Anisotropic Organic Materials - ACS Publications - American Chemical

C h a p t e r 18

Order-Disorder and Order-Order Transitions in Smectic C* Liquid Crystalline Diblock Copolymers

Downloaded by NORTH CAROLINA STATE UNIV on September 17, 2012 | http://pubs.acs.org Publication Date: November 2, 2001 | doi: 10.1021/bk-2001-0798.ch018

Mitchell Anthamatten and Paula T. Hammond Department of Chemical Engineering, Massachusetts Institute of Technology, Cambridge, M A 02139

B l o c k copolymers containing side-chain l i q u i d crystalline (LC) moieties enable materials to have a combination o f p o l y m e r i c and l i q u i d crystalline properties. I f the two blocks are incompatible, microphase segregation can occur resulting i n isolated LC and amorphous domains. W e are interested i n h o w LC o r d e r i n g overlaps w i t h b l o c k c o p o l y m e r microstructure. A series o f diblocks w i t h an amorphous, polystyrene block and a methacrylate based side-chain LC block were prepared. R o o m temperature and morphological phase diagrams were evaluated using a combination o f SAXS and TEM, and microstructure transitions were identified using SAXS at elevated temperatures. These results are compared to LC transitions studied w i t h microscopy and calorimetry. The LC phase effectively changes the F l o r y - H u g g i n s interaction parameter. Effort is also b e i n g made to m o d e l the observed transitions and to optimize electro-optic responses.

Thermotropic liquid crystals are rod-like organic molecules that form an ordered fluid w i t h anisotropic optical and electro-magnetic properties. O f particular interest are chiral smectic C * liquid crystals w h i c h arrange into layers and have a net spontaneous p o l a r i z a t i o n J These materials are ferroelectric, and an electric field coupled to the spontaneous polarization can be used to switch the material between two stable states. F o l l o w i n g their discovery, ferroelectric l i q u i d crystals received widespread attention because o f their m i c r o s e c o n d response times and bistable s w i t c h i n g capabilities. M a n y applications can be envisioned i n c l u d i n g light valves, optical switches, and m e m o r y storage devices. H o w e v e r , to realize any application, the h e l i c a l twist associated w i t h the smectic C * phase must be u n w o u n d ; this creates engineering challenges such as maintaining a constant c e l l thickness and c o n t r o l l i n g mesogen orientation across the electo-optic cell. P o l y m e r liquid crystals ( P L C ' s ) are formed by incorporating l i q u i d crystals into the polymer backbone or as a side-chain moiety.^ Their thermotropic phase behavior depends not only on mesogen characteristics such as shape and polarity, but also the f l e x i b i l i t y , length, and tacticity o f the p o l y m e r backbone. M e s o g e n i c ordering is

© 2002 American Chemical Society

In Anisotropic Organic Materials; Glaser, R., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 2001.

239

240 3

affected strongly by interactions w i t h the polymer b a c k b o n e . Generally, side-chain P L C ' s have broader smectic phases than their small molecule analogs. D i b l o c k copolymers are formed i f two p o l y m e r chains A and Β are covalently j o i n e d at their ends. I f A and Β are dissimilar enough, and i f the total molecular weight is large enough, microphase segregation occurs resulting i n a microstructure on the scale o f the radius o f gyration o f the polymer. The extent o f phase segregation can be quantified b y the product χ Ν where χ is the F l o r y - H u g g i n ' s interaction parameter, describing the enthalpic penalty o f m i x i n g , and Ν is the copolymer degree o f polymerization ( N = N + N ) , reflecting the entropie drive o f an Ν dimensional c h a i n s U p o n microphase segregation, periodic domains rich i n A and Β form a selfassembled morphology, and, as Figure 1 suggests, the type o f microstructure depends largely on the composition o f the diblock.


A

B

Figure 1. Cartoons illustrating the molecular architecture of LC diblock (top) and possible phase-segregated microstructures (bottom)

copolymers

This paper relates to diblock copolymers w i t h a high T amorphous b l o c k and a more flexible, side-chain ferroelectric liquid crystalline ( L C ) block. This architecture a l l o w s a self-supportable L C material, e.g. a glassy s o l i d w i t h electro-optical properties. The idea is to combine the mechanical integrity associated w i t h the P S block w i t h the electro-optic properties arising from the L C block. A d d i t i o n a l l y , it has been s h o w n that the b l o c k c o p o l y m e r intermaterial d i v i d i n g surface ( I M D S ) is capable o f stabilizing l i q u i d crystalline side-chain moieties m u c h as a buffed electrooptic cell stabilizes small molecule L C ' s A ^ W i t h sufficient control over the type and orientation o f microdomain morphology it should be possible to achieve macroscopic spontaneous polarization over large domains. A number o f research groups, including ours, have already taken significant steps toward developing this technology.^" ^ g


241 P r e v i o u s l y we reported the synthesis and characterization o f a series o f w e l l defined smectic C * d i b l o c k c o p o l y m e r s . ^ R o o m temperature morphologies and mesomorphic superstructures i n solvent cast, roll-cast, and fiber drawn samples were also reported. * % The objective o f this paper is to probe h o w thermotropic L C order transitions affect diblock copolymer morphology. The approach is to systematically measure and compare L C transition temperatures to m o r p h o l o g i c a l order-order transition ( O O T ) and order-disorder transition ( O D T ) temperatures.


Materials The molecular structure o f the diblock copolymers considered here is shown in Figure 2. Synthesis f o l l o w e d the literature procedure that i n v o l v e s direct a n i o n i c p o l y m e r i z a t i o n o f styrene f o l l o w e d by a d d i t i o n o f a mesogenic methacrylate monomer. 14 T h i s technique enables control o f both the total molecular weight and the L C volume fraction § w h i c h determine the phase segregated microstructure as s h o w n i n F i g u r e 3. F o r φιχ less than ~ 0.25, dispersed morphologies ( D S ) are observed that consist o f unaligned L C spheres i n a continuous P S d o m a i n . A t intermediate v o l u m e fractions the d i b l o c k s form layered mophologies that can be s u b d i v i d e d into c o m p l e t e l y l a m e l l a r ( L ) a n d p r e d o m i n a t e l y l a m e l l a r ( P L ) m o r p h o l o g i e s . T h e P L m o r p h o l o g y consists o f lamellae that coexist w i t h P S cylinders w h i c h are arranged either hexagonally ( P L / H C P ) or i n a m o d i f i e d layer ( P L / M L ) fashion. Cylinders i n the M L form retain periodic spacing between planes o f cylinders but not necessarily between cylinders themselves. F i n a l l y , at high φιχ, only hexagonally-close-packed P S cylinders ( H C P ) self-assemble i n a continuous L C phase. N o t e the phase diagram is asymmetric i n that microstructures w i t h continuous L C phases occupy a m u c h larger part o f phase space than do structures w i t h a continuous P S phase. L C

P S

C4H9-

Figure 2. Molecular structure materials in this study.

of PS-HBPB

diblock copolymers;

n = 6 for all



242

Figure 3. Experimentally determined morphological phase diagram for PS-PHBPB diblock copolymers. Morphologies observed include: dispersed LC spheres (DSu), lamellae (L), predominately lamellae (PL), and hexagonally-packed PS cylinders

(HPCrsf

Experimental L i q u i d crystalline phase transition temperatures were assigned u s i n g differential scanning calorimetry ( D S C ) and p o l a r i z e d optical m i c r o s c o p y ( P O M ) . ^ A P e r k i n E l m e r D S C - 7 w i t h a c o o l i n g accessory was used at scanning rates o f 20 °C / m i n . M i c r o s c o p y observations were made using a L e i t z optical microscope equipped w i t h a C C D camera and a Metier F P - 8 2 heating stage and controller (heating/ cooling rate o f 10 °C / m i n ) . L C phases were assigned by e x a m i n i n g textures after samples were annealed for hours just beneath their phase transition temperature. S m a l l angle X - r a y scattering ( S A X S ) was used to confirm phase assignments and to measure block copolymer order-disorder transition and order-order transition ( O D T & O O T ) temperatures. B e f o r e X - r a y analysis, i n order to encourage phase segregation, a l l samples were annealed i n a vacuum oven for 48 hours at 110 °C, just above the T o f the P S block. The X - r a y sample chamber was evacuated, and samples were suspended vertically, using p o l y i m i d e adhesive tape, i n an Instec temperature controlled hotstage (model: H S 2 5 0 ) . X - r a y s were generated from a rotating copper g



243 anode producing C u Κ α radiation ( λ = 1.54 Â ) and operating at 40 k V χ 30 m A . X rays were scattered on to a detector, 54.1 c m from the sample, that consists o f a pressurized x e n o n chamber w i t h a w i r e g r i d assembly (512 χ 512). D a t a were acquired as scattered X - r a y intensity, I, taken as a function o f the scattering vector q = 4 π / ( λ sin Θ) and temperature. T o a v o i d sample degradation, shorter scattering times were used for high-temperature scans. Preceding data analysis, a l l intensities were n o r m a l i z e d to one hour to a l l o w direct comparison between temperatures. H o w e v e r , due to differences i n sample thickness, direct c o m p a r i s o n o f intensity between samples o f different composition is not appropriate. O D T and O O T transitions were determined b y plotting both inverse m a x i m u m intensity 1/Imax and the wavelength o f concentration fluctuations D vs. inverse temperature T" . In the disordered state, 1/Imax is predicted to decrease linearly with T" and D should be nearly independent o f T" ; the theory behind this analysis is addressed i n reference 19. Other observations were also considered when assigning O D T ' s , such as decreases i n f i l m v i s c o s i t y and the disappearance o f higher-order S A X S peaks. 1

1

Transmission electron micrographs were obtained using a J O E L 2 0 0 C X electron microscope operating at 160kV. Preceding microscopy, samples were cut into 40 n m sections using a Reichert-Jung F C 4 E microtome and were then stained w i t h R u 0 4 vapor for 20 minutes.

Results & Discussion S i x L C d i b l o c k copolymers representing different regions o f the phase diagram were studied at elevated temperatures. Table I is a summary o f the material characteristics and L C phase transitions—systematic synthesis & characterization o f L C phases is described i n elsewhere.^ N o t i c e that for samples w i t h intermediate to h i g h L C v o l u m e fractions, P S - H B P B 5 0 , 56, & 58, that the smectic mesophases appear destabilized. These samples have isotropization temperatures ( T ) as l o w as 159 °C and o n l y exhibit the smectic C * mesophase. F o r samples w i t h l o w e r L C fractions t y p i c a l l y have smectic phases that are stable up to 180 - 200 ° C . ^ The destabilization o f smectic phases in P S - H B P B 5 0 , 56, & 58 results from an entropie frustration arising from conformationally asymmetric diblocks situated at a lamellar interface. T h i s idea w i l l be discussed in more detail later; however, note that these three sample are the o n l y ones that exhibit the P L morphology. Conversely, P S - H B P B 7 9 , w i t h a higher L C content ( φ ι χ = 0.77), has a very stable smectic phase and forms a H P C microstructure. The temperature dependent S A X S scattering patterns for P S - H B P B 3 2 are shown in F i g u r e 4. T h e smaller angle reflections at q ~ 0.34 nm" correspond to b l o c k copolymer m i c r o d o m a i n ordering o f about 186 Â . Higher-order peaks are difficult to discern since o n l y two-hour scans were acquired at each temperature. The r o o m temperature m o r p h o l o g y o f these samples was confirmed i n an earlier study using m u c h longer scans. ^ Weaker signals, not shown i n the figure, are present at w i d e r angles (q ~ 2.2 nm' ) and arise from smectic ordering. These peaks appeared broader than expected for polymer s m e c t i c s ^ ' ^ i , however, upon prolonged annealing, they became narrower and more defined. A d d i t i o n a l l y , some broadening i n the w i d e r angle regime arises from equipment limitations. P S - H B P B 4 3 had nearly identical thermal behavior and thus is not shown. iso

1

1


244 Table I. Material Characteristics and L C phase transitions


Diblock

Block M„ (kg/mol)

PSHBPB32 PSHBPB43 PSHBPB50* PSHBPB56* PSHBPB58* PSHBPB79*

PS

LC

17.7

7.3

Total M„ MJM (kg/mol)

n

24.6

1.07

LC Phase

RT Morph.

Transitions

(°Q

H : S *(163)S (181)I

LAM

C

A

C : I(171)S (156)S * A

10.3

7.4

17.7

1.08

LAM

8.4

11.0

19.4

1.14

PL/ML

10.5

11.0

21.5

1.07

PL/HPC

8.8

12.0

20.8

1.11

PL/ML

4.8

16.3

21.1

1.08

HPC

C

H : S *(136)S (158)Ch(177)I c

A

C : I(166)S * C

H : S *(161)I C

C : I(145)S * C

H : S *(159)I C

C : I(153)S * C

H : S *(177)I C

C : I(171)S * H : S *(176)S (212)I C : I(203)S (173)S * C

C

A

A

r

B o t h P S - H B P B 3 2 & 43 have O D T temperatures that are slightly higher than their isotropization temperature T . This correlation indicates that the isotropization o f the L C d o m a i n triggers the O D T transition. T h i s phenomenon can be explained by slightly m o d i f y i n g the criterion for phase segregation to account for L C order as w e l l as segmental interactions. F o r amorphous-amorphous d i b l o c k copolymers, phase segregation occurs i f the product χ Ν exceeds a critical value. F o r the L C d i b l o c k s considered here, there is an additional increment o f free energy % c N i n v o l v e d w i t h m i x i n g ordered mesogens and disordered P S chain segments. Thus, as suggested i n our earlier p a p e r s ' ^ phase segregation occurs w h e n (% c + % ) N exceeds a critical value. The quantity (%LC + %) represents the thermodynamic incompatibility between the two b l o c k s . F o r P S - H B P B 3 2 & 43, the % c term disappears as the sample is heated through the L C isotropization point, thus reducing % ffN b e l o w its c r i t i c a l value. In a study o f symmetrical P S - L C diblocks (φ ~ 0.5) Y a m a d a et. al. observed a similar phenomena where for lower molecular weight samples (< 10 k g / m o l ) , T D T was coincident with T i . ^ P S - H B P B 5 0 , 56, and 58 a l l exhibit P L microphase segregated morphologies at r o o m temperature. * ° P S - H B P B 5 8 was unique since it showed easily discernible scattering peaks for both lamellae and the c y l i n d r i c a l defects in the shorter (2 hr) scans used at elevated temperatures. The S A X S profiles for P S - H B P B 5 8 change significantly as temperature is increased as shown in Figure 5. The morphology at lower temperatures is predominately lamellae w i t h c y l i n d r i c a l defects arranged i n modified layer fashion ( P L / M L ) . A t lower temperatures the first-order B r a g g peak at q « 0.39 nm" corresponds to a lamellar spacing o f 156 Â , and the less intense peak q « 0.50 nm" (123 Â ) signifies the layers o f M L cylinders. B o t h peaks have higherorder signals that correspond to a scattering ratio o f 1:2:3. Smectic layer ordering is apparent from B r a g g scattering at wider angles ( q ~ 2.1 nm" ). Between 160 and 180 °C this reflection disappears w h i c h is consistent with S * to isotropic transition at 177 iso

L

6

L

L

e

0

S 0

1

1

1

c


245

ODT : 1 8 0 ± 1 0 ° C


T

iso

=181 °C

1 ^ 1 9 0 °C R l 8 0 °C 160 °C 140 °C 120 °C 100 °C 60 °C q

1

(nm' )

Figure 4. SAXS ID diffraction patterns at elevated temperatures for PS-HBPB32. Data are plotted as the logarithm of relative scattered intensity log I(q) vs. the scattering factor q.

ODT

195 ± 5 °C 170±5°C = 177 ° C

t

210°C 200 °C •190°C »180 °C h*170 °C »160 °C •140 °C ^ 1 2 0 °C * 1 0 0 °C Γ*»60 °C

1

q (nm )

Figure 5. SAXS ID diffraction patterns at elevated temperatures for PS-HBPB58. Data are plotted as the logarithm of relative scattered intensity log I(q) vs. the scattering factor q.


246 °C observed using P O M . Between 120 °C and 160 °C, the first-order peaks gradually m o v e closer together until they are completely merged at 180 °C, where, based on analysis o f S A X S intensities, the m o r p h o l o g y is strongly phase-segregated. At temperatures exceeding 190 ° C , the intensity o f the r e m a i n i n g peak decreases, i n d i c a t i v e o f an O D T . U p o n c o o l i n g , t w o separate sets o f peaks reappeared, indicating that this transition is reversible; this has been explained i n detail i n a recent publication.


22

T o confirm P S - H B P B 5 8 ' s morphology, a sample was annealed at 110 °C for 72 hours. A s discussed i n reference 2 2 , resulting T E M micrographs c o m b i n e d w i t h S A X S diffraction patterns confirm that alternating lamellae are observed coexisting w i t h P S cylinders. T o examine the m o r p h o l o g y at the higher-temperature state, a S A X S specimen was q u i c k l y cooled from 210 °C to r o o m temperature, and another T E M was acquired; this time a completely lamellar morphology was observed. Based on these and other e x p e r i m e n t s , the sample undergoes a thermoreversible O O T , between 160-170 ° C , from a predominately lamellar m o r p h o l o g y to a completely lamellar one. This type o f transition was an unexpected result; based on the theory o f amorphous-amorphous b l o c k copolymers, one w o u l d expect the opposite: an order t r a n s i t i o n u p o n h e a t i n g f r o m a l a m e l l a r m o r p h o l o g y to a c y l i n d r i c a l morphology. 19,24-26 22

Figure 6 is a cartoon that explains h o w liquid-crystalline order triggers the O O T in P S - H B P B 5 8 ( P L / M L -> L ) . W h e n the L C phase is ordered, the p o l y m e r molecule's natural shape is 'wedge-shaped', and it is conformationally asymmetric i n the x - y plane. Consequently, it natural for this shape to self-assemble into a structure w i t h a curved I M D S ^ > 8 w i t h the L C domain on the convex side. H o w e v e r , the d i b l o c k ' s v o l u m e f r a c t i o n (φ ~ 0.56) is more c o m p a t i b l e w i t h a l a m e l l a r microstructure. U p o n heating, as the L C order disappears, the d i b l o c k is less asymmetric and becomes more suited for a lamellar interface. In summary, i f conformationally asymmetric diblock are situated at a planar interface the L C chains become entropically frustrated. This explains 1) the existence o f a near-equilibrium microstructure w i t h coexisting lamellae and P S cylinders 2) the destabilization o f the smectic phases observed i n P S - H B P B 5 0 , 5 6 , & 58, and 3) the O O T that occurs in P S HBPB58. 2

2

Free Energy Description of Side-Chain L C Diblock Copolymers Theories for diblock copolymer phase segregation are n o w advanced and agree w e l l w i t h experiment. F o r a recent review o f this topic the reader is referred to H a m l e y . 4 Theories accurately predict microphase structures as a function o f diblock copolymer composition φ and the product χ Ν . A classic example is Semenov's meanfield theory for h i g h l y incompatible b l o c k s ^ ; it predicts m i c r o d o m a i n structure by balancing interfacial free energies between the two blocks with elastic free energies associated w i t h stretching polymer chains away from the I M D S . The Semenov theory enables the O D T phase boundary to be estimated however it predicts the phase boundaries between morphologies to be independent o f χ Ν . Figure 7a shows a simplified phase diagram for amorphous-amorphous b l o c k copolymers. I f one b l o c k contains, as part o f its m o n o m e r repeat unit, a l i q u i d crystalline moiety, then one w o u l d expect the morphological phase diagram to change 2



247

Figure 6. Top: order-order transition observed in PS-HBPB58 from a predominately lamellar morphology to a completely lamellar one. Three different Bragg periods: Dl (156 À), D2 (123 Â), & D3 (145 Â) were observed in SAXS (see Figure 5). Bottom: cartoon showing how LC smectic order can lead to conformationally asymmetric diblocks in the x-y plane—upon isotropization, conformational asymmetry is reduced.

significantly. Figures 7b and 7c are projections o f phase diagrams for main-chain and side-chain liquid crystalline diblock copolymers. The upper regions o f Figures 7b and 7c represent l o w e r , anisotropic states and are a s y m m e t r i c about φ ^ 0.5. E x p e r i m e n t a l l y , this is indeed the case; we recently confirmed that m o r p h o l o g i c a l phase boundaries as w e l l as O D T temperatures are strongly affected by L C ordering.^ The lamellar samples P S - H B P B 3 2 & 43 are cases i n w h i c h O D T ' s are initiated b y i s o t r o p i z a t i o n o f the L C phase. O t h e r researchers have o b s e r v e d s i m i l a r p h e n o m e n a . ^ H o w e v e r , it is less frequent for transitions i n L C order to trigger an O O T . 2 9 i P S - H B P B 5 8 the sample's morphology transforms, upon heating, from a P L / M L morphology to a completely lamellar one; this transition is unusual because curvature is lost upon heating; as Figure 7a indicates, this transition is counter to theoretical predictions for amorphous-amorphous block copolymers. Theoretical studies have focused on diblocks w i t h mesogens incorporated tandem w i t h i n the L C m a i n chain. Consequently, the extension o f the L C backbone perpendicular to the I M D S is facilitated below the isotropization temperature by L C ordering. This results i n a bias o f morphological phases toward microstructures w i t h the L C phase r e s i d i n g i n m i c e l l e e o r e s ^ ^ l as s h o w n i n F i g u r e 7b. T o our =

n


248 knowledge, there have been no theoretical phase diagram models formulated for L C d i b l o c k s that have an amorphous b l o c k and a side-chain L C b l o c k . A n extensive analysis o f order in nematic side-chain homopolymers was carried out by W a n g and W a r n e r . ^ T h i s m o d e l i n v o l v e d self-consistent calculation o f the mesogenic and L C backbone order parameters, w h i c h i n turn were used to predict backbone c h a i n configurations, both parallel and perpendicular to mesogenic ordering. W e are seeking a free energy expression for side-chain L C d i b l o c k copolymers that captures k e y features o f the m o r p h o l o g i c a l phase diagram such as L C induced O D T ' s and O O T ' s . F o u r free energy components are considered: FTOT = F c + F M + FIMDS + F M,LC where F OT is the total free energy per d i b l o c k copolymer, F c is the total free energy associated w i t h the L C b l o c k ( i n c l u d i n g stretching), F M is due to entropie stretching o f the amorphous block, FIMDS is a surface tension term, and F M,LC represents F l o r y - H u g g i n s type interactions between the L C and amorphous repeat units. FLC is described u s i n g W a n g and Warner's m o d e l ^ and includes elastic stretching o f the L C backbone. Since this model was intended for L C homopolymers, FLC is independent o f diblock copolymer morphology. Calculation o f the amorphous stretching energy is s i m i l a r to that o f G i d o * * where the amorphous b l o c k is considered a p o l y m e r brush grafted to a surface with mean and Gaussian curvatures H and K. A 'wedge-like' geometry is assumed subjecting both domains to the v o l u m e constraint: L


A

T

A

L

A

A

2

-Khf+Hhf+k 3— _

:

=

(1)

σ

W h e r e h is the brush height and σ is the graft density at the I M D S . The brush energy per c h a i n is calculated by progressively adding chains to a surface. B y differentiating eqn 1 da/dh can be calculated and the stretching free energy obtained using: * AM *h j a

AM

F

AM

= - \ d & s(a) = - L f dti ^ s ( h σ i σ I dh

)

(2)

T h i s equation is exact for lamellae and for grafting to a concave side o f the I M D S ; it is a g o o d a p p r o x i m a t i o n for c o n v e x surfaces w i t h l o w e r curvatures. H o w e v e r , equation 2 requires an expression for /*AM. This is obtained by a p p l y i n g the wedge-like v o l u m e constraint (eqn. 1) to each phase, equating the two resulting graft densities, and s o l v i n g for h as a function o f Η, K, and h . T h e third term, the interfacial surface energy Fj„t, is the interfacial tension γ d i v i d e d b y the graft density σ . F o r the case o f a homogeneous melt Fj is assumed zero. L a s t l y , FAM,LC describes the F l o r y - H u g g i n s interactions between amorphous and L C m o n o m e r s : AM

LC

nt

FAM,LC = %NLCAM>.

In s u m m a r y , a free energy d e s c r i p t i o n has been o u t l i n e d to r e c o n c i l e our experimental observation o f L C triggered O D T and O O T transitions. B y m i n i m i z i n g the F at a g i v e n temperature, the e q u i l i b r i u m m o r p h o l o g y for a p r e - s p e c i f l e d architecture can be determined. B y iteration over a temperature and c o m p o s i t i o n range, phase diagrams r e s e m b l i n g those i n Figures 7b & 7c s h o u l d be predicted. O n g o i n g w o r k includes refinement and validation o f this model. T



249

Figure 7. Morphological phase space predictions for a) amorphous-amorphous, b) (main-chain LC)-amorphous, and c) (side-chain LC) - amorphous diblock copolymers. Regions include LC spheres (SLC) and cylinders (C c), lamellae (L), amorphous cylinders (C ) and spheres (S ). L

A

A


250


Conclusions A series o f L C side-chain diblock copolymers were investigated The data presented indicate that L C isotropization can induce O D T ' s for the l o w e r molecular weight samples. P S - H B P B 5 8 exhibits a thermoreversible O O T that progresses, upon heating, from a predominately lamellar morphology to a completely lamellar one. C o n f o r m a t i o n a l asymmetry explains 1) the destabilization o f smectic phases i n samples w i t h φιχ ~ 0.50 - 0.58, 2) the presence o f curvature at unusally l o w L C v o l u m e fractions, and 3) the loss o f curvature i n P S - H B P B upon heating through its O O T . Efforts are currently underway to describe the interplay between mesogenic and morphological order by using a free energy modeling approach.

Acknowledgements The research was supported by the N a t i o n a l Sciece Foundation P o l y m e r Program for funding under Grant N u m b e r D M R - 9 5 2 6 3 9 4 . W e are also thankful to Professor N e d Thomas for use o f his microtome equipment.

Literature Cited 1) 2) 3) 4) 5) 6) 7) 8) 9) 10) 11) 12) 13) 14) 15)

Clark, Ν . Α . ; L a g e r w a l l , S. T. Appl. Phys. Lett. 1980, 36, 899-901. B l a c k w o o d , Κ. M. Sciencs 1996, 273, 909-912. W a n g , X. J.; Warner, M. J. Phys. A: Math Gen. 1987, 20, 713-731. H a m l e y , I. W . The Physics of Block Copolymers; O x f o r d U n i v e r s i t y Press: N e w Y o r k , 1998. M a o , G.; W a n g , J . ; C l i n g m a n , S. R . ; Ober, K.; C h e n , J . T . ; T h o m a s , E . L. Macromolecules 1997, 30, 2556-2567. Zheng, W . Y.; H a m m o n d , P. T. Macromolecules 1998, 31, 711 -721. R u o k o l a i n e n , J . ; Saariaho, M.; I k k a l a , O . ; ten B r i n k e , G . ; T h o m a s , E. L.; Serimaa, R . ; T o r k k e l i , M. Macromolecules 1999, 32, 1152-1158. C h i e l l i n i , E.; G a l l o , G.; A n g e l o n i , A. S.; L a u s , M.; B i g n o z z i , C . Macromol Symp. 1994, 77, 349-358. A d a m s , J . ; G r o n s k i , W . Makromol. Chem., Rapid Commun. 1989, 10, 553-557. Z a s c h k e , B.; Frank, W . ; Fischer, H.; Schmutzler, K.; A r n o l d , M. Polymer Bulletin 1991, 27, 1-8. Bohnert, R . ; Finkelmann, H. Macromol. Chem. Phys. 1994, 195, 689-700. Y a m a d a , M.; Iguchi, T.; Hirao, Α . ; Nakahama, S.; Watanabe, J . Macromolecules 1995, 28, 50-58. Fischer, H.; Poser, S. Acta Polymer. 1996, 47, 413-428. Zheng, W. Y.; H a m m o n d , P. T. Macromol. Rapid Commun. 1996, 17, 813-824. Omenat, Α . ; H i k m e t , R. A. M.; L u b , J.; van der Sluis, P . Macromolecules 1996, 29, 6730-6736.



251 16) M a o , G.; W a n g , J . ; Ober, C. K.; Brehmer, M.; O ' R o u r k e , M. J.; Thomas, E . L. Chem. Mater. 1998, 10, 1538-1545. 17) Brehmer, M.; M a o , G.; Ober, C. K. Macromol. Symp. 1997, 117, 175-179. 18) Anthamatten, M.; Z h e n g , W . Y.; H a m m o n d , P . T. Macromolecules 1999, 32, 4838-4848. 19) Leibler, L. Macromolecules 1980, 13, 1602-1617. 20) Lipatov, Y. S.; Tsukruk, V. V.; Shilov, V. V. Rev. Macromol. Chem. Phys. 1984, C24, 173-238. 21) Shilov, V. V.; Tsukruk, V. V.; Lipatov, Y. S. Journal of Polymer Science 1984, 22, 41-47. 22) Anthamatten, M.; H a m m o n d , P. T. Macromolecules 2000, 32, 8066-8076. 23) Y a m a d a , M.; Iguchi, T.; Hirao, Α . ; Nakahama, S.; Watanabe, J. Polymer J. 1998, 30, 23-30. 24) Fredrickson, G. H.; Helfand, E. J. Chem. Phys. 1987, 87, 697-705. 25) Semenov, A. N. Sov. Phys. JETP 1985, 61, 733-742. 26) Lescanec, R . L.; Muthukumar, M. Macromolecules 1993, 26, 3908-3916. 27) M i l n e r , S. T. Macromolecules 1994, 27, 2333-2335. 28) G i d o , S. P.; Z.G., W . Macromolecules 1997, 30, 1997. 29) Saenger, J . ; G r o n s k i , W . ; M a a s , S.; Stuehn, B.; H e c k , B. Macromolecules 1997, 30, 6783-6787. 30) Sones, R . Α . ; Petschek, R . G. Physical Review Ε 1994, 50, 2906-2912. 31) W i l l i a m s , D . R . M.; Halperin, A. Phys. Rev. Letters 1993, 71, 1557-1560.


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