalchli, D. Sentenac, J. W. Mays and J. J. Benattar, Free-

standing Black Films of Polymers: A Model of Charged Brushes in

Interaction, Phys. Rev. Lett., 1995, 74, 3628–3631.

117. Y. Mir, P. Auroy and L. Auvray, Density Profile of Polyelectrolyte Brushes,

Phys. Rev. Lett., 1995, 75, 2863–2866.

118. M. Balastre, F. Li, P. Schorr, J. Yang, J. W. Mays and M. V. Tirrell, A Study

of Polyelectrolyte Brushes Formed from Adsorption of Amphiphilic Diblock

Copolymers Using the Surface Forces Apparatus, Macromolecules, 2002,

35, 9480–9486.

119. X. Guo and M. Ballauff, Spherical polyelectrolyte brushes: Comparison

between annealed and quenched brushes, Phys. Rev. E, 2001, 64,

0514061–0514069.

120. G. Romet-Lemonne, J. Daillant, P. Guenoun, J. Yang and J. W. Mays,

Thickness and Density Profiles of Polyelectrolyte Brushes: Dependence on

Grafting Density and Salt Concentration, Phys. Rev. Lett., 2004, 93,

148301–148304.

121. H. Ahrens, S. Förster and C. A. Helm, Polyelectrolyte Brushes Grafted at

the Air/water Interface, Macromolecules, 1997, 30, 8447–8452.

122. N. P. Shusharina and P. Linse, Oppositely Charged Polyelectrolytes

Grafted onto Planar Surface: Mean-field Lattice Theory, Eur. Phys. J. E,

2001, 6, 147–155.

147

123. L. Chen, H. Merlitz, S. He, C.-X. Wu and J.-U. Sommer, Polyelectrolyte

Brushes: Debye Approximation and Mean-Field Theory, Macromolecules,

2011, 44, 3109–3116.

124. R. Israels, F. A. M. Leermakers and G. J. Fleer, On the Theory of Grafted

Weak Polyacids, Macromolecules, 1994, 27, 3087–3093.

125. E. B. Zhulina, T. M. Birshtein and O. V. Borisov, Theory of Ionizable Polymer

Brushes, Macromolecules, 1995, 28, 1491–1499.

126. T. Jiang, Z. Li and J. Wu, Structure and Swelling of Grafted Polyelectrolytes:

Predictions from a Nonlocal Density Functional Theory, Macromolecules,

2007, 40, 334–343.

127. O. J. Hehmeyer, G. Arya, A. Z. Panagiotopoulos and I. Szleifer, Monte Carlo

Simulation and Molecular Theory of Tethered Polyelectrolytes, J. Chem.

Phys., 2007, 126, 244902–244913.

128. F. S. Csajka and C. Seidel, Strongly Charged Polyelectrolyte Brushes: A

Molecular Dynamics Study, Macromolecules, 2000, 33, 2728–2739.

129. C. Seidel, Strongly Stretched Polyelectrolyte Brushes, Macromolecules,

2003, 36, 2536–2543.

130. N. A. Kumar and C. Seidel, Polyelectrolyte Brushes with Added Salt,

Macromolecules, 2005, 38, 9341–9350.

131. N. Ayres, S. G. Boyes and W. J. Brittain, Stimuli-Responsive Polyelectrolyte

Polymer Brushes Prepared via Atom-Transfer Radical Polymerization,

Langmuir, 2007, 23, 182–189.

148

132. Z. Xiong, B. Peng, X. Han, C. Peng, H. Liu and Y. Hu, Dual-stimuli

Responsive Behaviors of Diblock Polyampholyte PDMAEMA-b-PAA in

Aqueous Solution, J. Colloid Interface Sci., 2011, 356, 557–565.

133. K. Yu and Y. Han, Effect of Block Sequence and Block Length on the

Stimuli-responsive Behavior of Polyampholyte Brushes: Hydrogen Bonding

and Electrostatic Interaction as the Driving Force for Surface

Rearrangement, Soft Matter, 2009, 5, 759–768.

134. Y. K. Jhon, S. Arifuzzaman, A. E. Özçam, D. J. Kiserow and J. Genzer,

Formation of Polyampholyte Brushes via Controlled Radical Polymerization

and Their Assembly in Solution, Langmuir, 2012, 28, 872–882.

135. A. B. Lowe and C. L. McCormick, Synthesis and Solution Properties of

Zwitterionic Polymers, Chem. Rev., 2002, 102, 4177–4190.

136. N. Ayres, C. D. Cyrus and W. J. Brittain, Stimuli-responsive Surfaces using

Polyampholyte Polymer Brushes Prepared via Atom Transfer Radical

Polymerization, Langmuir, 2007, 23, 3744–3749.

137. S. Sanjuan and Y. Tran, Synthesis of Random Polyampholyte Brushes by

Atom Transfer Radical Polymerization, J. Polym. Sci. Part Polym. Chem.,

2008, 46, 4305–4319.

138. S. Sanjuan and Y. Tran, Stimuli-Responsive Interfaces Using Random

Polyampholyte Brushes, Macromolecules, 2008, 41, 8721–8728.

139. E. Ji, D. G. Whitten and K. S. Schanze, pH-Dependent Optical Properties

of a Poly(phenylene ethynylene) Conjugated Polyampholyte, Langmuir,

2011, 27, 1565–1568.

149

140. P. G. Higgs and J.-F. Joanny, Theory of Polyampholyte Solutions, J. Chem.

Phys., 1991, 94, 1543–1554.

141. A. Akinchina, N. P. Shusharina and P. Linse, Diblock Polyampholytes

Grafted onto Spherical Particles: Monte Carlo Simulation and Lattice Mean-

Field Theory, Langmuir, 2004, 20, 10351–10360.

142. N. P. Shusharina, E. B. Zhulina, A. V. Dobrynin and M. Rubinstein, Scaling

Theory of Diblock Polyampholyte Solutions, Macromolecules, 2005, 38,

8870–8881.

143. L.-J. Qu, X. Man, C. C. Han, D. Qiu and D. Yan, Responsive Behaviors of

Diblock Polyampholyte Brushes within Self-Consistent Field Theory, J.

Phys. Chem. B, 2012, 116, 743–750.

144. V. Yamakov, A. Milchev, H. J. Limbach, B. Dünweg and R. Everaers,

Conformations of Random Polyampholytes, Phys. Rev. Lett., 2000, 85,

4305–4308.

145. R. K. Das and R. V. Pappu, Conformations of Intrinsically Disordered

Proteins are Influenced by Linear Sequence Distributions of Oppositely

Charged Residues, Proc. Natl. Acad. Sci., 2013, 110, 13392–13397.

146. D. Srivastava and M. Muthukumar, Sequence Dependence of

Conformations of Polyampholytes, Macromolecules, 1996, 29, 2324–2326.

147. J. N. Bright, M. J. Stevens, J. Hoh and T. B. Woolf, Characterizing the

Function of Unstructured Proteins: Simulations of Charged Polymers Under

Confinement, J. Chem. Phys., 2001, 115, 4909–4918.

150

148. Q. Cao and H. You, Morphologies of Spherical Polyampholyte Brushes:

Effects of Counterion Valence and Charged Monomer Sequence, Polymer,

2017, 113, 233–246.

149. M. Baratlo and H. Fazli, Brushes of Flexible, Semiflexible, and Rodlike

Diblock Polyampholytes: Molecular Dynamics Simulation and Scaling

Analysis, Phys. Rev. E, 2010, 81, 011801–011807.

150. M. Baratlo and H. Fazli, Molecular Dynamics Simulation of Semiflexible

Polyampholyte Brushes--The Effect of Charged Monomers Sequence, Eur.

Phys. J. E, 2009, 29, 131–138.

151. N. A. Koratkar, J. Suhr, A. Joshi, R. S. Kane, L. S. Schadler, P. M. Ajayan

and S. Bartolucci, Characterizing Energy Dissipation in Single-walled

Carbon Nanotube Polycarbonate Composites, Appl. Phys. Lett., 2005, 87,

063102–063105.

152. G. M. Russo, G. P. Simon and L. Incarnato, Correlation between

Rheological, Mechanical, and Barrier Properties in New Copolyamide-

Based Nanocomposite Films, Macromolecules, 2006, 39, 3855–3864.

153. E. R. Chan, X. Zhang, C.-Y. Lee, M. Neurock and S. C. Glotzer, Simulations

of Tetra-Tethered Organic/Inorganic Nanocube−Polymer Assemblies,

Macromolecules, 2005, 38, 6168–6180.

154. C. R. Iacovella, A. S. Keys, M. A. Horsch and S. C. Glotzer, Icosahedral

Packing of Polymer-tethered Nanospheres and Stabilization of the Gyroid

Phase, Phys. Rev. E, 2007, 75, 040801–040805.

151

155. Zhang, M. A. Horsch, M. H. Lamm and S. C. Glotzer, Tethered Nano

Building Blocks: Toward a Conceptual Framework for Nanoparticle Self-

Assembly, Nano Lett., 2003, 3, 1341–1346.

156. Y. Li, T. M. Krentz, L. Wang, B. C. Benicewicz and L. S. Schadler, Ligand

Engineering of Polymer Nanocomposites: From the Simple to the Complex,

ACS Appl. Mater. Interfaces, 2014, 6, 6005–6021.

157. J. Pyun, S. Jia, T. Kowalewski, G. D. Patterson and K. Matyjaszewski,

Synthesis and Characterization of Organic/Inorganic Hybrid Nanoparticles:

Kinetics of Surface-Initiated Atom Transfer Radical Polymerization and

Morphology of Hybrid Nanoparticle Ultrathin Films, Macromolecules, 2003,

36, 5094–5104.

158. K. Matyjaszewski and N. V. Tsarevsky, Nanostructured Functional

Materials Prepared by Atom Transfer Radical Polymerization, Nat. Chem.,

2009, 1, 276–288.

159. J. Nicolas, Y. Guillaneuf, C. Lefay, D. Bertin, D. Gigmes and B. Charleux,

Nitroxide-mediated Polymerization, Prog. Polym. Sci., 2013, 38, 63–235.

160. M. Nishimoto, H. Keskkula and D. R. Paul, Blends of Poly(styrene-co-

acrylonitrile) and Methyl Methacrylate Based Copolymers, Macromolecules,

1990, 23, 3633–3639.

161. S. Tugulu, A. Arnold, I. Sielaff, K. Johnsson and H.-A. Klok, Protein-

functionalized Polymer Brushes, Biomacromolecules, 2005, 6, 1602–1607.

162. K. Matyjaszewski and J. Xia, Atom Transfer Radical Polymerization, Chem.

Rev., 2001, 101, 2921–2990.

152

163. G. R. Brannock, J. W. Barlow and D. R. Paul, Blends of Styrene/maleic

Anhydride Copolymers with Polymethacrylates, J. Polym. Sci. Part B Polym.

Phys., 1991, 29, 413–429.

164. T. Fukuda, M. Nagata and H. Inagaki, Light Scattering from Ternary

Solutions. 1. Dilute Solutions of Polystyrene and Poly(methyl methacrylate),

Macromolecules, 1984, 17, 548–553.

165. T. Fukuda, M. Nagata and H. Inagaki, Light Scattering from Polymer Blends

in Solution. 2. Nondilute Solutions of Polystyrene and Poly(methyl

methacrylate), Macromolecules, 1986, 19, 1411–1416.

166. A. C. Balazs, I. C. Sanchez, I. R. Epstein, F. E. Karasz and W. J. MacKnight,

Effect of Sequence Distribution on the Miscibility of Polymer/copolymer

Blends, Macromolecules, 1985, 18, 2188–2191.

167. N. Mohan, P. Senthil, S. Vinodh and N. Jayanth, A review on composite

materials and process parameters optimisation for the fused deposition

modelling process, Virtual Phys. Prototyp., 2017, 12, 47–59.

168. S. Hwang, E. I. Reyes, K. Moon, R. C. Rumpf and N. S. Kim, Thermo-

mechanical Characterization of Metal/Polymer Composite Filaments and

Printing Parameter Study for Fused Deposition Modeling in the 3D Printing

Process, J. Electron. Mater., 2015, 44, 771–777.

169. D. B. G. Williams and M. Lawton, Drying of Organic Solvents: Quantitative

Evaluation of the Efficiency of Several Desiccants, J. Org. Chem., 2010, 75,

8351–8354.

153

170. X. Wang, J. L. Davis, J. P. Hinestrosa, J. W. Mays and S. M. Kilbey, Control

of Self-Assembled Structure through Architecturally and Compositionally

Complex Block Copolymer Surfactant Mixtures, Macromolecules, 2014, 47,

7138–7150.

171. J. L. Davis, X. Wang, K. Bornani, J. P. Hinestrosa, J. W. Mays and S. M.

Kilbey, Solution Properties of Architecturally Complex Multiarm Star Diblock

Copolymers in a Nonselective and Selective Solvent for the Inner Block,

Macromolecules, 2016, 49, 2288–2297.

172. X. Wang, J. L. Davis, B. M. Aden, B. S. Lokitz and S. M. Kilbey, Versatile

Synthesis of Amine-Reactive Microgels by Self-Assembly of Azlactone-

Containing Block Copolymers, Macromolecules, 2018, 51, 3691–3701.

173. D. Zhao, M. Di Nicola, M. M. Khani, J. Jestin, B. C. Benicewicz and S. K.

Kumar, Role of Block Copolymer Adsorption Versus Bimodal Grafting on

Nanoparticle Self-assembly in Polymer Nanocomposites, Soft Matter, 2016,

12, 7241–7247.

174. Y.-P. Wang, X.-W. Pei, X.-Y. He and K. Yuan, Synthesis of well-defined,

polymer-grafted silica nanoparticles via reverse ATRP, Eur. Polym. J.,

2005, 41, 1326–1332.

175. A. Zhu, A. Cai, J. Zhang, H. Jia and J. Wang, PMMA-grafted-silica/PVC

Nanocomposites: Mechanical Performance and Barrier Properties, J. Appl.

Polym. Sci., 108, 2189–2196.

176. P. Pincus, Colloid Stabilization with Grafted Polyelectrolytes,

Macromolecules, 1991, 24, 2912–2919.

154

177. M. Chen, W. H. Briscoe, S. P. Armes and J. Klein, Lubrication at

Physiological Pressures by Polyzwitterionic Brushes, Science, 2009, 323,

1698–1701.

178. J. Klein, Repair or Replacement: A Joint Perspective, Science, 2009, 323,

47–48.

179. A. Wittemann, B. Haupt and M. Ballauff, Adsorption of Proteins on Spherical

Polyelectrolyte Brushes in Aqueous Solution, Phys. Chem. Chem. Phys.,

2003, 5, 1671–1677.

180. M. Ballauff and O. Borisov, Polyelectrolyte Brushes, Curr. Opin. Colloid

Interface Sci., 2006, 11, 316–323.

181. F. Zhou and W. T. S. Huck, Surface Grafted Polymer Brushes as Ideal

Building Blocks for “Smart” Surfaces, Phys Chem Chem Phys, 2006, 8,

3815–3823.

182. T. E. Patten and K. Matyjaszewski, Atom Transfer Radical Polymerization

and the Synthesis of Polymeric Materials, Adv. Mater., 1998, 10, 901–915.

183. T. Ando, M. Kato, M. Kamigaito and M. Sawamoto, Living Radical

Polymerization of Methyl Methacrylate with Ruthenium Complex: Formation

of Polymers with Controlled Molecular Weights and Very Narrow

Distributions, Macromolecules, 1996, 29, 1070–1072.

184. K. Matyjaszewski, T. E. Patten and J. Xia, Controlled/“living” Radical

Polymerization. Kinetics of the Homogeneous Atom Transfer Radical

Polymerization of Styrene, J. Am. Chem. Soc., 1997, 119, 674–680.

155

185. W. Jakubowski and K. Matyjaszewski, Activators Regenerated by Electron

Transfer for Atom-Transfer Radical Polymerization of (Meth)acrylates and

Related Block Copolymers, Angew. Chem. Int. Ed., 2006, 45, 4482–4486.

186. K. Min, H. Gao and K. Matyjaszewski, Use of Ascorbic Acid as Reducing

Agent for Synthesis of Well-Defined Polymers by ARGET ATRP,

Macromolecules, 2007, 40, 1789–1791.

187. J. Pietrasik, H. Dong and K. Matyjaszewski, Synthesis of High Molecular

Weight Poly(styrene- co -acrylonitrile) Copolymers with Controlled

Architecture, Macromolecules, 2006, 39, 6384–6390.

188. H. Dong, W. Tang and K. Matyjaszewski, Well-Defined High-Molecular-

Weight Polyacrylonitrile via Activators Regenerated by Electron Transfer

ATRP, Macromolecules, 2007, 40, 2974–2977.

189. L. Mueller, W. Jakubowski, W. Tang and K. Matyjaszewski, Successful

Chain Extension of Polyacrylate and Polystyrene Macroinitiators with

Methacrylates in an ARGET and ICAR ATRP, Macromolecules, 2007, 40,

6464–6472.

190. J. Pietrasik, C. M. Hui, W. Chaladaj, H. Dong, J. Choi, J. Jurczak, M. R.

Bockstaller and K. Matyjaszewski, Silica-Polymethacrylate Hybrid Particles

Synthesized Using High-Pressure Atom Transfer Radical Polymerization,

Macromol. Rapid Commun., 2011, 32, 295–301.

191. Y. Shen, H. Tang and S. Ding, Catalyst Separation in Atom Transfer Radical

Polymerization, Prog. Polym. Sci., 2004, 29, 1053–1078.

156

192. W. Lin, S. Nie, Q. Zhong, Y. Yang, C. Cai, J. Wang and L. Zhang,

Amphiphilic Miktoarm Star Copolymer (PCL)3-(PDEAEMA-b-PPEGMA)3

as pH-sensitive Micelles in the Delivery of Anticancer Drug, J. Mater. Chem.

B, 2014, 2, 4008–4020.

193. P. V. Mendonça, S. E. Averick, D. Konkolewicz, A. C. Serra, A. V. Popov,

T. Guliashvili, K. Matyjaszewski and J. F. J. Coelho, Straightforward ARGET

ATRP for the Synthesis of Primary Amine Polymethacrylate with Improved

Chain-End Functionality under Mild Reaction Conditions, Macromolecules,

2014, 47, 4615–4621.

194. H. Tang, M. Radosz and Y. Shen, CuBr2/N,N,N′,N′-Tetra[(2-

pyridal)methyl]ethylenediamine/Tertiary Amine as a Highly Active and

Versatile Catalyst for Atom-Transfer Radical Polymerization via Activator

Generated by Electron Transfer, Macromol. Rapid Commun., 2006, 27,

1127–1131.

195. C. Deodhar, E. Soto-Cantu, D. Uhrig, P. Bonnesen, B. S. Lokitz, J. F.

Ankner and S. M. Kilbey, Hydration in Weak Polyelectrolyte Brushes, ACS

Macro Lett., 2013, 2, 398–402.

196. Y. Shen, S. Desseaux, B. Aden, B. S. Lokitz, S. M. Kilbey, Z. Li and H.-A.

Klok, Shape-Persistent, Thermoresponsive Polypeptide Brushes Prepared

by Vapor Deposition Surface-Initiated Ring-Opening Polymerization of α-

Amino Acid N -Carboxyanhydrides, Macromolecules, 2015, 48, 2399–2406.

197. S. B. Rahane, J. A. Floyd, A. T. Metters and S. M. Kilbey, Swelling Behavior

of Multiresponsive Poly(methacrylic acid)-block--poly(N-

157

isopropylacrylamide) Brushes Synthesized Using Surface-Initiated

Photoiniferter-Mediated Photopolymerization, Adv. Funct. Mater., 2008, 18,

1232–1240.

198. J. P. Hinestrosa, J. Alonzo, J. W. Mays and S. M. Kilbey, Role of Surface

Reorganization on Preferential Adsorption of Macromolecular Ensembles at

the Solid/Fluid Interface, Macromolecules, 2009, 42, 7913–7918.

199. J. Alonzo, J. P. Hinestrosa, J. W. Mays and S. M. Kilbey, Kinetics of

Preferential Adsorption of Amphiphilic Star Block Copolymers that Tether

by Their Corona Blocks at the Solid/Fluid Interface, Macromolecules, 2014,

47, 4048–4055.

200. A. Y. Sankhe, S. M. Husson and S. M. Kilbey, Effect of Catalyst Deactivation

on Polymerization of Electrolytes by Surface-Confined Atom Transfer

Radical Polymerization in Aqueous Solutions, Macromolecules, 2006, 39,

1376–1383.

201. K. A. Davis and K. Matyjaszewski, Atom Transfer Radical Polymerization of

tert -Butyl Acrylate and Preparation of Block Copolymers, Macromolecules,

2000, 33, 4039–4047.

202. B. Zhao and W. J. Brittain, Synthesis, Characterization, and Properties of

Tethered Polystyrene- b -polyacrylate Brushes on Flat Silicate Substrates,

Macromolecules, 2000, 33, 8813–8820.

203. X. Kong, T. Kawai, J. Abe and T. Iyoda, Amphiphilic Polymer Brushes

Grown from the Silicon Surface by Atom Transfer Radical Polymerization,

Macromolecules, 2001, 34, 1837–1844.

158

204. M. Ejaz, K. Ohno, Y. Tsujii and T. Fukuda, Controlled Grafting of a Well-

Defined Glycopolymer on a Solid Surface by Surface-Initiated Atom

Transfer Radical Polymerization, Macromolecules, 2000, 33, 2870–2874.

205. H. Dong and K. Matyjaszewski, ARGET ATRP of 2-(Dimethylamino)ethyl

Methacrylate as an Intrinsic Reducing Agent, Macromolecules, 2008, 41,

6868–6870.

206. J. F. Weiss, G. Tollin and J. T. Yoke III, Reactions of Triethylamine with

Copper Halides. II. Internal Oxidation-reduction of

Dichlorobis(triethylamine) Copper (II), Inorg. Chem., 1964, 3, 1344–1348.

207. J. D. Willott, T. J. Murdoch, B. A. Humphreys, S. Edmondson, G. B. Webber

and E. J. Wanless, Critical Salt Effects in the Swelling Behavior of a Weak

Polybasic Brush, Langmuir, 2014, 30, 1827–1836.

208. P. van de Wetering, E. E. Moret, N. M. E. Schuurmans-Nieuwenbroek, M.

J. van Steenbergen and W. E. Hennink, Structure−Activity Relationships of

Water-Soluble Cationic Methacrylate/Methacrylamide Polymers for

Nonviral Gene Delivery, Bioconjug. Chem., 1999, 10, 589–597.

209. J. D. Willott, T. J. Murdoch, B. A. Humphreys, S. Edmondson, E. J. Wanless

and G. B. Webber, Anion-Specific Effects on the Behavior of pH-Sensitive

Polybasic Brushes, Langmuir, 2015, 31, 3707–3717.

210. J. D. Willott, B. A. Humphreys, T. J. Murdoch, S. Edmondson, G. B. Webber

and E. J. Wanless, Hydrophobic Effects within the Dynamic pH-response of

Polybasic Tertiary Amine Methacrylate Brushes, Phys. Chem. Chem.

Phys., 2015, 17, 3880–3890.

159

211. Y. Shen, S. Desseaux, B. Aden, B. S. Lokitz, S. M. Kilbey, Z. Li and H.-A.

Klok, Shape-Persistent, Thermoresponsive Polypeptide Brushes Prepared

by Vapor Deposition Surface-Initiated Ring-Opening Polymerization of α-

Amino Acid N -Carboxyanhydrides, Macromolecules, 2015, 48, 2399–2406.

212. T. Wu, P. Gong, I. Szleifer, P. Vlček, V. Šubr and J. Genzer, Behavior of

Surface-Anchored Poly(acrylic acid) Brushes with Grafting Density

Gradients on Solid Substrates: 1. Experiment, Macromolecules, 2007, 40,

8756–8764.

213. P. Gong, T. Wu, J. Genzer and I. Szleifer, Behavior of Surface-Anchored

Poly(acrylic acid) Brushes with Grafting Density Gradients on Solid

Substrates: 2. Theory, Macromolecules, 2007, 40, 8765–8773.

214. R. Kumar, B. G. Sumpter and S. M. Kilbey, Charge Regulation and Local

Dielectric Function in Planar Polyelectrolyte Brushes, J. Chem. Phys., 2012,

136, 234901.

215. R. Dong, M. Lindau and C. K. Ober, Dissociation Behavior of Weak

Polyelectrolyte Brushes on a Planar Surface, Langmuir, 2009, 25, 4774–

4779.

216. S. Ulrich, M. Seijo and S. Stoll, A Monte Carlo Study of Weak

Polyampholytes: Stiffness and Primary Structure Influences on Titration

Curves and Chain Conformations †, J. Phys. Chem. B, 2007, 111, 8459–

8467.

160

217. Turner Alfrey, Jr, Raymond M. Fous, H. Morawetz and H. Pinner,

Amphoteric Polyelectrolytes. 11. Copolymers of Methacrylic Acid and

Diethylaminoethyl Methacrylate, J. Am. Chem. Soc., 1952, 74, 438–441.

218. E. Stellwagen, J. D. Prantner and N. C. Stellwagen, Do Zwitterions

Contribute to the Ionic Strength of a Solution?, Anal. Biochem., 2008, 373,

407–409.

219. Y.-T. Liao, A. C. Manson, M. R. DeLyser, W. G. Noid and P. S. Cremer,

Trimethylamine N -oxide Stabilizes Proteins via a Distinct Mechanism

Compared with Betaine and Glycine, Proc. Natl. Acad. Sci., 2017, 114,

2479–2484.

220. S. Jamal, A. Kumari, A. Singh, S. Goyal and A. Grover, Conformational

Ensembles of α-Synuclein Derived Peptide with Different Osmolytes from

Temperature Replica Exchange Sampling, Front. Neurosci., 2017, 11, 1–

11.

221. R. Politi and D. Harries, Enthalpically Driven Peptide Stabilization by

Protective Osmolytes, Chem. Commun., 2010, 46, 6449.

222. S. Sukenik, L. Sapir and D. Harries, Balance of Enthalpy and Entropy in

Depletion Forces, Curr. Opin. Colloid Interface Sci., 2013, 18, 495–501.

223. N. Srinivasan, M. Bhagawati, B. Ananthanarayanan and S. Kumar, Stimuli-

sensitive Intrinsically Disordered Protein Brushes, Nat. Commun., 2014, 5,

5145.

161

Appendices

162

Appendix A - Chapter 3: Fused deposition modeling of Poly(methyl methacrylate-r-cyclohexyl methacrylate) Copolymer-grafted Nanoparticles

in a Chemically Dissimilar Matrix

163

Dynamic Mechanical Analysis

Figure A-1. Stress-strain curves measured by DMA used to determine the Young’s modulus of the printed pure PS. All samples were run at 230 ⁰C.

Figure A-2. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 235 ⁰C.

164

Figure A-3. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 240 ⁰C.

Figure A-4. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 245 ⁰C.

165

Figure A-5. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 250 ⁰C.

Figure A-6. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 255 ⁰C.

166

Figure A-7. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 260 ⁰C.

Figure A-8. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 265 ⁰C.

167

Figure A-9. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed pure PS. All samples were run at 270 ⁰C.

Figure A-10. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 235 ⁰C.

168

Figure A-11. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 265 ⁰C.

Figure A-12. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % bare silica nanoparticles in a PS matrix. All samples were run at 270 ⁰C.

169

Figure A-13. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % 90:10

CHMA:MMA grafted nanoparticles in a PS matrix. All samples were run at 265 ⁰C.

Figure A-14. Stress-strain curves measured by DMA used to determine the

Young’s modulus of the printed nanocomposites comprised of 1 wt % 90:10

CHMA:MMA grafted nanoparticles in a PS matrix. All samples were run at 270 ⁰C.

170

Appendix B - Chapter 5: ARGET-ATRP Synthesis and Swelling Response of

Compositionally Varied Poly(methacrylic acid-co-N, N-diethylaminoethyl

methacrylate) Polyampholyte Brushes

171

Analysis of Parent and Deprotected Brush Thicknesses and Molecular

Weights

Compositional information and molecular weight of the chains produced from sacrificial initiator allow the molecular weight of the p(MAA-co-DEAEMA) polyampholytes to be determined. This assumes that the chains recovered from solution are appropriate proxies of the brushes and that all tert-butyl groups

“protecting” groups are cleaved during conversion of p(tBMA-co-DEAEMA) to form the p(MAA-co-DEAEMA) polyampholyte brushes.

As shown by Murata and Rühe,1 when a polymer brush is chemically modified and if assumptions of constant mass density and grafting density are invoked, then ratio of brush heights (before and after chemical modification) equals the ratio of molecular weights of the parent and modified chains comprising the brushes:

퐻 푀 1 = 푛,1 퐻 푀 2 푛,2 (S1)

Recasting this equation for the p(tBMA-co-DEAEMA) “parent” and corresponding p(MAA-co-DEAEMA) polyampholyte (PA) system, and using the assumption that the conversion from “protected” form to polyampholyte is complete results in

퐻 푀푛, polyampholyte ∑ 푛 푚 1 = = 푖 표,푖 퐻2 푀푛,"parent" 푀푛,"parent" (S2)

172 where ni is the number of electrolytic repeating units of type i (i = MAA, DEAEMA) comprising the polyampholyte and mo,i their molar mass. The numbers of repeating units of each type are readily determined from the composition and molecular weight of the parent brush, which is represented by the chains in solution.

Mol % Mn Mn, PA

tBMA:DEAEMA ntBMA nDEAEMA DEAEMAa (kDa)a (kDa)b

80:20 19 55.8 318 57 37.9

50:50 48 68.4 250 177 54.3

20:80 83 55.3 66 248 51.6

a Composition and molecular weight determined from free chains in solution.

b Calculated based on numbers of repeating units.

173

Buffer Recipes and Preparation Procedure

Buffers were prepared following literature procedures using the following acids and salt.2

Table B-1. Amounts of buffer species and sodium chloride required to make 100 ml buffer solution of 30 mM ionic strength at pH values given.

Mass of Buffer Mass of NaCl pH Buffer Species Species (g) (g)

3 Phosphoric Acid 0.294 0.018

5 Acetic Acid 0.180 0.057

7 MOPSa 0.627 0.111

8 TAPSb 0.729 0.128

a 3-(N-morpholino)propanesulfonic acid

b N-[Tris(hydroxymethyl)methyl]-3-aminopropanesulfonic acid

174

60% DEAEMA; and G and H are for a copolymer targeted to be 30% tBMA and

70% DEAEMA. All reactions were run in anisole at T=35°C for 8 h.

175

Figure B-2. Swelling response of p(DEAEMA-co-MAA) brushes as a function of pH. The curves also display the refractive index from the “slab” model used in the analysis of the ellipsometric data.

176

[tBMA]:[DEAEMA] (left) and [MAA]:[DEAEMA] (right). RMS roughness values are inset in each image.

177

Vita

Rachel Irene Ramirez was born in Winter Park, Florida and raised in

Fayetteville, Georgia. Rachel attended Whitewater High School where she graduated in 2009. Upon graduation in 2009, Rachel enrolled at the University of

North Georgia in Dahlonega, Georgia where she earned a B.S. in Chemistry. In

2013, Rachel enrolled at the University of Tennessee-Knoxville to major in polymer chemistry. She joined the research group of Professor S. Michael Kilbey II, where her research involved how the structural design of polymer thin films impacts their response.

178