‘v 15: x: .3. . ‘cc V- ~ ue‘ ‘~t4-;p ' “. 1 if? ‘ II I ' n. I I .w. K. .‘x 'v" I 9 1 ' -"’ Mam 3 x .4» . x .fi‘fig‘g ’1“ e fax , “1%- : 12" -l 'k‘i. “‘r 11%: ~ 3-51 u [5, ' kiwi?- ,{‘ I A x L. “52%"? $5, r; '” I; . “fife; 1 . ‘ 3’1‘1,‘ ‘ Ln nu , . I ~ ~ i'é ’3 2: "1‘-3’-:?:‘ wfififi “‘F'ff" ' Y. ‘1 5‘ This is to certify that the dissertation entitled THE EFFECT OF INTERPHASE PROPERTIES ON ADHESION IN POLYPHENYLENE SULFIDE/ CARBON FIBER COMPOSITES presented by Gregory Scott Fisher has been accepted towards fulfillment of the requirements for . Ph . D . degree in Chemical Engineering MS U i: an Affirmative Action/Equal Opportunity Institution 0-12771 IGAN STATE ‘ will l r121: m tilllllllllllll 93 01027 6149 . LIBRARY M'Chigan State University *- PLACE N RETURN BOXtonmavotNochockmmm younocord. TO AVOID FINES Mum on or baton date duo. DATE DUE DATE DUE DATE DUE fiClfi éfl fifiE—j MSU IoAnNflrmutlvo Adlai/Emu Opportunly Im Mani-a1 THE EFFECT OF INTERPHASE PROPERTIES ON ADHESION IN POLYPHENYLENE SULFIDE/CARBON FIBER COMPOSITES By Gregory Scott Fisher A DISSERTATION Submitted to Michigan State University in partial fulfillment of the requirements for the degree of DOCTOR OF PHILOSOPHY Department of Chemical Engineering ABSTRACT THE EFFECT OF INTERPHASE PROPERTIES ON ADHESION IN POLYPHENYLENE SULFIDE/CARBON FIBER COMPOSITES By Gregory Scott Fisher The level of adhesion between reinforcing fibers and a thermoplastic matrix can have a strong effect on composite mechanical properties. Several factors can affect the adhesion between carbon fibers and a semicrystalline thermoplastic matrix. These factors include fiber surface structure, morphology of . the matrix near the fibers, chemical bonding between fiber and matrix, and surface adsorption of matrix onto fibers. The objectives of this study were to determine the interfacial shear strengths of several types of carbon fibers in a new composite grade of polyphenylene sulfide (PPS), an extrusion grade of PPS, and an amine terminated PPS and to determine how the factors mentioned above affect fiber-matrix adhesion. In this work, the surface free energies of carbon fibers and PPS resins as they relate to surface chemistry and adhesion as well as the effect of composite processing conditions on carbon fiber surface chemistry are discussed. The crystallinity of PPS resin is described. Finally, the interfacial shear strengths between several types of carbon fibers and both a composite grade, an extrusion grade, and an amine terminated version of PPS resin are presented and compared in light of the interphase characteristics previously described. The composite grade of PPS established stronger adhesion to A84 carbon fibers than either the extrusion grade or amine terminated PPS. All of the resins have similar surface free energies, eliminating differences in wetting behavior as the reason for the difference. No chemical reaction between fiber surfaces and PPS resins occurs at processing conditions. An amine end group on PPS does not increase the level of adhesion to carbon fibers. The key to the higher interfacial shear strength for the composite grade PPS is the formation of a tough interphase with fine crystalline spherulites (0.3-0.5 p in diameter). Both the amine terminated PPS and extrusion grade PPS formed transcrystalline interphases with high modulus HMS4 fibers. The transcrystallinity was determined to be a result of epitaxy from surface graphitic basal plane edges present on these fibers. The interfacial shear strength of transcrystalline composites has a strong dependence on the isothermal crystallization temperature. ACKNOWLEDGMENTS This work was sponsored by the Phillips Petroleum Company, which provided materials and funding. Additional support was provided by the Office of Naval Research. iv TABLE OF CONTENTS List of Tables .............................. n ' fleer ............. n' frnFi ............. viii 10 MIL Materials and Experimental Methees ......... 37 Re in t ri l ............................... 37 We: ................................... 39 Mel rf Fr En Measur m n .............. 42 r n i r F En M m n ........... 47 Mi i Examina ' n f llization f Thin il f P ................................... 49 WW ...................... 50 E hin fP ml fr nnin El When ............................ 51 IMI r E m n eLEESMaterials ................................. 51 l b m1 I v: f wn h r n 131an ........................................ 52 rf i l h r n h M m n ................. 53 EM ........................................ 57 W ........................ 57 W228 .................... 57 WW .................. 59 vi WWW n P n r n ................ 64 W ........................ 64 II' in fN in ..................... 64 WW ....................... 71 lliza'n f R inin h n f r n [Eben ........................................ 76 hl n h n f r nFI ......... 78 N L I N ................................. 86 WW ........... 88 W ........................ 88 l v h n I rf ....... 88 l h n h P r n I m I 91 W ................................. 100 W ..................... 102 APPENDIX WWW QanhIehens .................................... 107 vii LIST OF TABLES m1 Tensile and Shear Modulus Values for Neat PPS Resins ..................................... 39 M Carbon Fiber Diameters and Tensile Modulus Values ........................................ 41 Table 2,3 Contact Angle Liquids and Their Surface Free Energies ................................... 49 m Surface Free Energies for PPS Resins ............. 58 W HMS4 Contact Angles and Work of Adhesion ...................................... 61 W Surface Free Energies of Carbon Fibers ............. 62 13121944,], Relative PPS Crystallite Nucleating Activity of Carbon Fibers ................................. 77 W Interphase Morphologies for Fiber Reinforced PPS Samples .................................... 85 m XPS Results for AS4, HMS4, and Pitch Fibers ........ 89 M XPS Results for 1M6 Fibers with 0%, 100%, and 600% Surface Treatment ............................. 90 viii LIST OF FIGURES Figure 1.1 Schematic of a Pendant Drop .................. 12 Figure 1.2 Possible Interphase Morphologies ............... 17 Figure 2.1 Load Versus Displacement for Neat PR10X2 Resin ......................................... 40 Figure 2.2 Pendant Drop Environmental Cell Schematic ........ 43 Figure 2.3 Pendant Drop Surface Free Energy Measurement Apparatus ............................. 45 Figure 2.4 Pendant Drop Sample Capillary ................. 45 Figure 2.5 Wilhelmy Balance for Contact Angle Measurements . . . . 48 Figure 2.6 Interfacial Testing System Composite Specimen Test Surface .................................... 55 Figure 2.7 Interfacial Testing System Apparatus ............. 55 Figure 3.1 Surface Free Energy Plot for HMS4 Carbon Fibers ........................................ 60 Figure 4.1 Crystallization Endotherm of PR10X2 PPS at 265°C ........................................ 66 Figure 4.2 Crystallization Endotherm of PR09 PPS at 250°C ........................................ 66 Figure 4.3 Crystallization Endotherm of PR09 PPS at 235°C ........................................ 67 ix Figure 4.4 DSC Scan of PR09 PPS Heated at 5°C/minute ....... 68 Figure 4.5 DSC Scan of PR10X2 PPS Heated at 5°C/minute ..... 68 Figure 4.6 PR09 PPS Spherulites Crystallizing at 240°C ........ 72 Figure 4.7 Same Spherulites from Figure 4.7 at 240°C, 16 Minutes Later ................................. 72 Figure 4.8 SEM of PR10X2 PPS Resin, Plasma Etched ......... 74 Figure 4.9 SEM of PR10X2 Resin, Chemically Etched ......... 74 Figure 4.10 Confocal Micrograph of PR09 PPS Resin Crystallized at 250°C ............................... 75 Figure 4.11 A34 Fiber/PR09 Resin Crystallized 1 hour at 265°C ...................................... 80 Figure 4.12 A84 Fiber/PR09 Resin Crystallized 1 hour at 235°C ...................................... 80 Figure 4.13 HMS4 Fiber/PR09 PPS Resin Crystallized 1 hour at 265 °C .................................. 81 Figure 4.14 HMS4 Fiber/PR09 PPS Resin Crystallized at 235°C for 1 hour ................................ 82 Figure 4.15 HMS4 Fiber/Amine Terminated PPS Crystallized 1 hr at 250°C ............................ 84 Figure 4.16 A84 Fiber/Amine Terminated PPS Crystallized 1 hr at 250°C ............................ 84 Figure 5.1 Interfacial Shear Strengths of PR10X2 PPS with Carbon Fibers ............................. 92 Figure 5.2 Interfacial Shear Strengths of PR10X2 PPS Resin with 1M6 Fibers ........................... 92 Figure 5.3 Interfacial Shear Strengths of PR09 PPS with AS4 Fibers .................................. 93 Figure 5.4 Interfacial Shear Strengths of PR09 PPS with HMS4 Fibers ................................ 93 Figure 5.5 SEM of PPS/ Pitch Fiber Composite Failure ........................................ 95 Figure 5.6 Bulk Nucleated Crystalline Interphase (PR09 PPS/A84) ................................. 97 Figure 5.7 Transcrystalline Interphase (PR09 PPS/ HMS4) ....................................... 97 Figure A.l Schematic of Pendant Drop ................... 114 INTRD TIN The nature of the fiber-matrix interphase in a polymer composite can have a large effect on its mechanical behavior. It is through this region between the fiber and the bulk resin matrix that load is transferred from the matrix phase to the high strength, high modulus fibers. In the case of a semicrystalline thermoplastic matrix composite, polymer adsorption on the fiber during processing, fiber-matrix chemical and physical interactions, and interphase crystalline morphology are interfacial mechanisms which can determine the level of fiber-matrix adhesion for the system. The semicrystalline interphase under investigation in this work is polyphenylene sulfide (PPS). PPS is synthesized by the following reaction: Cl—Ph-Cl + N018 .. (-PIr-S-LI + NaCl The mechanism of the polymerization has been described by Fahey and Ash‘. PPS is a semicrystalline thermoplastic with a glass transition temperature around 85°C and a melting range peak around 285°C. The molecular weight of commercial grade resins ranges from 18,000 to 80,000}3 When left above its melting point for an extended period of time in the presence of air, the resin undergoes complicated changes, which increase its molecular weight. The resin will also blacken somewhat. This process 2 crosslinking is often desirable and an air curing step is sometimes used for molded composites. In the present research, however, it is necessary to have all composites and resins processed in a reproducible and identical manner. In addition, surface crosslinking and oxidation may mask many of the surface effects under investigation. All processing at melt temperatures was therefore conducted in an inert (e. g. argon) environment to prevent reactions with oxygen. In composites of PPS with carbon fibers, Phillips Petroleum Company has reported an increase in transverse tensile strength from 30 MPa to 74 MPa and an increase in transverse flexural strength from 44 to 141 MPa when their extrusion grade PPS resin is replaced with their composite grade PPS resin‘. Since both of these mechanical tests are highly sensitive to the strength of the fiber-matrix interphase, it is suspected that the increase in strength is a consequence of a difference in interphase properties between composites made with the extrusion grade PPS and the composite grade PPS. The present work is an investigation of the interphase of PPS/carbon fiber composites that will determine what characteristics of the interphase bring about composite mechanical property changes such as that noted above. One possible reason for the increase in interphase strength for the composite grade resin is that the surface free energy of the composite grade resin may be significantly lower than that of the extrusion grade, which would provide a greater reduction in interfacial free energy between PPS and the fibers during processing. A significantly higher surface free energy for the composite grade would suggest that it is a cleaner resin that contains less surface active impurities which interfere with 3 significantly lower than that of the extrusion grade, which would provide a greater reduction in interfacial free energy between PPS and the fibers during processing. A significantly higher surface free energy for the composite grade would suggest that it is a cleaner resin that contains less surface active impurities which interfere with adhesion than the extrusion grade resin. Surface active contaminants could include low molecular weight oligomers, unreacted starting materials, or other impurities. Surface active components will concentrate at surfaces and interfaces due to their lower surface free energy relative to the bulk polymer. In this way, impurities that are present at low concentrations in the bulk may be present in significant concentrations at surfaces and interfaces, and may interfere with efficient adhesion of the matrix to the fiber. The extent of these surface free energy effects has been determined by measuring the surface free energies of PPS melts at the composite processing temperature of 330°C. For this purpose, a pendant drop surface tension measurement apparatus capable of measuring surface free energies at high temperatures in an inert environment was constructed. A chemical reaction between the phases is another factor that could affect the fiber-matrix interphase strength in a composite. Chemical bonding is commonly used to increase the interfacial adhesion of thermoset / glass fiber composites. The possible chemical bonding of PPS to carbon fibers, which could be hindered by impurities present in the extrusion grade PPS but absent from composite grade PPS has been examined by subjecting carbon fibers to processing conditions in the presence of compounds which contain the same functional groups as PPS. 4 The levels of PPS/carbon fiber adhesion for different resins, carbon fibers, and processing conditions were measured by an in-situ fiber indentation test and reported as interfacial shear strengths. The combination of the results from the surface free energy, chemical reactivity, PPS morphology, and interfacial shear strength studies have been used to explain which characteristics of the PPS/carbon fiber interphase are most important in determining the level of interfacial adhesion attained. EFERENQE l. D. H. Fahey and C. E. Ash, Macromolecules, 34, 4242 (1991). . J. S. Dix, Chemical Engineering Progress, 42 (1985). C. J. Stacy, Journal of Applied Polymer Science, 32, 3959 (1986) eww Phillips Petroleum Company Data. HAPTER I B r n This chapter is a review of the literature that is pertinent to the results chapters (chapters 3-5). It includes a discussion of surface free energies of polymer melts and fibers, thermodynamics of wetting, semicrystalline morphology in neat resins and composites, and fiber-matrix adhesion. F En Eff Surface compositions can differ greatly from the compositions of bulk mixtures. In a mixture of two or more components, thermodynamics dictates that the components with lower surface free energy will be overrepresented at the surface. This surface segregation serves to minimize the surface free energy of the overall mixture, but is never a complete segregation for miscible blends due to other considerations, such as entropy effects. These surface phenomena occur at surfaces and interfaces in polymers and polymer composite systems. Koberstein‘ has demonstrated this for the miscible blend of polystyrene (PS) and poly(vinyl methyl ether) (PVME). The surface free energy of PVME is about 10 dyne/cm lower than that of PS. The surface free energy partitioning effect was so great that addition of 20% PVME to PS dropped the overall surface free energy to within 20% of the value for pure PVME. In this way, constituents present in low amounts in the bulk can have large concentrations at surfaces or interfaces. This means that an interface between bulk polymer and a fiber 6 can also contain significant amounts of other material not characteristic of the bulk composition. The removal of certain impurities or even the addition of a small amount of a different material could drastically improve polymer/fiber adhesion and, consequently, polymer and composite performance. In addition, molecular weight has been found to have a significant effect on surface tension. For linear polymers, surface tension generally increases with increasing molecular weight, indicating it might be favorable for low molecular weight oligomers to concentrate at surfaces and interfaces. With their smaller size also increasing their mobility, it is quite possible that low molecular weight materials can concentrate at interfaces and interfere with effective polymer/fiber adhesion. Chemical interactions, such as acid base behavior, can also play a role in interfacial adhesion. Greenberg2 conducted a study of poly(acrylic acid) which shows that acid-base interactions between polymer and filler are important. Silica and aluminum oxide fillers had little effect on the polymer’s glass transition temperature. However, the addition of calcium silicate, a basic filler, raised the glass transition temperature 15°C. This was attributed to the restriction of chain movement caused by the acid-base interactions between polymer and filler. In a fiber-polymer composite material, surface and interface thermodynamic considerations can lead to the significant concentration of impurities which are present at low levels in the bulk. The wetting of the fiber by the molten or curing matrix is taken as a necessary condition for the proper formation of any composite material’. It is during this step that intimate contact between the two constituents is established, which leads to good adhesion of matrix to fiber in the final composite. Microscopic 7 examination of composite failure surfaces that reveal matrix material still adhering to broken fibers is taken as evidence that good wetting was established during composite processing. On the other hand, the presence of clean, bare fibers in the locus of failure indicates that poor wetting occurred during processing. For unidirectional fiber composites, mechanical tests which measure off-axis strength, such as short beam shear and transverse tensile tests, give a measure of the level of fiber matrix adhesion, and consequently are believed to indicate whether or not good wetting was established during composite processing. Thermodynamics provides a starting point for a study of fiber and matrix surface energetics and for quantifying the interactions which can occur when fiber and matrix are brought together. For small molecule liquids where the time constants for the rearrangement of molecules is short, equilibrium is easily established and the thermodynamics which describe the interactions between the solid surface and the liquid phase can be extrapolated to predict polymer matrix interaction with the reinforcing fiber“. In the case of high molecular weight polymers, time constants may be so long that thermodynamic equilibrium is not attainable within a practical processing time. Nevertheless, the equilibrium energetics may still be useful indicators of whether or not good wetting should be expected. When a liquid having a surface tension 7”, (liquid in equilibrium with its vapor) is placed on a solid surface with surface tension st (solid in equilibrium with vapor), the liquid will either form a draplet on the surface or spread over the surface and form a film. If a droplet is formed a relationship between the solid surface free energy and the liquid surface free energy can be expressed in thermodynamic terms. 8 The surface free energy of the solid-liquid interface is labelled 7L5 (liquid in equilibrium contact with the solid) and the equilibrium can be expressed as YLV = Yrs + YLV “’56 where 9 is the angle formed between the stable, equilibrium drop surface and the solid surface as measured through the liquid. Neither the solid surface tension nor the solid-liquid interfacial tension can be measured directly but the difference between the two is the product of the cosine of the contact angle and the liquid surface tension. Liquids that form contact angles greater than 90° are called ”non-wetting”. For the situation of a sessile drop, i.e. a droplet sitting on a solid surface, the liquid could be said to "bead up”, as water does on a freshly waxed automobile hood. Liquids that form a contact angle less than 90° are termed ”wetting". If the liquid does not form a droplet, i.e. the contact angle is 0°, the liquid will spontaneously spread into a film, and the equilibrium relationship does not hold. The equilibrium is then expressed by an inequality where st ' Yrs 2 YLV because as the liquid flows it is decreasing the solid-vapor interface, covering the solid surface with the spreading liquid. This inequality indicates that for all cases of wetting and spreading, the surface tension of the wetting liquid must be lower than the solid surface tension in order to provide the thermodynamic driving force for spontaneous wetting or spreading. 9 The Work of Adhesion (W A), which is an expression for the reversible thermodynamic work required to create a solid-vapor surface and a liquid-vapor surface by pulling apart the solid-liquid interface is defined by the relationship W. = m + my - us By substituting the equilibrium expression for the droplet forming the contact angle, the work of adhesion is given by W4 = y”, (l + 0086) This work of adhesion is an ideal, reversible equilibrium expression that is not a measure of the actual mechanical adhesive bond strength between two adhesively joined interfaces. The actual adhesion of surfaces is not an equilibrium situation and the energy involved in separating the two surfaces also involves many other factors, such as toughness and crack propagation. The thermodynamic WA expression indicates that the maximum value of the work of adhesion occurs when the contact angle is equal to zero and the work of adhesion is twice the surface tension of the liquid. Although these thermodynamic expressions apply to equilibrium situations they can also be used to understand the wetting process in a polymer/fiber composite. 10 F En i fPl rMel Surface and/or interfacial free energy (also called surface tension and expressed in dyne/cm or erg/sq cm) is the key factor in determining whether or not a constituent in a mixture will concentrate at a surface or interface. This important surface parameter can be thought of as the amount of work required to create a new surface, i.e. the work required to pull some material out of the bulk and into the surface. It is therefore energetically favorable that constituents with low surface free energy will migrate to the surface. This tendency is offset somewhat by entropic effects and can also be influenced by interaction effects, such as the acid-base effect described earlier, but the minimization of surface and interfacial free energy is the overall driving force. Since the interphase composition of a polymer composite can be far different from the bulk composition, it is valuable to know surface energy and chemical properties of all the constituents in a polymer so that the interphase composition may be understood. Once the interphase has been characterized, it may be possible to ”tailor" polymers for optimum interfacial properties in the same manner as fibers are tailored with coatings and sizings. Several options for the determination of polymer surface free energies are available. The pendant drop method was chosen from among these methods because it is well suited to systems with long equilibration times and can be coupled with a microcomputer to automate the analysis. The method is based on the shapes of drops hanging from capillary tips. 11 The shape of a hanging drop is the result of the balance between gravitational forces, which tend to stretch the drop out, and surface tension forces, which tend to keep the drop spherical to minimize its surface area. Figure 1.1 is a diagram of a typical pendant drop. The Bashforth and Adams equation describes the resultant shape of a droplet and can be written as 1 are + (TI/7) W) cob) + 2 where R1 = the first radius of curvature at the point (x,z) b = the radius of curvature at the drop apex x,z, and o are defined on Figure 1.1 B = Apgb’h Ap = the density difference between the two fluids (the liquid of the drop and the surrounding atmosphere) g = acceleration due to gravity 7 = the surface tension of the liquid. Andreas’ developed a method using a shape dependent quantity, S. This quantity is a ratio of two diameters on the droplet, the denominator being the equatorial diameter ((1,) and the numerator being the diameter ((1,) measured a length d.= up from the apex of the drop. The parameters B and b are combined to form the 12 \/ Figure 1.1 Schematic of a Pendant Drop. 13 quantity H (H = -B(d,/b)2). The values of H as a function of 8 have been tabulated using empirical information from water droplets? Surface free energy can be calculated as 2 4984 H This treatment can be good to a few tenths of a percent with a good optical setup. An alternative to measuring (1, and d, and using tables is to solve a set of differential equations numerically. These equations are derived from the same Bashforth-Adams equation. The resulting equations are integrated numerically and compared to the actual digitized drop profile. The shape factor parameter is varied until the analytical profile fits the experimental profile. This has been done by Rotenberg‘, who used an Euler method approach, and by Koberstein”, who used a Runga-Kutta type of algorithm. En i f n Fibe Surface free energies are often partitioned into two different components. These components are based on the chemical nature of the material. The partitioning generally involves London-van der Waals forces, usually referred to as dispersive forces, and other higher order interactions. Kaelble’ built upon some earlier work of Fowkes and proposed that the surface free energy of a liquid or a solid is composed of dispersion interactions and higher order ones such as polar interactions. Other higher order interactions that have been used include acid-base components and hydrogen bonding components.‘0 If liquids with known surface free energy 14 components are used as contacting fluids with solid surfaces, the polar and dispersive components of the solid could be determined indirectly by measurement of solid-liquid contact angles. This results in an expression for the work of adhesion which involves the contact angle and the liquid surface free energy which could be measured directly through the expression: P” P“ YLV (“0056) _ om YLV st DV‘ ‘ 5V * T 2hr Yrv Where the surface tensions refer to the dispersive (7°) and polar (7”) components of the solid (SV) and liquid (LV) phases. A plot of the left hand side of this equation versus the ratio of the square roots of the polar to dispersive ratios of a series of contacting liquids results in a straight line whose slope and intercept determine the dispersive and polar component of the surface free energy of the solid. This procedure was used by Hammer and Drzal" to determine polar and dispersive components of the surface free energies of carbon fibers. The contact angles were measured with a gravimetric Wilhelmy technique. The surface free energies of the carbon fibers will therefore be partitioned into dispersive and polar components. Carbon fibers of the type used in this study have been characterized previously with this method.”13 The increase in the polar component of surface free energy has been used to explain the increase in adhesion between carbon fibers and an epoxy matrix as the amount of commercial fiber surface treatment time was increased.“ The polar component of the surface free 15 energy of the fibers increased as the oxidative surface treatment increased the amounts of oxygen on the fiber surface. Th Mo h l f P Polyphenylene sulfide is a semi-crystalline thermoplastic. Crystallization of polymers occurs in two steps. First, a crystallite must be nucleated. This occurs when a PPS chain first begins to fold into a lamellar structure at a nucleation site. A nucleation site can be of several types, including impurities (i.e. dirt particles), PPS molecules that are structurally or chemically different from the bulk PPS, and reinforcing fibers in composite materials. The rate at which nucleation occurs is extremely sensitive to the temperature or degree of supercooling of the polymer melt. The higher the degree of supercooling, the more nucleation sites become available and the more rapidly nucleation occurs. Carbon fibers can act as nucleation sites for PPS crystallization. If the number of active nucleation sites on a fiber is high enough, the crystallites are so close that they impinge on one another. In this case lateral growth is hindered and the only direction available for continued growth is perpendicular to the axis of the fiber. This results in the formation of a cylindrical sheath structure which surrounds the fiber and is referred to a transcrystalline region. The second step is crystallite growth or propagation. During the growth process the chain continues to fold in a lamellar structure. Crystallite growth proceeds symmetrically outward from the nucleation site. During isothermal crystallization, the radial growth rate (the increase in radius of the growing spherulite) 16 will remain constant. Growth in a given direction will be halted when the growing spherulite encounters an obstacle in that direction. Obstacles to crystallite growth include other crystallites and, in the case of composite materials, reinforcing fibers. As with the nucleation step, a lower crystallization temperature leads to more rapid activity. Lowering the crystallization temperature thus increases the rate at which both nucleation and crystallite growth occur, leading to much faster crystallization rates at lower temperatures. During isothermal crystallization, both nucleation and crystalline growth are occurring continuously. Once the temperature of the PPS is brought below the glass transition temperature of 85°C, both processes stop. The resulting morphology of the fiber/matrix interphase region will fall into one of three major categories. Schematics of the three interphase types are presented in Figure 1.2. 1) Amorphous Interphase: In this case the interphase region of the composite material is devoid of crystallites. This will be the case when the melt is brought from the processing temperature to below the glass transition point rapidly, giving the crystallization process little time to occur. 2) Crystalline Interphase: In this case, the majority of crystallization has nucleated from the bulk. Although there may be some crystallites that have grown from the fiber, they did not nucleate densely enough to form a transcrystalline sheath around the fiber. Crystallites nucleated in the bulk continue to grow toward the fiber and either continue to grow around the fiber or stop growth at the fiber surface. 17 Amorphous Chm »yswe ara- < Transcrystalline Figure 1.2 Possible Interphase Morphologies. l8 3) Transcrystalline Interphase: In this case, high density nucleation of crystallites on the fiber surface has occurred. Adjacent crystallites grew together until they impinged on one another and crystallization from the fiber surface continued radially out into the bulk polymer. The result of the transcrystalline layer is a cylindrical sheath of polymer crystalline material surrounding the fiber. The nucleating ability of carbon fibers in a semicrystalline polymer melt varies from fiber to fiber and has been studied by several researchers. In general, higher modulus fibers have been found to nucleate crystallinity more effectively than lower modulus fibers. Several factors have been proposed as the cause of the higher nucleating ability of higher modulus carbon fibers. One factor is epitaxy, or geometric matching of the surface crystallite sites on the fiber with the lamellar fold geometry of a nucleated crystallite. Epitaxy has been found to be a driving force for almost any type of fiber", but a higher degree of graphitic order leading to more numerous epitaxial fiber surface sites leads to higher nucleating activity for higher modulus pitch and PAN based carbon fibers. This higher degree of order can be attributed to the treatment of the fiber during graphitization. The fibers are converted from the organic precursors to high modulus graphite fibers by stressing the fibers at very high furnace temperatures in an inert environment. The higher temperature treatment will result in a more highly graphitized and higher modulus fiber. The higher temperature treatment leads to longer and more regular crystalline sites on the fiber surface. Donnet'6 reports that as the high temperature treatment temperature is increased from 1000°C to 3000°C, the characteristic surface graphitic crystallite 19 length increases from around 3 nm to more than 10 nm. Higher modulus fibers therefore have longer and more perfected crystallite sites that provide epitaxial regions to promote polymer crystallite nucleation. The increase from 3 nm to more than 10 nm is significant for the nucleation of PPS because the chain fold length of PPS is 7 nm to 9 nm. This suggests that high modulus PAN based fiber surface graphitic basal plane edges are long enough to support one fold length, whereas the lower modulus fibers have surface graphitic crystallites which are not long enough to support a chain fold and nucleate the crystallization process of PPS. In addition, the surface graphitic crystallites of the high modulus fibers are straighter and more parallel to the fiber axis, which also contributes to the formation of a transcrystalline layer. Epitaxy can be enhanced through the use of a surface nucleating agent. Phillips and Greso‘ looked at AS4 carbon fibers in PEEK resin before and after treatment of the fibers with a B-Cu pthalocyanine surface agent. The surface agent produced a surface epitaxial structure that enhanced PEEK nucleation so that the AS4 fibers that had not developed transcrystallinity before treatment did so after treatment. Another factor that has been suggested as playing a role in determining how well a fiber promotes crystalline nucleation is the match of surface energetic factors between the fiber and the resin. Lopez and Wilkesl7 found that as the polar component of surface free energy of carbon fibers increased, the tendency of the fibers to generate transcrystallinity in PPS decreased. This was attributed to the fact that the PPS resin is very nonpolar and would not interact well with a highly polar fiber surface. Strong fiber-matrix attractive interactions could cause segments of 20 polymer chains to lie parallel to fiber surfaces, providing the first fold length for a lamellar structure. Still another factor that has been suggested as important is the higher modulus of the fibers which would lead to higher interfacial stresses at the fiber resin interface during melt solidification. It has been well documented that inducing shear stress at a fiber matrix interface can induce transcrystallinity where it would not normally exist.” This is accomplished by dragging the fiber through the melt. The stress induced at the interface leads to ordering of the polymer chains into a configuration favorable to transcrystalline formation. This factor may be important in generating transcrystallinity in processes such as pultrusion where large amounts of shear are generated at the fiber-matrix interface by pulling the fibers through the melt. It has also been speculated that the differences in coefficient of thermal expansion between the fiber and the matrix could lead to transcrystalline formation. As the melt cools, the matrix tends to shrink more than the fiber due to the lower coefficient of expansion for fibers. This relative motion could build up stresses that could lead to chain ordering by the same mechanism as dragging fibers through melts. The fact that higher modulus fibers, which have lower coefficients of themal expansion than do lower modulus fibers, are more likely to produce transcrystallinity has been used to support this idea. The difference in coefficients of thermal expansion, however, is unlikely to be important in the present work. The difference in expansion coefficients for high modulus fibers which do produce transcrystallinity and for lower modulus fibers which do not is very small. This suggests that the differences in crystalline nucleating ability between the fibers must be due in large part to other factors. 21 Some or all of the above factors may play a role in determining the nucleating ability of carbon fibers. All of the theories predict that the higher modulus carbon fibers, which generally have fewer polar groups, have a higher degree of graphitic order, and by definition have a higher tensile modulus, will promote crystallization of PPS more readily than will the lower modulus carbon fibers. ininf mi llinThrmli 11' in Crystallinity of thermoplastics can be studied by observation of crystallization as it occurs from the melt state. This can be done by measuring the crystallization exotherm as the sample is crystallized, using a thermal analysis method such as differential scanning calorimetry (DSC). This method has been used extensively for several thermoplastics and for PPS in particular by several researchers. Crystallization can also be studied optically as it occurs by observing thin films in an environmental hot stage under a microscope. Polarized light microscopy is particularly valuable for this task as the crystalline portions of the sample refract light and will appear as bright objects against a dark background under cross polarized light. Crystallization of polymers in the presence of fibers can be studied by sandwiching fibers between two layers of polymer film before melting of the polymer. 22 The starting point for studies of polymer crystallization kinetics is the well known Avrami'9 equation: C(t) = 1 — cxp(-Kt") Where: C = volume fraction crystallized at time t K = avrami rate constant n = avrami exponent. Avrami plots of ln{-ln(l-C)} versus 1n(t) are straight lines with slope = n and y intercept of ln(l(). The parameters n and K do not have generally accepted physical meanings. A value of n roughly equal to 3, however, is often taken to indicate three dimensional crystalline growth while a value of 2 indicates two dimensional growth. The Avrami equation contains several built-in assumptions. These include a spherulitic geometry with diameters increasing linearly with time, negligible density changes upon crystallization, evenly distributed nucleation sites, and primary (pre- impingement) crystallization. Systems which do not rigidly adhere to these assumptions may also be fit with Avrami analysis, although this sometimes requires modifications. Cebe and Chung20 studied the effect of crystallization temperatures on the melting endotherms of PPS. They observed two melting peaks for PPS. The melting peak at the higher temperature was attributed to primary crystallization, defined as crystallization occurring in the form of growing spherulites. The lower melting temperature was attributed to secondary crystallization, which occurs in regions between the primary spherulites and constitutes approximately half of the total 23 crystallinity for PPS material. The secondary crystallites are generally made up of less perfect chains and may run out of crystallizable material before forming complete spherulites. Less energy is therefore required to destroy secondary crystalline material, which accounts for its lower melting temperature. The melting temperature of the secondary peak can be moved closer to the primary peak temperature by annealing at high temperatures, thereby allowing the material to reorganize into a more perfect structure. The effect of repeated high temperature cycling of PPS on crystallization kinetics was reported by Mehl and Rebenfeld.“ Samples of PPS were heated to high process temperatures as many as ten times in 3 Differential Scanning Calorimetry (DSC) unit under dry nitrogen. Chain extension and crosslinking led to slower crystallization for samples treated at very high temperatures. The authors also observed that melt processing of PPS at temperatures below the equilibrium melting temperature, Tm", led to incomplete destruction of crystalline structure, leading to much more rapid crystallization of samples upon cooling. A Hoffman-Weeks22 technique of extrapolating a melting temperature (Tm) versus crystallization temperature (Tc) plot until it intersected the TIn = Tc line was used to determine the equilibrium crystalline melting temperature. The effect of high temperature melt processing history was also considered by Budgell and Day.23 They calculated Avrami parameters for PPS that had undergone varying melt treatment. The parameters were found to be highly dependent upon the melt history of the polymer. Long times at high temperatures were found to decrease the crystallization rates. This effect is caused by the destruction of pre-existing 24 crystalline substructure and by degradation of the resin for extremely long times or high temperatures. They estimated the equilibrium crystalline melt temperature to be 340 i 15°C for PPS. The more complicated situation of nonisothermal crystallization was considered by Collins and Menczel.“ They found that Avrami analysis was insufficient for nonisothermal processing due to nonlinearities in the data. It was also conjectured that for rapid crystallization at high cooling rates the mechanism of crystalline growth shifts from three dimensional to two dimensional. The effect of seeding PPS melts with nucleating agents was examined by Song and coworkers.” Several nucleating agents, including kaolin, talc, tin oxide and calcium oxide were used to increase the rate of PPS crystallization by acting as crystallization nucleation sites. Particles in the size range of 0.1 microns to 35 microns were used to nucleate PPS. It was found that the effectiveness of the nucleating agent depended on the type used as well as the size of the particles. Smaller particles were found to nucleate crystallization more effectively. It was also found that a higher concentration of nucleating particles led to more rapid crystallization. The crystallization of PPS in the presence of carbon fibers was considered by Caramaro and coworkers."5 They found that the crystallization of PPS could be slowed considerably when transcrystallinity did not occur. The Avrami exponent was found to be in the range of l-l.8 for fiber filled resins as opposed to 3 for unfilled resins. It is generally assumed that an Avrami exponent of 3 indicates a three dimensional spherulitic growth rate, whereas a value around 2 indicates two 25 dimensional growth. The lower values for the filled systems may indicate a change in morphology or simply the slowing of crystalline growth due to the hindrance of the fibers. Short beam shear tests indicated that failure of PPS/carbon fiber composites went from adhesive to cohesive for high temperature treatment transcrystalline specimens. There was no additional strength, however, as the high temperature treatment weakened the matrix strength substantially due to thermal degradation. The highest shear strengths were obtained with transcrystalline samples that had not undergone extremely high temperature processing. The failure for these systems was still interfacial, but the matrix had retained its strength and had not been degraded by severe processing. Desio and Rebenfeld27 studied the crystallization kinetics of PPS/fiber composites. They found that Avrami analysis of the crystallization was nonlinear but that it could be modeled well with either a parallel or series Avrami analysis. The parallel model was fit with a five parameter nonlinear regression. Two processes were assumed to occur in parallel. The five parameter fit included an exponent and constant for each process plus the fraction of crystallization that occurred by the first process. The series model assumed that process two occurred after the completion of process one and two lines were fit to the Avrami curve rather than the usual one for linear systems. Cold crystallization, or crystallization above the glass transition temperature but well below the melt state was observed by Kenny and Maffezzoli.28 They found that under these conditions the presence of carbon fibers slowed the crystallization kinetics of PPS and also lowered the maximum possible crystallinity in the material 26 from 55% to 45%. This can be attributed to the fact that cold crystallized samples already have nucleating sites present (undestroyed crystalline precursors) so the presence of fibers under these conditions serves only to hinder crystalline growth. Th Mo hol f emi st lline Resins nd om osit The second approach to studying the crystalline structure of a semicrystalline thermoplastic is the observation of the structure after it has formed. Once the melt has solidified, crystalline material is no longer discernible since the sample has become Opaque due to the large crystalline content present and to the surrounding amorphous content of the sample. Microscopic examination of the crystalline structure of semicrystalline thermoplastics is feasible with Scanning Electron Microscopy (SEM) if the surface layer of the sample has been enhanced. This enhancement is done by the removal of amorphous material by chemical or physical etching. The goal of the etching process is to remove as much of the amorphous material while leaving as much of the crystalline structure intact as possible. The etching can be either chemical (using an agent such as chromic acid) or physical (using a beam such as an argon or oxygen plasma). Both chemical and physical etching have been used with excellent results. The crystalline morphology of PPS as revealed with chemical etching was studied by Lopez and Wilkes.29 Chromic acid and aluminum trichloride in toluene were both used to preferentially remove amorphous PPS and leave the crystalline structure intact. Concentration of etching solutions, etching temperature, and exposure time were all found to be critical in getting the right degree of texture 27 enhancement. The aluminum trichloride solution was selective enough to reveal very fine detail of the texture of the crystalline spherulites. The elucidation of crystalline structure of PEEK for subsequent SEM analysis was described by Lustiger.” 3‘ Etching was conducted using an argon plasma beam in an ion mill. The morphology of APC-2 (optimized PEEK/AS4 carbon fiber) composite was found to be transcrystalline in regions near fibers with bulk crystals existing in resin rich areas of composites. The orientation of crystalline lamellae on AS4 fibers was found to be flat on, with the folded chain surface lying parallel to the fiber surface, allowing good contact with amorphous polymer content. The orientation on HMS4 fibers was edge on, with the folded chain surface protruding perpendicularly outward from the fiber surface, allowing less interaction with amorphous polymer content with the exception of amorphous loops on the edges of the lamellar structure. This can have an effect on the adhesion level between fiber and polymer excluding material which had adsorbed during processing from the fiber surface. Another post—crystallization characterization approach is X-ray diffraction, which relies on the fact that X-ray scattering angles for crystalline and amorphous phases are different. Wide Angle X-ray diffraction was used by Johnson and Ryan32 to determine the crystalline content of PPS/carbon fiber composites. The X-ray diffraction peaks of interest were the PPS crystalline peaks (20 = 20.5° and 18.8°), the broad amorphous PPS peak (20 = 19°), and the carbon fiber peak (29 = 25.1°). The peaks were mathematically separated and integrated to determine the percent crystallinity of various samples. Composites that were quenched and annealed at 28 200°C for 2 hours had a maximum crystallinity of 33%. Samples that were slow cooled and subsequently annealed at 200°C for 2 hours had a maximum crystallinity of 55%. These results suggest that annealing at 200°C increases crystallinity somewhat, but is not sufficient to provide the high amount of crystallinity obtainable by slow cooling or isothermal crystallization directly from the melt. Another very useful technique for observing crystalline polymer samples that has come into use in the past few years is Confocal Scanning Optical Microscopy (CSOM). In the reflectance mode, this technique uses essentially the same optics as a conventional light microsc0pe, with the key difference being that the conventional light source is replaced with a scanning laser and that the reflected light is detected with a photoelectric detector rather than the microscopist’s eyes. As the laser scans across the sampling region, the reflected light that is not in the focal plane of the sampling region is eliminated by passing the light through a pinhole. As a result, semicrystalline samples that would normally be opaque due to scattering of reflected light caused by the presence of the out-of-focus signal, are viewable at depths of up to several microns. The microscope stage can be moved up toward the laser source, allowing the operator to obtain images at increasing depths, in effect optically sectioning the sample into vertical layers. This technique, which requires no more sample preparation than conventional light microscopy, provides a method for viewing crystallinity and transcrystallinity in solid, standard composites and eliminates the nwd for extrapolating results from thin film studies. The final composite is viewed rather than the crystallization as it occurs from the melt. This technique was used 29 previously to study aramid fiber in a polypropylene matrix33 and clearly delineated the existence of transcrystalline and bulk crystalline regions in solid samples. bur-Matrix Aggien; The level of adhesion of reinforcing fibers to the polymer matrix is an important factor in the determination of composite mechanical properties. Weak adhesion between fiber and matrix can prevent the efficient transfer of stress from the matrix to the fibers. Very strong adhesion, on the other hand, can lead to a brittle fiber-matrix interface, providing a easy path for crack propagation and premature composite part failure under stress. Knowledge of the level of adhesion in a composite material can lead to predictions of failure modes and the determination of interphase properties that will lead to the formation of a strong, yet tough interphase. The adhesion of fibers to polymer matrices may be evaluated by several approaches. For unidirectional composites, bulk mechanical tests which are sensitive to the level of fiber/matrix adhesion include tests such as transverse tensile strength, short beam shear strength, and transverse flexural strength. Micromechanical tests have also been develOped that test the adhesion of fibers to matrices on a single fiber level. These single fiber techniques focus on the interfacial shear strength between the fiber and matrix in a composite material and give a more direct measure of the adhesion between fiber and matrix than do bulk mechanical tests. Micromechanical results will be presented in this work. 30 he Ii e iz an Matrix Mechani l Pr rties The influence of crystalline spherulite size on the mechanical properties of semicrystalline polymers has been described by several researchers. Way and coworkers“ looked at the effect of spherulite size on the fracture morphology of polypropylene. They reported that larger spherulites, which form in the bulk at slow cooling rates, lead to higher concentrations of impurities at the spherulitic boundaries. Although the larger spherulites that form at lower cooling rates are internally stronger due to a more perfect microstructure, the overall material is weaker and less tough due to weak interspherulite boundaries. In addition to the concentration of impurities at spherulitic boundaries, the scarcity of tie molecules for large spherulitic structures has been pointed to as a source of weakness.3"3°'”'” Tie molecules are polymer chains that are contained in the crystalline lamellar structure in more than one spherulite. They act as crosslinks between the spherulites, toughening the boundary regions. Tie molecules are more abundant for smaller spherulites and can disappear as polymers are annealed. The evidence in the literature suggests that finer spherulitic structure leads to tougher semicrystalline materials. A tougher fiber-matrix interphase should also lead to higher interfacial shear strengths, due to improved crack resistance which can delay interfacial failure until higher loads are applied. Toughness of the interphase as determined could be an important factor for the PPS/carbon fiber system. The tensile moduli for all grades of PPS are very close to one another. Since the toughness of the matrix is equal to the area under its stress-strain curve, this means that the toughness of the resins will be proportional to their strain to failure. Rao and Drzal”, 31 in a study of epoxy/carbon fiber composites, reported interfacial shear strengths proportional to the matrix strain to failure times the square root of the matrix shear modulus, which suggests that the toughness of the interphase is indeed a strong factor in determining interfacial shear strengths. Since the extrusion grade and amine terminated PPS used in the present work were determined to fail at 1.5-2% strain and the composite grade PPS fails at 4-5% strain, the composite grade PPS should be roughly twice as tough. The increase in toughness can be attributed to the much smaller spherulitic texture of the composite grade PPS (0.3 to 0.5 micron diameters as opposed to 10-50 microns). mi 1 rxllin a: ML}! 2 Mn “T n 11".. ii ht“; 'nau Most of the relevant literature addressing the interfacial adhesion between carbon fibers and semicrystalline thermoplastics has emphasized the importance of the morphological structure of the fiber-matrix interphase. For crystalline (but not transcrystalline) interphases, it has been shown that larger spherulites present in the interphase reduce toughness and lead to lower adhesive strengths. In their study of PEEK composites, Lustiger and coworkers“0 observed that large spherulites generally lead to increased brittleness in the polymer. Large spherulites in the interphase also lead to lower fracture toughness and impact strength in PEEK/carbon fiber composites. Xiao and coworkers“ noted that fracture toughness in PEEK/carbon fiber composites increased as the characteristic spherulite size decreased from 8-10 microns to 1 micron. Larger spherulites reduce toughness because they lead to large intercrystallite flaw regions. Griffith’s flaw theory of mechanical failure states that 32 flaws below a critical size will not lead to failure but that flaws at or above the critical size will lead to failure.“2 Cracks propagate either around the edges of spherulites or through the spherulitic centers. Large flaw regions lead to rapidly propagating cracks and early failure of a crystalline polymer with large spherulitic structures. The effect of transcrystallinity on adhesion has been reported both to increase and decrease interfacial shear strengths. On the one hand, the fact that the crystalline structure is mechanically anchored to the fiber suggests that higher adhesion levels should be observed for transcrystalline composites. On the other hand, the presence of the transcrystalline structure around the fiber provides a straight, unhindered path along which interfacial cracks may propagate. In a study of PEEK/carbon fiber composites, Lustiger‘3 found that PEEK/AS4 composites, which did not have transcrystallinity, had no bare fibers in failure regions. Composites with HMS4 fibers that had a transcrystalline interphase, however, had failure regions full of fibers that had pulled out of the PEEK matrix cleanly. The HMS4 composites also had a lower interfacial shear strength than did the nontranscrystalline AS4 composites. In another study of PEEK/carbon fiber composites, Nardin“ reported that transcrystallinity tended to lower interfacial shear strength, particularly for rapidly crystallized samples. The minimum value of interfacial shear strength resulted from cyrstallization at the temperature where the crystallization rate was the maximum. The rate of cooling can have an effect on adhesion for amorphous systems as well. The dependence of the interfacial shear strength of amorphous composites on the rate of cooling was observed by 0am“ for polycarbonate/ aramid fiber 33 composites, who suggested that the slower rate of cooling allowed time for a more ordered structure to form. This suggests that even though the interphase may be amorphous, quenching rapidly through the rubbery state to below T, will yield a weaker interphase than either a slow cooling or isothermal hold step below the crystalline melting point for PPS. Folkes and Wong“ reported that transcrystallinity lowered the interfacial shear strength of polypropylene/ glass fiber composites. Other research has indicated that the presence of transcrystallinity increases the interfacial shear strength between the fiber and the matrix in a composite. Transcrystallinity was found to increase the interfacial shear strength in polypropylene/glass fiber composites by Schoolenberg and Van Rooyen". They also observed that higher interfacial shear strengths were observed for slowly crystallized composites. Chen and Hsaio48 looked at several resin/fiber systems and found that transcrystalline composites consistently had over 40% higher interfacial shear strength than noncrystalline composites. They also found that this increase in interfacial shear strength became less and less significant as the fiber loading increased and the transcrystalline regions impinged on one another. The fact that there is no clear consensus on the effect of transcrystallinity on interfacial shear strength suggests that the effect may well depend on the physical properties of the specific fibers and matrices used. In addition, it is evident that the strength of adhesion of transcrystalline, as well as crystalline, composites may be a strong function of the crystallization temperature and other processing variables. 34 REFEREN ° 1. J. T. Koberstein, et al., ”Interface Partitioning in Multicomponent Polymer Systems”, Materials Research Society Preprint, 1989. 2. A. R. Greenberg, Journal of Materials Science Letters, 6, 78 (1987). 3. D. H. Kaelble, Ehysiee! Chemistry ef Adhesion, Wiley Interscience, 1971. 4. W. D. Bascom and Drzal, L.T., The Surface Prep_erties ef Cnmon Fibers and flfheir Adhesien te Qrganie Pelymers, NASA Technical Report No. 4084, July, 1987. 5. Adamson, A.W., Physigl Chemistg ef Serfaces, 5th edition,Wiley-Interscik%6, 6. Y. Rotenberg, et al., Journal of Colloid and Interface Science, 93, 169 (1983). 7. Q. S. Bhatia, et al., Journal of Colloid and Interface Science, 106, 353 (1985). 8. S. H. Anastasiadas, et al., Polymer Engineering and Science,26, 1410 (1986) 9. D. H. Kaelble, Ehysieal Chemistg ef Adhesien, Wiley Interscience, 1971. 10.F. W. Fowkes, J. Adhesion Sci. Tech, 1, 7 (1987). 11. G. E. Hammer and L. T. Drzal, Application of Surface Science, 4, 340 (1980). 12. L. T. Drzal, et al., Carbon, 17, 375 (1979). 13. L. T. Drzal, Carbon, 15, 129 (1977). 14. L. T. Drzal, M. Madhukar, and M. C. Waterbury, Composite Structures, 27, 65 (1994). 15. P. J. Phillips and A. J. Greso, ANTEC I992, 779. 16. J. B. Donnet and R. C. Bansal, Careen Fimrs, 2nd Ed., Marcel Dekker, New York, 1990. 17. L. C. Lopez and G. L. Wilkes, Journal of Thermoplastic Composite Materials, 4, 58 (1991). 18. J. L. Thomason and A. A. van Rooyen, ICC] Proceedings 1990, 423. 19. L. Mandelkem, Castnllizntien ef Pelymers, McGraw-Hill, 1969. 20. P. Cebe and S. Chung, Polymer Composites, 11, 265 (1990). 21. N. Mehl and L. Rebenfeld, Polymer Engineering and Science, 32, 1451 (1992). 35 22. J. D. Hoffman and J. 1. Weeks, Journal of Chemical Physics, 37, 2072 (1962). 23. D. R. Budgell and M. Day, Polymer Engineering and Science, 31, 1271 (1991). 24. G. L. Collins and l. D. Menczel, Polymer Engineering and Science ,32, 1270 (1992). 25. S. S. Song, 1. L. White, and M. Cakmak, Polymer Engineering and Science, 30, 944 (1990). 26. L. Caramaro, B. Chabert, and J. Chauchard, Polymer Engineering and Science, 31, 1279 (1991). 27. G. P. Desio and L. Rebenfeld, Journal of Applied Polymer Science, 45, 2005 (1992). 28. J. M. Kenny and A. Maffezzoli, Polymer Engineering and Science, 31, 607 (1991). 29. L. C. Lopez and G. L. Wilkes, Journal of Polymer Science: Polymer Letters Edition, 24, 573 (1986). 30. A. Lustiger, Polymer Composites, 13, 408 (1992). 31. A. Lustiger, F. S. Uralil, and G. M. Newaz, Polymer Composites, 11, (1990). 32. T. W. Johnson and C. L. Ryan, Int. SAMPE Symp., 32, 1537 (1986). 33. J. L. Thomason and A. Knoester, Journal of Materials Science Letters, 9, 258, (1990). 34. J. L. Way, J. R. Atkinson, and J. Nutting, Journal of Materials Science, 9, 293 (1974). 35. J. T. Yeh and J. Runt, Journal of Materials Science, 24, 2637 (1989). 36. R. Greco and G. Ragosta, Journal of Materials Science, 23, 4171 (1988). 37. M. J. McCready and J. M. Schultz, Journal of Polymer Science: Polymer Physics Ed., 17, 725 (1979). 38. T. J. Pecorini and R. W. Hertzberg, Polymer, 34, 5053 (1993). 39. V. Rao and L. T. Drzal, Polymer Composites, 12, 48 (1991). 40. A. Lustiger, F. S. Uralil, and G. M. Newaz, Polymer Composites, 11, 65 (1990). 36 41. X. R. Xiao, J. Denault, and T. Vu-Khanh, Journal of Thermoplastic Materials, 5, 64 ( 1992). 42. N. M. Bikales, Meehanical Progrties of Polymers, Wiley-Interscience, 1971. 43. A. Lustiger, Polymer Composites, 13, 408 (1992). 44. M. Nardin, E. M. Asloun, and J. Schultz, Polymers for Advanced Technologies, 2, 109 (1991). 45. U. Gaur, G. Desio, and B. Miller, Plastics Engineering, p 43 , October 1989. 46. M. J. Folkes and W. K. Wong, Polymer, 28, 1309 (1986). 47. G. E. Schoolenberg and A. A. Van Rooyen, Composite Interfaces, 1, 243 (1993). 48. E. J. H. Chen and B. S. Hsiao, Polymer Engineering and Science, 32, 280 (1992). HAPTER H Materials and Exmrimenthl Method_s The principal materials used in this research were PPS resins and carbon fibers. The PPS resins were provided by Phillips Petroleum. The resins included, Phillips PR09 (extrusion grade resin), PR10X2 (composite grade resin), and an amine terminated resin. Three types of polyacrylonitrile precursor fibers (Hercules) of different tensile moduli and a high tensile modulus pitch based type of fiber (DuPont) were used. Experimental methods used included the pendant drop method for liquid surface tension measurement, various types of microscopy, fiber surface chemisorption measurement, and an interfacial shear strength measurement technique. In this chapter, the materials and methods used in the research will be described, laying the groundwork for the results and discussion chapters to follow. Mil—Materials The PPS resins studied were provided in granular form and were free from the additives (such as fillers, processing aids, or carbon black) sometimes blended in commercial grades. In addition, the PR09 and PR10X2 resins were provided in film form, which was ideal for hot stage thin film crystallization studies and for fabricating composites. The recommended processing for PPS composites is one hour of melt processing at 330°C followed by an appropriate crystallization schedule. Therefore, the first step in processing a PPS or PPS composite sample for analysis was always a one hour isothermal hold at 330°C. This temperature is high enough to ensure 37 38 melting of crystalline precursors, yet low enough to limit degradation. As discussed earlier, PPS materials are susceptible to attack by oxygen at the high temperatures required for melt processing. Therefore, all melt processing was conducted under an argon atmosphere. The long melt time of one hour ensures good consolidation of PPS composite parts, since PPS viscosity is low compared to many thermoplastics but is still appreciable. The viscosity of the PPS at 300°C as measured with steady shear in a Rheometrics RMSSOO parallel plate geometry was 250 Pa-s for low shear rates. The most pertinent mechanical properties for the neat resins are the tensile modulus and shear modulus for each resin. These are listed in Table 2.1. The tensile moduli of the materials were determined directly by measuring axial load versus displacement using a United Testing Systems screw driven mechanical testing device. Each material’s shear modulus was calculated from its tensile modulus with the relation‘: 2(1 +v) Where: G = Shear Modulus E = Tensile Modulus u= Poisson’s Ratio. Poisson’s ratio was assumed to be 0.45, which is a general value for thermoplastic materials.2 39 flfnhle 2,1 Tensile and Shear Modulus Values for Neat PPS Resins. Wization Tensile Modulus (GPa) swarm—ll Temperature PR10X2 , 250°C 3.8 i 0.2 1.3 i 0.1 PRO9, 235°C 3.8 i 0.2 1.3 i 0.1 PRO9, 250 °C 3.7 i 0.1 1.3 i 0.1 PRO9, 265°C 4.1 :1; 0.2 1.4 ;t 0.1 hAmine Term., 265°C 3.9 i 0.1 _ 1.3 :1; 0.1 A typical load versus displacement plot for the PR10X2 material is illustrated in Figure 2.1. The plot is linear, indicating elastic behavior, until the strain reaches about 2% (ASTM E 8M-93). At this point, curvature in the stress/strain behavior is evident, indicating that the deformation is now plastic. This makes the use of the embedded single fiber fragmentation test for interfacial shear strength invalid for the material, since one of the assumptions of the test is an elastic matrix material. Slamming There were three types of polyacrylonitrile (PAN) precursor carbon fibers used in this work. They include AS4 fibers (standard tensile modulus), HMS4 fibers (high tensile modulus), and 1M6 fibers (intermediate tensile modulus). All of these PAN based fibers were produced by Hercules. In addition, the 1M6 fibers were available with 0%, 20%, 100%, and 600% of the standard commercial oxidative surface treatment. The percentages refer to the processing times in an oxidative bath. For instance, 20% surface treatment indicates that the fibers were treated in the bath for 40 90 *‘i 80~ 70- 50- 4o~ Stress (MP0) 30~ 20» 10~ l A i A 1 A l A l 0.5 1.0 1.5 2.0 2.5 7. Strain 3.0 figure 2.1 Load Versus Displacement for Neat PR10X2 PPS Resin. 41 one fifth of the standard treatment time (standard time would be 100%). This series of 1M6 fibers allows the study of fiber adhesion for fibers that differ only in the extent of their surface treatment. Some of the properties of these carbon fibers are listed in Table 2.2. For comparison purposes, a high modulus pitch fiber was also used. The fiber was produced by DuPont and designated as a type G, pitch based carbon fiber with a tensile modulus of 827 GPa (120 million psi). All of the fibers appear to be smooth when viewed with SEM. Differences in roughness on the nanoscale have been observed with Scanning Tunneling Microscopy and are being further investigated for quantification.’ man Carbon Fiber Diameters and Tensile Modulus Values. Fiber Type Nominal Diameter Tensile Modulus 0““) (GPa) AS4 7 230 1M6 8 280 HMS4 5 345 — J Due to the surface sensitive nature of the work, all fibers used in the study were handled only with clean tweezers and only allowed to contact clean foil or glass surfaces. Allowing the fibers or PPS films or powders to come into contact with contaminants (such as oil from bare fingers) would compromise the integrity of any subsequent interfacial shear strength or surface composition measurements. 42 M urfa e Free Ene Measurements The surface free energy of the molten polymer matrix is an important factor in determining the eventual interfacial shear strength of the resulting polymer composite. Polymer melt surface free energy lower than that of the fiber is favorable since this situation provides a thermodynamic driving force for the polymer to wet the fiber surface. The surface tension measurement technique chosen was the pendant drop technique. This method uses the equilibrium shape of a hanging droplet to determine the surface free energy of the material. The environment for the pendant drop formation had to be capable of reaching temperatures as high as 330°C in a reasonable amount of time. In addition, due to the tendency of PPS to oxidize and crosslink in the presence of oxygen, the atmosphere surrounding the drop had to be inert. Argon was chosen as the inert gas, due to its availability with very low trace levels of oxygen. In order to provide a clean, leak free system, the environmental cell was made from high vacuum stainless steel components (Huntington). All connections in the cell were steel knife edge flanges sealed with copper gaskets. This setup eliminated complications of contamination from outgassing of rubber seals and O«rings. Heating was supplied by two 275 Watt coiled cable heaters located above and below the drop formation area. Viewports were optically flat glass to prevent distortion of pendant drop images viewed through them. A schematic of the pendant drop cell is presented in Figure 2.2. The pendant drop images were captured with a Panasonic solid state black and white video camera. The camera was selected because of the low light requirements 43 ”WWW figure 2.2 Pendant Drop Environmental Cell Schematic. 44 and its resistance to "image burn-in" because of solid state construction. The camera was equipped with a Nikon macro lens and bellows assembly, capable of magnifying the image of a 2 mm diameter drop enough to fill the monitor screen at a focal distance of several inches. Images were captured on a 640x400 pixel resolution frame grabber, manufactured by Progressive Peripherals. The high resolution of the frame grabber allowed the same reproducibility of measurements obtained by earlier video systems, which used 250x250 resolution frame grabbers, without resorting to complicated algorithms. In these experiments, surface free energies were calculated using the measurement of the width of the drops at two heights and tabulated shape factors. A schematic of the overall pendant drop apparatus is presented in Figure 2.3. The procedure for obtaining a pendant drop measurement was as follows: 1) Glass capillaries (Drummond Scientific) were scrupulously cleaned with Chromerge' solution and reverse osmosis purified water. The clean capillaries were left to dry in a clean oven at 100°C for several days. 2) PPS powder was loaded into a clean glass culture tube. The tube was sealed with aluminum foil and flushed continuously with argon to expel oxygen. 3) The culture tube with PPS was loaded into a heated aluminum block. The block temperature was maintained at 360°C with an insulated band heater. The oven block assembly was wrapped in fiberglass insulation to prevent heat loss. 4) After the PPS had melted in the culture tube, a clean capillary connected to a vacuum pump was inserted below the surface of the melt. Material was pulled into the capillary from below the surface, supplying undegraded PPS. 45 , Pressure All metal I ; Bakeout p1 93099 valve isolation Cryopump valve Linear #7 fl n9.33" 1 , j I > 7m,“ is attained for all resin/fiber combinations. larger differences may lead to more favorable wetting. 63 N L I NS' The surface free energies for the PPS resins and carbon fibers presented here indicate that good wetting should occur for all resin-fiber combinations, provided that the melt time is long enough. Higher temperatures, such as 330°C as opposed to 300°C, improve the chances of establishing good adhesion due to lower melt viscosity, as well as lower resin surface free energy. The surface free energies of the composite grade PR10X2 PPS and the extrusion grade PR09 are very close to one another. This suggests that higher adhesion of the composite grade resin to carbon fibers must be attributed to factors other than better wetting during composite processing, such as morphological differences resulting from a higher crystallization nucleation density for the PR10X2 resin. The similar surface free energies for the resins, which are assumed to be of similar molecular weight, also indicate that neither resin has an appreciable amount of low surface free energy material concentrating at surfaces, which would lead to poor adhesion. The surface free energy of the amine terminated resin is also similar to those of the other resins. This indicates that the amine end groups are not particularly surface active. They may, however, be interface active when brought into contact with a material which has a chemical interaction, such as copper metal. REFERENCES; l. V. Rao, PhD Dissertation from Michigan State University, 1991. 2. L. T. Drzal, M Madhukar, and M. C. Waterbury, Composite Structures, 27, 65 (1994). HAPTER IV a 1-: in zno M ahllor' f _’ ’_ ' i :n- P l: .u n Fi r In when. The morphology of the interphase region in a PPS/carbon fiber composite can have a strong effect on the level of adhesion and resulting interfacial shear strength between matrix and fiber. This makes an understanding of the morphology resulting from different resins, fibers, and processing conditions necessary in understanding adhesion of resin to fiber. The morphology of the PPS resins used in this study will be presented in this chapter. Crystallinity was examined using environmental hot stage optical microscopy, Confocal Scanning Optical Microscopy (CSOM), Scanning Electron Microscopy (SEM), and Differential Scanning Calorimetry (DSC). The crystallization and resulting morphology of the PPS resins in the absence of fibers will be discussed first. The PPS resins studied varied drastically in their crystallization rates and resulting morphological structure. Once the crystallization and resulting morphology have been discussed in detail, the morphology of PPS resins in the presence of various carbon fibers will be examined. The effect of interphase morphology on fiber matrix adhesion will be discussed in a later chapter. W: W The three PPS resins studied exhibited different crystallization rates at a given temperature. The PR10X2 resin crystallized the most rapidly and displayed the highest concentration of active nucleation sites. The PR09 resin was the slowest to 64 65 crystallize at any given temperature. The amine terminated resin was intermediate in crystallization rate. A DSC plot for PR10X2 resin at 265°C is presented in Figure 4.1. The crystallization at a temperature only 17°C below the melting peak for the material is effectively complete after only 20 minutes. Crystallization at the other temperatures studied, 235°C and 250°C , occurred in a matter of seconds. A DSC plot for the isothermal crystallization of extrusion grade Ryton' PR09 resin at 250°C is presented in Figure 4.2. The crystallization for this resin at this temperature is essentially complete after 2 hours. Even with an additional 15 degrees of supercooling, the crystallization of this resin is still much slower than that of the PR10X2 resin as shown in Figure 4.1. When the temperature of the isothermal crystallization is lowered even further to 235°C, the crystallization of the PR09 resin now occurs at a rate comparable with that of the PR10X2 resin at 265°C. As shown in Figure 4.3, crystallization of the PR09 resin at 235°C is essentially complete after 20 minutes. The PPS samples crystallized above were subsequently melted in the DSC by heating the samples under nitrogen to 330°C at a rate of 5°C/minute. The crystalline melting endotherm for PR09 isothermally crystallized at 250°C is presented in Figure 4.4 and the scan for PR10X2 isothermally crystallized at 250°C is presented in Figure 4.5. The melting scans for the two types of resin are essentially identical. Both materials exhibit dual melting peaks. The larger peak represents the primary crystalline regime and the smaller peak located about 25°C lower than the larger peak represents the secondary crystalline regime, as discussed in the background section. 66 “are: one I!) "I. EAT DSC File: m.” Ilu: ”.8000 - “orator: 0' name: I" m DOC/m M Date: “III/92 03: 87 We: mmrnmnaacmrmmnmc that fle- We) 160 More: van we moo do n- hanl figure 4.1 Crystallization Exotherrn of PR10X2 PPS at 265°C. 80-910: mm ran. m: 11 Fire: ssrnvroupz am I 0000 I. D S C aerator: 65F mm»: In mt nee/aloe run one: 00/15/32 :1: to ”at: VININ HUI rum It!" M 0.: that no- Ila! O 2 l mid Standard Byron PPS at 'C WI, complete after 2 hrs. .'=u r60 do :00 so are u. bin) We] V21! me me figure 4.2 Crystallization Exotherrn of PR09 PPS at 250°C. 67 )0: mm ‘47. m 7 DSC F110: m.” flu: 11.“ - Water: 0' name: In out? In one: “In!!! Mu cogent: vrmru nut cum lrm m 1 DJI- 3 a re- i 3 fl . i o u- "; .4 0.1:- Md Rnon PPS Crystallizing at 235°C financially mplete after 20 lab. 0.: 7 v at too do no do are n- hml Calm: va.u we not: figure 4.3 Crystallimtion Exotherm of PR09 PPS at 235°C. 68 ' scoot..- man mt. m memo: D S C rue: man.“ One: 0.0000 n. war-tor: 09' memo: m to 33°C At ac/uru am note: oo/ra/oe at no ell-oat: our no no com at 25°C. 0cm in s nee/urn 0.: DJ“ E e '5' t S 1 ‘4‘ Sum Ryton res Isothermlly Crystallized d 250°C Heated at 5 Chain 4.: - . - - - a . U 100 1.0 200 no ”0 35° Incl-euro (‘61 We! VI.” We .00 figure 4.4 DSC Sean of PR09 PPS Heated at 5 °Clminute. 0.1.: C” m '8. EAT DSC '11.: m.“ flu: It.” - tor-tor: . “that I“. to ace 0? MD! .1! “ti: ”III/D! ID: 13 Gin-gnu til ”I RATED AT 8 Elm!!! r0 noc NY “NJ! 0.0- 1 '5 4.:- i I 8 d s i -e.o~ “Grade PPS MCMd 250°C ‘ mu S'Clllin -o.su 4.: v v - v r v - a - I: do do coo do ado no rm 1°C! Donor-o1 Vt." that non Figure 4.5 DSC Sm of PR10X2 PPS Heated at 5°C/minute. 69 The secondary crystalline regime is made up of material that crystallizes between branches of the major spherulitic structure and also includes smaller crystallites that were formed rapidly. The secondary crystallites have higher surface energies and therefore melt at lower temperatures than the primary crystallites. The DSC results presented above show that during melt crystallization of a given resin, the rate of crystallization increases as the degree of supercooling of the melts is increased. This rate increase is due to two major factors. The first factor is an increase in the rate of crystallite growth. The second is an increase in the number of active crystallite nucleation sites with the increase in supercooling of the melt. The observation that the PR10X2 resin crystallizes faster than the PR09 at a given temperature, however, is due primarily to a greater abundance of active nucleation sites in the PR10X2 resin. This is confirmed by examining crystallized samples of the resins. The PR10X2 resin has many more spherulites than the PR09 resin. If the faster crystallization were due to faster crystallite growth, the PR10X2 would have fewer and larger spherulites, rather than more and smaller ones. In addition to the DSC thermal analysis of PPS crystallization, the crystallization of PPS in thin films from melts was directly observed. Films of 4 mil in thickness were observed crystallizing under argon in the Linkham THM 600 hot stage under an optical microscope equipped with cross polarizers. The films were melt processed following the processing schedule used for PPS composites in other sections of this study. The first step was an isothermal hold at 330°C for one hour. The second step was an isothermal hold at a prescribed crystallization temperature. 70 The PR10X2 resin crystallized rapidly and with such a high degree of nucleation that individual spherulites were too small to resolve with the optical microscope. A sample ramped to 250°C after the one hour hold at 330°C crystallized to the extent that the view field was completely darkened within 60 seconds. A sample ramped to 265°C from the 330°C melt step crystallized sufficiently to darken the view field in less than 3 minutes. The crystallization ability of PR10X2 was damaged by more extreme heating. After a hold at 400°C for one hour, the film would not crystallize even after a half hour at 250°C. After a quench to room temperature, regions resembling spherulites a few microns in size were visible. Another sample of the film was held at 350°C for one hour. This sample did not crystallize visibly after 45 minutes at 265°C. The sample was subsequently cooled to 250°C and spherulites, again in the size range of a few microns, formed over a three to five minute period. The effect of the extreme heating can slow crystallization in two ways. It can destroy active nucleation sites present in the melt and it can thermally cross-link the melt which can slow or even prevent crystallization. In either case, the condition of the film even after the milder melt treatment at 350°C indicated that this extreme processing is not feasible with the PPS resin even under argon. The film had darkened and bubbled even under the argon seal of the hot stage. The hot stage examination of the PR10X2 PPS resin indicates that the characteristic spherulite size is very small, probably on the order of one micron or less. Thus, the resulting morphology of these specimens will need to be investigated with a higher resolution technique than optical microscopy. 71 As suggested by the DSC results, the density of active nucleation sites in the PRO9 PPS resin is much lower than that of the PR10X2 resin. This leads to a spherulitic structure in a size range large enough to be observed with an optical microscope. As with the PR10X2 resin, conditions for the film in the hot stage matched those in the composite processing. Samples were held at 330°C for one hour before ramping to an isothermal crystallization temperature. As isothermal crystallization of the PR09 PPS proceeded, both nucleation and chain growth occurred continuously. During the first few minutes, only a few small spherulites were visible. These would keep growing and as they grew other spherulites would appear randomly distributed throughout the sample. This led to a distribution of spherulite sizes that ranged up to several dozen microns in diameter. Growing spherulites are shown at two different times during isothermal crystallization at 240°C in Figure 4.6 and Figure 4.7. The hot stage observations of the PR10X2 and PR09 PPS films in the absence of fibers indicate that the former is densely populated with active nucleation sites that crystallize rapidly to a size on the range of one micron or less. The PR09 has many times fewer active nucleation sites and crystallizes more slowly, leading to a structure with spherulites that may be several dozen microns in diameter. 1' P i The morphology that results from the crystallization of a semicrystalline thermoplastic, whether in a neat resin or a composite material, can have a marked 72 Figure 4.6 PRO9 PPS Spherulites Crystallizing at 240°C. Figure 4.7 Same Spherulites from Figure 6 at 240°C, 16 Minutes later. 73 effect on mechanical properties. The Confocal Scanning Optical Microscope (CSOM) allows the direct examination of crystalline morphologies without excessive sample preparation or etching techniques. The characteristic spherulite size of the PR10X2 resin was found to be at the lower edge of optical resolution. The nucleation density is so high that the spherulites appear to be less than one micron in diameter when viewed in the microscope. To verify this, both chemical and plasma etching, as described previously were used to bring out the crystalline structure of the PR10X2. SEM micrographs of plasma and chemically etched surfaces are presented in Figure 4.8 and Figure 4.9. In both cases, the surfaces have been severely etched, but what remains confirms that the characteristic spherulite size is 0.2 to 0.5 microns. The morphology of this resin, whether in an interphase or in the bulk, whether rapidly or slowly crystallized, remains constant. This leads to reproducible mechanical properties of the resin and rapid crystallization times for molded parts. As observed in the hot stage crystallization studies, the characteristic size of PR09 PPS spherulites is much larger. The sizes range up to 50 microns in diameter, depending on how soon during the crystallization process they nucleated and on how soon they impinged on other spherulites to stop their growth. A CSOM of neat PR09 resin is presented in Figure 4.10. 74 Figure 4.8 SEM of PR10X2 PPS Resin, Plasma Etched. Figure 4.9 SEM of PR10X2 Resin. Chemically Etched. 75 Figure 4.10 Confocal Micrograph of PR09 PPS Resin Crystallized at 250°C. 76 Cgstallization of PPS Resin in the Presence of Carbon Fibers The crystallization of PPS resin in the presence of carbon fibers can differ from that of the neat resin. The fibers can act as obstacles for the growing sperulites, with growth either stopping once the fiber surface is reached or continuing around the fiber. In addition, the fiber surfaces can act as nucleation sites and promote the growth of spherulites from themselves. If the distribution of active nucleation sites is dense enough, transcrystallinity may develop. In the case of the PR10X2 resin, the presence of fibers does not have a dramatic effect on the crystallization process. The spherulites are so densely nucleated in the resin itself that they are far more likely to encounter other spherulites to impair their continued growth than they are to encounter a fiber. Furthermore, the high density of nucleation sites in the interphase region of the resin prevents the development of transcrystallinity from the fiber due the high amount of competition for crystallizable material. The presence of fibers does, however, have an impact on the crystallization of the PR09 resin. The crystallization of PPS in the presence of A84, HMS4, 1M6, and a high modulus pitch fiber was observed. The tensile modulus, precursor, and relative crystallization nucleation activity observed are presented in Table 4.1. 77 Tghle 4,1. Relative PPS Crystallite Nucleating Activity of Carbon Fibers. Fiber Type Tensile Precursor Relative Crystallite Modulus (GPa) Nucleation Activity HMS4 345 Polyacrylonitrile highest 1M6 280 Polyacrylonitrile second highest (tie) I} AS4 230 Polyacrylonitrile second highest (tie) Pitch 825 Pitch lowest ll Of the four fiber types studied, only the HMS4 fiber provided a sufficient number of nucleation sites to lead to transcrystallinity. Spherulites growing from the HMS4 surface had impinged on adjacent sperulites after their diameter had reached about 5 microns. The A84 and 1M6 fibers did nucleate crystallization to some extent, but the majority of crystallites were still nucleated in the bulk resin and the crystallites growing from the fiber were too sparsely populated to grow into a transcrystalline layer. Spherulites growing from the fiber surface were still several dozen microns apart even after the diameter of the spherulites was more than 20 microns. Interestingly, the high modulus pitch fiber nucleated crystallization to the lowest extent in spite of possessing the highest tensile modulus of any fiber studied. The graphitic structure of the pitch fibers is quite different than that for the PAN fibers. The PAN fibers have graphitic basal plane strutures oriented longitudinally along their surfaces. The HMS4 fibers, which have undergone the most severe high temperature treatment, have straighter, longer basal plane edges at their surface than do AS4 or 1M6 fibers. The surface graphitic crystallites are on the order of 10 nm in length and are straight and well ordered for the HMS4 fiber, making the edges long 78 enough to support the 7 to 9 nm fold lengths of PPS lamellar chains. The initial growth of the lamellar structure should therefore be perpendicularly outward from the fiber surface. The graphitic basal plane structure tends to be oriented axially rather than longitudinally along the fibers for pitch fibers, leaving no exposed basal plane edges for epitaxial activity. This suggests that fiber surface morphological structure and epitaxy may indeed play a role in the nucleation of polymer crystallization and that the tensile modulus alone is not an important factor in determining nucleation ability of carbon fiber surfaces. It is the straight, long graphitic basal surface plane edges that are responsible for the high degree of PPS crystalline nucleation on HMS4 fibers. h l f i in h n f n Fi The morphological studies of PPS resins in the presence of carbon fibers was limited to the PR09 and amine terminated PPS. As mentioned previously, the abundance of active bulk nucleation sites in the PR10X2 resin prevents the fibers from having an impact on resin crystallinity. Confocal microscopy was performed on composite specimens made from film stacked resin and fiber tow as described earlier. Sections of composite were embedded in polyester casting compound and polished to produce a flat surface for microscopic investigation. The PR09 samples crystallized at 235°C for one hour after the standard one hour hold at 330°C had crystalline material throughout the sample. The spherulites had grown to diameters of from 10 to 50 microns, depending on the number of other 79 spherulites in the region. The PR09 samples crystallized at 250°C and 265°C for one hour did not have sufficient time to crystallize to such a complete extent. Both contained spherulites of from 5 to 15 microns in diameter. The different types of fibers were found to nucleate crystallinity just as observed in the hot stage microscopy work. The AS4, 1M6, and high modulus pitch fiber samples had some spherulites nucleated from their surfaces, but most crystallinity clearly came from the bulk. This is most easily viewed with a sample crystallized at 265°C, such as the AS4 sample presented in Figure 4.11. In this picture, where the spherulites have not yet grown to impingement, it is obvious that the spherulites nucleated in the bulk resin dominate and that the spherulites nucleated from the fiber surface are not abundant enough to produce transcrystallinity. In Figure 4.12, another AS4 sample is presented. This sample was crystallized at 235°C so that the spherulites were allowed to grow to impingement in their one hour of crystallization time. This image also clearly shows that the majority of the spherulites were nucleated in the bulk resin and not on the fiber surface. The HMS4 fibers, on the other hand, were again found to lead to a transcrystalline interphase morphology. The rate of growth of the transcrystalline layer was equivalent to that of concurrently growing spherulites. This structure is seen to have developed for all crystallization temperatures studied (235°C, 250°C, and 265°C). A transcrystalline layer was also observable for a sample that was held at 330°C for an hour and quenched to below T, in a 2-3 minute period. This layer, which was approximately 15 microns in diameter, formed rapidly under nonisothermal 80 ’ -- . . J - ZOOH'ZI T'ZI C3119 IISII Fl L488 Figure 4.11 AS4 Fiber/PRO9 PPS Resin Crystallized 1 hour at 265°C. ZOOH'Z. T'ZS Figure 4.12 AS4 Fiber/PR09 PPS Resin Crystallized 1 hour at 235°C. 81 ZOOIIZO T'Zs C'll9 I354} F8 L488 Figure 4.13 HMS4 Fiber/PR09 PPS Resin Crystallized 1 hour at 265°C. . fl ‘. _ HMS4 Fibers/Stan. PPS, Cryst. 1 hr at 235°C, 5 am deep A HMS4 Fibers/Stan. PPS. Cryst. 1 hr at 235°C. 7 pm deep Figure 4.14 HMS4 Fiber/PR09 PPS Resin Crystallized at 235°C for 1 hour. 83 conditions. An HMS4 sample crystallized at 265°C is presented in Figure 4.13. The "optical sectioning” capabilities of the confocal microscope are used in Figure 4.14. In this series of images, the depth of analysis was increased one micron at a time, in effect stepping through the transcrystalline regions of the two fibers shown. This verifies that the transcrystalline region has actually formed around the fibers and is not above or below the fiber. The amine terminated PPS led to morphologies very much like the PR09 resin. An HMS4 amine terminated PPS transcrystalline sample is shown in Figure 4.15 and an AS4 amine terminated PPS crystalline sample is presented in Figure 4.16. The morphologies for all of the composite samples are summarized in Table 4.2. 84 Figure 4.15 HMS4 Fiber/Amine Terminated PPS Crystallized 1 hr at 250°C. Figure 4.16 AS4 Fiber/Amine Terminated PPS Crystallized 1 hr at 250°C. 85 Tehle 4,2 Interphase Morphologies for Fiber reinforced PPS samples. F Sample Type Cryst. Cryst. Interphase Bulk Temp Time Morphology Spherulite (°C) Diameter PR09/AS4 235°C 1 hour bulk nucleated 10—50 a crystals PR09/AS4 250°C 1 hour bulk nucleated 10—30 a crystals + amorphous regions PR09/AS4 265°C 1 hour Bulk nucleated 5-20 is crystals + amorphous regions PR09/AS4 330°C-- 2-3 minute Amorphous + 5-10 )4 85°C quench some spherulites PR09/HMS4 235°C 1 hour Transcrystalline, 10-50 a 2050 it thick PR09/HMS4 250°C 1 hour Transcrystalline, 10-30 a 2050 p. thick PR09/HMS4 265°C 1 hour Transcrystalline, 5-20 a 20-30 a thick PR09/HMS4 330°C-- 2-3 minute Transcrystalline, 5-10 a 85°C quench 20—25 a thick PR10X2/AS4 or 250°C 1 hour Bulk nucleated 0.3-0.5 a II HMS4 small spherulites Amine terminated 250°C 1 hour Bulk nucleated 1050 Il- PPS/AS4 spherulites Amine terminated 250°C 1 hour Transcrystalline 10-50 It PPS/HMS4 20-50 a thick II 86 NL IN: Active crystalline nucleation sites are much more plentiful in the composite grade PR10X2 resin than in the amine terminated or PRO9 PPS resins. This is evident by the much more rapid crystallization of the composite grade resin and the much smaller characteristic crystallite size. The bulk morphology of all resins studied was found to be spherulitic for both neat and fiber filled resins. The characteristic spherulite diameter for the composite grade resin was less than 1 micron and that for the PR09 and amine terminated resins was several dozen microns. The extremely high nucleation density in the composite grade PR10X2 PPS resin eliminated the possibility of transcrystallinity. Any crystallizable material in the interphase became incorporated in bulk crystallites before fiber associated nucleation could occur. All of the PPS resins studied exhibited dual melting endothermic peaks for samples that were crystallized isothermally. This indicates the presence of both primary and secondary crystalline regimes for all resin types studied. Extremely severe thermal treatment was found to inhibit crystallization of PPS, even when the treatment was conducted in an argon environment. This can be attributed to both the destruction of nucleation sites and thermal crosslinking of the resin, both of which inhibit crystallization. Of the PAN based fibers studied, the highest modulus HMS4 fibers nucleated crystallinity the most effectively, leading to transcrystalline morphologies with both the PR09 and amine terminated resins. The lower modulus A84 and 1M6 fibers, which contain less graphitic crystalline plane structure, nucleated a small degree of 87 crystallinity, but not enough to lead to transcrystallinity for any of the conditions studied. The pitch fibers, which had the highest tensile modulus of any of the fibers studied nucleated crystallinity the least. This indicates that for this system, the fiber surface graphitic morphology is a more important factor in determining nucleating ability than is fiber modulus. This suggests an epitaxial mechanism for the nucleation of PPS by carbon fibers. A ER V A i n f P rbon Fibers The level of adhesion between a semicrystalline thermoplastic such as polyphenylene sulfide (PPS) and reinforcing carbon fibers in a composite can be the result of interphase morphology, composite processing, fiber and matrix mechanical and surface properties, and the extent of matrix to fiber chemisorption. In this paper, values of interfacial shear strength are reported for PPS/carbon fiber composites. In addition, the effect of the previously mentioned potentially important physical and processing factors are discussed. The chemical reactivity between PPS compounds and carbon fibers was determined by subjecting the fibers to PPS functional groups at processing conditions and examining the surface chemical compositions of the fibers with X-ray Photoelectron Spectroscopy (XPS) before and after processing. The interfacial shear strengths of the resin/fiber interphases were measured in situ using an indentation system. I I N: hmi IR i' w’h nFir rf The XPS surface composition results are reported in Tables 5.1 and 5.2. XPS measurements represent the outer 40 A of the fiber, limited by the mean free path of the ejected electrons. 88 89 m1 XPS Results for AS4, HMS4, and Pitch Fibers. F Fiber Type Na (%) O(%) N (%) S (%) AS4, ar’d 0.3 i 0.1 5.1 :1; 1.1 2.1 :1; 0.4 0.2 :1; 0.1 AS4, bp 0.4 i 0.1 5.0 i 1.0 1.8 3; 0.1 0.2 :1; 0.1 AS4, dsp 0.4 i 0.1 3.7 i 0.9 1.8 -_i- 0.2 0.1 i 0.1 HMS4, ar’d 1.0 i 0.2 3.5 :1; 0.9 0.5 i 0.4 0.1 :1; 0.1 HMS4, bp 0.8 i 0.1 3.0 i 1.0 0.4 :1; 0.3 0.1 :1; 0.1 HMS4, dsp 0.8 :1: 0.1 2.7 i 0.6 0.2 i 0.1 0.1 i 0.1 Pitch, ar’d 0.1 i 0.1 4.0 i 0.8 0.4 :1; 0.1 0.1 i 0.1 Pitch, bp 0.2 :1; 0.1 2.7 i 0.5 0.3 :1; 0.1 0.1 :1: 0.1 Pitch, dsp 0.5 :1; 0.3 4.3 i 0.4 0.5 i 0.4 0.1 :1: 0.1 (ar’d = "as received”, bp = ”biphenyl processed”, dsp = "diphenyl sulfide processed”) m XPS Results for 1M6 Fibers with 0%, 100%, and 600% surface treatment. 90 Fiber Type Na (%) O (%) N (%) S (%) 0% ST, ar’d 0.3 i 0.1 3.0 j; 1.0 1.0 :t 0.2 0.1 :1; 0.1 0% ST, bp 0.5 i 0.1 3.0 :1; 1.0 1.4i 0.9 0.1 i 0.1 0% ST, dsp 0.5 i 0.1 2.3 i 0.3 1.0 i 0.1 0.1 :l: 0.1 100% ST, 0.2 i 0.1 4.2 i 0.7 1.1 :l: 0.3 0.1 i 0.1 ar’d 100% ST, bp 0.9 :1; 0.1 4.9 :1; 0.7 1.1 j; 0.1 0.2 i 0.1 100% ST, dsp 1.2 j; 0.1 5.0 :1; 1.0 1.1: 0.1 0.2 i 0.1 600% ST, 0.5 j: 0.3 13.0 i 0.3 3.5 :1; 0.3 0.2 :1; 0.2 ar’d 600% ST, bp 1.4 :1: 0.3 12.0 i 3.0 1.7 i 0.3 0.4 i 0.3 600% ST, dsp 1.3 i 0.2 12.0 :1; 1.0 1.6 :1; 0.2 0.4 :1; 0.1 The processing at high temperatures and subsequent refluxing in benzene ensured that any changes in surface composition would indicate that material was chemically bound to the fiber surface and not simply physically adsorbed onto it. The results for all fiber types studied indicate that there is no chemical bonding taking place between diphenyl sulfide and carbon fibers. The results for the PPS tetramer indicated the same. Therefore, chemical bonding between the polyphenylene sulfide and any fiber surface has not taken place and can be eliminated as a contributing 91 factor to fiber-matrix adhesion for PPS/ carbon fiber composites. Any differences between adhesion for different fibers must be attributed to other factors. rfilhr nhfP/anFirmi 1 X2 m it I Interfacial shear strengths for PPS/carbon fiber composites were measured with the ITS. The interfacial shear strength results for all the fiber and resin types, and for various processing conditions, are summarized in Figures 5.1-5.4. The effect of fiber type on adhesion of composite grade PPS (PR10X2) resin to carbon fibers is illustrated in Figure 5.1. The composites were processed for one hour at 330°C followed by one hour at 250°C, with 1.03 MPa pressure maintained throughout the processing. The bulk and interphase crystalline morphology is bulk nucleated spherulites 0.3-0.5 microns in diameter for all of these samples. All of the composites have a tough, nontranscrystalline interphase and these interfacial shear strengths represent the highest levels of adhesion attainable given the limitations of the carbon fibers. In comparing between different types of fibers, it is important to keep in mind that the interfacial shear strengths measured may be limited by the shear strengths of the fiber surfaces themselves. Higher modulus fibers tend to have lower shear strengths, that may lead to failure of the interphase before the mechanical bond between fiber and matrix would otherwise be destroyed. For the series of PAN based fibers, the adhesion of AS4 fibers was the highest, followed by 1M6 fibers, with the HMS4 adhesion being the lowest. SEM examination of composite samples broken in 92 Interfacial Shear Strengths l 10 r A i- : 100 . l” 2 90 ~ , ,, z 1 .. so I 5 g l a; o 70 l' :5: ’4: b ’ :/, ,1. en 60 ~ 21' i... ; T D 50 "‘ y; / E/ d _ 40 ~ .5 3/ 2:"? .3 t '2: 2;"? 0 b // " '3 .2 3° , r": 2.19 :3 20 r / "./ t: V .- /'/" — l 0 " ,v‘; .1]. f ', ,- '/ ’ 4 a i; 7/ i o 414’ 5 m A54 “[6 HMS4 Pitch Fiber Type Figure 5.1 Interfacial Shear Strengths of PR10X2 PPS with Carbon Fibers. Interfacial Shear Strengths no . 7 100 ~ 0- l 5 90: 5 80 ~ g 70 8 T 7/ 63 60 ~ g ~ e 3 50 - / = » r E”; 40 » g ‘2: 30L % a r 5 20 ~ g 5 10 ~ g 02 202 1001 2002 6002 “Extent of Surface Treatment Figure 5.2 Interfacial Shear Strength of PR10X2 PPS Resin with M6 Fibers. 93 l 10 1; 100 C... x 90 '5 so I 5 7o 1: w 60 ‘i o 50 .2 T. 40 .2 E 30 E 20 5 10 Interfacial Shear Strengths r PR 1 OX2 PPS * i I . 7,}; Amine Term. / , i g», PPS . 532 r I r ’/ i r. // i- V A r .4 ’i": L {’5 6’4 . : 3 {’33 #52 :2, 2a: ‘42; 75 /5: t 5.5. 5 4 '5 ’/ '1 r, ,v r, / , -:. 5 5w f '03; x5. .5, 5 i, [/z '71: ’/5,'/ i. {l ’5 '/" * M 2 5. 32", t .2, '22: I. r r ‘l 1;. 1.” l/ 250 Quench 23 265 Crystallization Temperature ( C) _ “fl“-.-- 4 Figure 5.3 Interfacial Shear Strengths of PR09 PPS with AS4 fibers. l 10 100 90 80 7O 60 4O 30 20 10 Interfacial Shear Strength (MPa) 50b 1 fiY PR1 Interfacial Shear Strengths 0X2 PPS % E ———.._ 250 Quench 235 250 265 280 Crystallization Temperature ( C) Figure 5.4 Interfacial Shear Strengths of PR09 PPS with HMS4 fibers. 93 Interfacial Shear Strengths 0 p ‘ ' PR 1 0x2 PPS :- 100 i 90 r "a 5 E 2 . .r: 80 , 3‘72 Amine Term. d /'-, I 1 7': PPS C I”: D 70 l» ’3‘ l 5 l m 50 » h ' r/ (A; 8 50 ~ .27; a I 1”,? .5, - 40 * 2'55 I ‘ {155. ° 30 r 271 55,” he I :0?’ g: :3 20 > :55 2;; E l E: A"! :61?!) " ‘0 i 502 O L‘ if? 250 Quench 235 265 Crystallization Temperature ( C) -‘--.—-_._—n—.«.- - Figure 5.3 Interfacial Shear Strengths of PR09 PPS with AS4 fibers. 1 10 100 90 80 70 60 40 30 20 10 Interfacial Shear Strength (MPa) 50* PR1 Interfacial Shear Strengths OX2 PPS i .ii . 250 Quench 235 265 280 Crystallization Temperature ( C) Figure 5.4 Interfacial Shear Strengths of PR09 PPS with HMS4 fibers. 94 a transverse flex mode revealed that the failure in all cases was adhesive. At the locus of failure, the fibers pulled out cleanly with very little or no PPS matrix left on their surface. The failure of the pitch sammes, however, was an entirely different situation. As illustrated in Figure 5.5, failure of the interphase region in this samples was a cohesive failure of the pitch fiber itself. As seen in the SEM micrograph, a surface layer of the fiber remains strongly adhering to the matrix with failure taking place within the fiber. This explains the remarkably low adhesion levels measured for the pitch samples with the ITS. The system was detecting low cohesive strength of the fibers themselves rather than a low level of adhesion between fiber and matrix. The series of 1M6 fibers, shown in Figure 5.2, also provides some useful information. With this series of fibers, the only difference from one type to another is the amount of electrolytic surface treatment to which each fiber was exposed. There is a significant increase in adhesion upon going from untreated fibers to fibers treated at 20% of the standard amount of treatment time. This increase agrees with results previously obtained by Waterbury and Drzal'. It has been noted that an increase in treatment time increases the amount of functional nitrogen and oxygen on the fiber surfaces. Using scanning tunneling microscopy, Sugiura2 has shown that the surface roughness also increases with treatment time. This increase in surface roughness, which is only clearly resolved at very high levels of magnification, may lead to increases in surface area of as much as 60% when the treatment level is increased from 0% to 600% . It is likely that this increase in surface energy is the primary factor in increased adhesion levels for fibers with more surface treatment. 95 figure 5 5 SEM of PPS / Pitch Fiber Composite Failure. 96 $92 and Amine jlfgrminatgd 225 Both the PR09 and amine terminated resin have a morphology that includes large spherulites. In addition, both resins exhibit transcrystallinity when combined with high modulus carbon fibers. With the effect of surface treatment and fiber type established, Figure 5.3 shows the effect of crystallization conditions on interfacial shear strength with AS4 fibers. The processing dependent studies were conducted with the PRO9 extrusion grade resin (low nucleation density) because the PR10X2 composite grade resin (high nucleation density) exhibits a morphology that is extremely insensitive to processing conditions. The composite grade resin crystallizes into small crystalline spherulites rapidly upon cooling from the melt process temperature to below the glass transition temperature. The crystallization of PRO9 resin, on the other hand, is dependent upon the crystallization temperature. The level of fiber-matrix adhesion, therefore, can be strongly influenced by the isothermal crystallization temperature. The PR10X2 adhesion value is shown first for comparison purposes. It is clear from this figure that the levels of adhesion with the PRO9 resin with the AS4 fibers which do not lead to transcrystallinity never approach the high value for the PR10X2 resin. A micrograph of a typical bulk nucleated crystalline sample is presented in Figure 5.6. This can be attributed to the lower toughness of the matrix interphase for the more 97 ZOOIIZI TIZI C31“ It!“ F. L“. Figure 5.6 Bulk Nucleated Crystalline Interphase (PRO9 PPS/AS4). "9 ‘ " saw HMS4 Fibers/Stan. PPS, Cryst. 1 hr at 235°C, 6 pm deep Figure 5.7 Transcrystalline Interphase (PRO9 PPS/HMS4). 98 coarse spherulitic structure, which has few tie molecules and many large interspherulitic regions with lots of imperfections. The sample that was quenched from 330°C to below the glass transition temperature in 2 to 3 minutes had a very low measured interfacial shear strength, compared to those held one hour at 235°C and 265°C, which had similar but much higher interfacial shear strengths. The interphase in the quenched case was amorphous with a few 5-10 micron spherulites and formed very rapidly, allowing little time for organization and chain equilibration as the melt solidified. Confocal microscopy and DSC investigations reveal that the 235°C sample is essentially completely crystallized, while the sample held at 265°C is still mostly amorphous due to the slower crystallization rate at the lower degree of supercooling. This indicates that at temperatures below the glass transition temperature, a primarily amorphous interphase, which has a chance to form during the hold at 265°C and a crystalline (but not transcrystalline) interphase formed at 235°C with large spherulites produce the same level of fiber-matrix adhesion. The amine terminated resin exhibited the same levels of adhesion as the extrusion grade PRO9 resin. The morphology in this case was similar, with large spherulites impinging upon the fiber surfaces. The amine end group appears to have little or no interaction with the AS4 fiber surface. In addition to AS4 carbon fiber/ PRO9 samples, HMS4 / PRO9 samples were also tested. The interphases for these specimens were transcrystalline and the results are displayed in Figure 5.4. Again, the PR10X2 composite grade result is displayed first for comparison purposes. Just as with the AS4 fibers, there is a large drop in interfacial adhesion for the quenched sample. The adhesion falls by a factor of 5 for 99 the HMS4 and by a factor of 9 for the A84. Confocal microscopic examination of the sample revealed that there was some surface nucleated crystallinity around many of the AS4 fibers, and that a transcrystalline region had developed for the HMS4 fibers. The fact that both transcrystalline and primarily amorphous interphases are very weak when no isothermal hold below the crystalline melting point is used indicates that contact between matrix and fiber when the matrix is in the rubbery state is essential to developing high adhesion levels. The next three values displayed are for different isothermal crystallization temperatures. All three of these samples exhibited a full transcrystalline morphology. Figure 5.7 shows a transcrystalline PPS interphase. The interfacial shear strength shows a significant increase for the slowly crystallized sample (265°C) over the more rapidly crystallized samples (235°C and 250°C). This dependence, which was not observed for the non-transcrystalline AS4 samples, shows that there is an effect of the rate of transcrystalline interphase formation on interfacial shear strength. In fact, the interfacial shear strength of the PRO9/HMS4 interphase was just as high or possibly higher than that of the PR10X2/ HMS4 interphase. To show that this increase was not simply a result of longer hold time at higher temperature, a sample was held at 280°C . This is the last result shown in Figure 5.4. The low value is similar to that for the quenched sample, indicating that a change in resin structure or composition at high temperature was not responsible for the high interfacial shear strength observed for the sample crystallized at 265°C. Again, it appears that a hold in the rubbery state, rather than only the liquid state (above the crystalline melting point) is crucial for developing satisfactory adhesion. 100 QQNQLQSIONS; High temperature processing studies of carbon fibers with model compounds for PPS indicate that there is no chemical bonding between fibers and PPS in a composite material. Interfacial shear strength measurements with a series of 1M6 fibers show that fiber surface chemistry, and surface roughness can be factors in determining the adhesion level reached in a composite. An amine endgroup attached to a PPS chain has no special interaction with the carbon fibers studied, as evidenced by the lack of measurable increase in interfacial shear strength when the group is included. The composite grade PPS had the highest interfacial shear strengths of the resins studied. The formation of a tough interphase consisting of fine bulk nucleated spherulites is a key in developing strong fiber-matrix adhesion. In the absence of transcrystallinity, the composite grade resin led to higher interfacial shear strengths than did the extrusion grade for all processing conditions studied. Processing conditions are impoflant in determining the interfacial shear strength of PPS/carbon fiber composites. Samples that were quenched from the melt processing temperature of 330°C developed very low fiber-matrix adhesion. Samples held at temperatures below the crystalline melting peak but well above Tg for an hour developed higher fiber-matrix adhesion. Composites with a transcrystalline interphase were particularly sensitive to processing conditions. Rapidly grown transcrystalline interphases, which formed at high degrees of supercooling, had lower interfacial shear strengths than did slowly grown transcrystalline regions. A more slowly grown transcrystalline layer should be 101 more free from crystalline defects. This is critical, particularly for the secondary crystallization that occurs in between and within spherulites, as demonstrated by Cebe’, who showed that the secondary crystallization peak could actually be shifted to higher temperatures by crystallizing at lower supercooling. In addition, slow growth of the crystalline lamellae from the fiber surface allows more time for the adjacent amorphous layers to establish good adhesion to the fiber surface. This is especially important since the lamellar growth direction is perpendicular to the fiber surface, which limits the contact of the crystalline component with the fiber surface. The interfacial shear strength of PPS/pitch fiber composites was severely limited by a weak surface layer on the pitch fiber itself. A weak surface layer may also limit the interfacial shear strength of HMS4 fiber composites. W 1. L. T. Drzal, M. Madhukar, and M. C. Waterbury, Composite Structures, 27, 65 (1994). 2. N. Sugiura and L. T. Drzal, unpublished work. 3. P. Cebe and S. Chung, Polymer Composites, 11, 265 (1990). YAND NL IN This investigation of extrusion (PRO9) and composite grade (PR10X2) PPS in composites with carbon fibers was an examination of interphase properties and their effect on interfacial shear strength. Surface Energetics / Wetting The surface free energies of all of the PPS resins studied are very similar. This indicates that none of the resins should adsorb more strongly to fibers during processing than the others. It also shows that low surface free energy impurities do not concentrate at surfaces or interfaces for any of the PPS resins during the composite processing cycles. The adhesion of PPS to 1M6 carbon fibers was found to increase with fiber surface treatment level which raised the fiber surface free energy and increased the surface area. Both of these factors contribute to higher interfacial shear strengths. Fiber-Matrix Chemical Reaction It was demonstrated that chemical reactions do not take place between PPS and carbon fiber surfaces at composite processing temperatures. This eliminates interfacial chemical bonding as a factor in PPS-carbon fiber adhesion. It was also shown that the addition of a reactive amine endgroup to the PPS chain does not lead to higher adhesion levels as it does for more active substrates. 102 103 Interphase Crystalline Morphlogy The composite grade PPS has a finer spherulitic (0.3-0.5 micron diameters) morphology than does the extrusion grade PPS (10-50 micron diameters). This morphology is responsible for the formation of tough, bulk nucleated crystalline interphases. The matrix is tougher due to the greater number of tie molecules, which are chains included in the lamellar structure of more than one spherulite. These shared chains strengthen and toughen the inherently weak interspherulitic amorphous regions. Smaller spherulites also lead to smaller interspherulitic regions and more tortuous paths for crack propagation, which also increase interphase toughness and consequently interfacial shear strength. It is this finer, tougher crystalline interphase morphology that leads to the higher adhesion levels for PPS/AS4 composites made from composite grade resin. Of the fiber types studied only the HMS4, which has a highly oriented graphitic crystalline structure, led to the formation of PPS transcrystallinity. The epitaxial effect of the edges of the graphitic basal planes was demonstrated to be the major cause for the transcrystallinity. Pitch based fibers, which have an even higher modulus than the HMS4, but lack the surface graphitic basal plane sites, did not lead to transcrystalline interphases. Transcrystallinity is not a factor for the composite grade PPS resin due to the high number of active nucleation sites present in the matrix. The isothermal crystallization temperature was found to have a marked effect on interfacial shear strengths for composites with transcrystalline interphases. Slowly grown (at high temperature) transcrystalline regions led to interfacial shear strengths 104 comparable to those measured for the tough composite grade PPS resin. Transcrystalline regions grown more rapidly (at lower crystallization temperatures) led to weaker adhesion. This can be attributed to the more perfect crystalline structure that is formed when the transcrystalline front advances slowly. More imperfections are likely for rapidly grown layers because the higher energetic driving forces can overcome chain imperfections that would not be overcome at higher crystallization temperatures. The slower growth rate also allows the polymer chains of PPS more time to establish optimum conformation with the fiber surface. In addition, the secondary crystallization regimes will be more ordered and stronger. A dwell time at a temperature below the crystalline melting peak but above the glass transition temperature was shown to be important in establishing high adhesion levels for all types of interphases (mostly amorphous, bulk nucleated crystalline, and transcrystalline). Specimens quenched from the liquid melt temperature (330°C) to below the glass transition temperature (85°C) without a hold at a temperature somewhat below the melting peak (265°C or lower) exhibited extremely low interfacial shear strengths. Specimens held at a temperature below the melting peak but well above the glass transition temperature exhibited much higher fiber-matrix adhesion. Similar results have been observed by other researchers with an amorphous polymer (polycarbonate) so it is unlikely that this phenomenon is directly related to the crystalline structure. 105 Overall Summary The improved adhesion of composite grade PPS to carbon fibers is due the resin‘s fine crystalline structure. This leads to the formation of tough, crack resistant interphases in PPS/carbon fiber composites. Other factors that were considered, including improved adsorption onto fibers during processing, chemical reactions between fiber and matrix, and the elimination of interphase active impurities, were shown to be of no importance. This suggests that in formulating semicrystalline polymer composites, producers should aim for a fine crystalline structure which will lead to tough, strong fiber-matrix interphases. At temperatures below the glass transition of PPS, interfacial shear strengths of bulk nucleated crystalline composites were shown to be insensitive to the degree of crystallinity or the isothermal crystallization temperature. The interfacial shear strengths of transcrystalline composites were found to be extremely sensitive to crystallization temperature. This may explain some of the discrepancies in the literature, where some researchers claim that transcrystalline interphases lead to stronger fiber matrix adhesion while others argue that transcrystallinity leads to lower interfacial shear strengths. In this research, transcrystalline composites that were crystallized at high degrees of supercooling had significantly lower interfacial shear strengths than bulk nucleated composites which had a tough, fine spherulitic interphase. On the other hand, transcrystalline composites that were crystallized at low degrees of supercooling developed roughly the same interfacial shear strength as the tougher composite grade resin. It was not shown, however, that transcrystallinity is a property that leads to significantly higher 106 interfacial shear strengths. The results discussed in this work suggest that composites with transcrystalline interphases should be crystallized slowly at low degrees of supercoolin g. W Pendan Dr Procedu and am le alculatio The pendant drop surface free energy measurement apparatus allows surface free energy determinations for liquids in an inert environment at temperatures up to 350°C. This procedure should be read and understood before any measurements with the system are attempted. Sample Preparation If the material for which the surface free energy measurement is required is a solid at room temperature, it will have to be melted and loaded into a clean capillary as a liquid. Glass micropipets from Drummond Scientific should be cleaned with deionized water and Chromerge' solution and dried in a clean oven that is reserved for drying clean glassware. Use of an oven which is used for other purposes (curing epoxies, drying thermoplastic pellets, etc.) could contaminate the glass surfaces and defeat the purpose of cleaning the capillaries in the first place. The sample capillaries used in this work were filled with molten PPS in the following manner. First, an aluminum heating block wrapped in a heating tape was wrapped in insulation and heated to 400°C. The temperature was controlled at the heating tape on the outside of the block. Once the block was up to the required 107 108 temperature, clean glass culture tubes were placed into 3/4" holes drilled into the aluminum block. The tubes were loaded with PPS film or granules and sealed with aluminum foil. A small purge flow of argon was maintained through the top of the culture tube to prevent oxidation of the PPS during heating. After the PPS had thoroughly melted a clean capillary connected to a vacuum pump was inserted below the melt surface in the culture tube. The vacuum pump was started and molten PPS was drawn up into the clean capillary. Once the capillary was full, the pump was turned off and the filled capillary was detached from the vacuum tubing. In order to obtain equilibrium surface free energy measurements, as many of the bubbles that form in the melt due to volume shrinkage upon melt solidification must be removed. To this end, the filled capillary was placed in an empty culture tube in the heating block. An argon purge was again supplied to the culture tube to prevent oxidation of the polymer. Once the polymer inside the capillary had remelted, a teflon tipped syringe plunger (supplied with the Drummond capillaries) was inserted into the end of the capillary. The molten polymer was compressed and cooled in a compressed state to remove as many voids as possible. This was not sufficient to remove all of the bubbles. Any pendant drops which have voids in them which do not go away during the drop equilibration process will not come to equilibrium or will not give an accurate value and must not be used for analysis. The surface tension calculation involves the ratio of the surface free energy to the molten polymer density. Voids will alter the melt density and throw off the surface tension calculation. 109 Filled capillaries for pendant drop measurement are held in the cell with a set screw cushioned with a small piece of teflon and by a lip at the end of the capillary itself. The empty ends of the loaded capillaries should be fluted out to a small lip by a glass technician. Sample Loading and Drop Cell Heating The loaded capillary should be placed into the holder at the end of the linear motion feedthrough on the top of the drop cell. The set screw at the side of the capillary should be cushioned with a small piece of teflon and tightened just enough to hold the capillary in place during the measurement. Tightening the screw too much will break the capillary and ruin the sample. The capillary should be checked to make sure that it is parallel with the holder rod on the motion feedthrough. Once the sample is loaded into the motion fwdthrough and is parallel with the holder rod, the feedthrough can be placed on top of the drop cell. Do not bolt the fwdthrough on at this point. The end of the capillary should be at about the middle of the Viewports in the drop cell. If the drop forms at an off center position, the temperature of the drop may not be accurately measured. The thermocouple which measures the drop temperature should be located directly next to the drop (it should be viewable or just offscreen when the drop is viewed in the computer monitor). This will ensure an accurate drop temperature measurement. The cell must first be purged with argon. Allow argon to flow into the cell and out the top of the cell through the linear feed attachment flange. After purging for 5 minutes, bolt the linear feed onto the cell, using a fresh copper gasket. Once 110 the cell is sealed, it should be pumped down to vacuum and refilled with argon several times. The pumping can be accomplished with the attached sorption pump. The pump will work more efficiently if it is baked out with the supplied bakeout heater before use. After the cell has been pumped down and refilled with argon several times, close both of the large all metal valves. Heating of the drop area in the environmental cell is accomplished with coiled cable heaters (275 Watts each) located above and below the drop formation area. The upper heater is primarily responsible for melting the polymer in the capillary and the lower heater for heating the drop formation area. Each heater can be controlled independently with its own controller and thermocouple loop. The outputs from the controller can be varied with the variacs to which each heater is connected. The heaters share a common ground in the electrical feedthrough at the bottom of the cell. Be sure that the neutral for each heater is the line connected to the common connection. The drop cell must be heated slowly in small increments (25°C or less at a time). Due to the space between the thermocouples and the heaters, there is a thermal lag which can lead to dangerous overheating if the temperature is increased too rapidly. The best method is to turn up the heaters 25°C at a time and wait for the heater to equilibrate at a stable temperature before turning up the setting again. Once the cell is at the desired temperature, the linear motion feedthrough can be advanced to push out enough material to form a hanging (pendant) drop. The first drop should be allowed to fall without using it for analysis as it may contain impurities and is most likely to have suffered some oxidation in heating. The material which falls from the capillary is collected in a culture tube at the bottom of the cell 111 which should be below the bottom heater. This tube must be removed and replaced periodically to prevent system contamination. Viewing the Drop and Recording Images The pendant drop images are captured with the frame grabber on the Amiga microcomputer which is dedicated to the system. The input to the frame grabber comes from a solid state video camera. The camera system is mounted on an optical bench so that image distances can be precisely reproduced. To record a drop image, focus the camera on the pendant drop. This is most easily done while running the frame grabber program on the Amiga. Selecting "Live Amiga” from the menu will provide a live view of the drop. Once the drop edge is sharply in focus, fix the camera securely in place. Images should be grabbed every five minutes (or another suitable time frame) until the drop is at equilibrium. The equilibration time will depend on the liquid used. On the first run, it is a good idea to allow an hour or two to be sure. As mentioned previously, any bubbles that are in the droplet make any measurements highly suspect and such droplets should not be used for analysis. After the all of the desired images of the drop have been obtained, the cell should be allowed to cool to room temperature. In order to calibrate the drop images and determine their actual size, an image of an optical grid or reticle must be grabbed at the same magnification and focal length as were the drop images. Distances can be measured in terms of pixels with the paint program on the Amiga and converted to actual metric dimensions using the optical reticle image. Alternatively, edge points 112 can be recorded and stored in a file for other image analysis. The Amiga is equipped with an IBM bridgede which can convert Amiga files to IBM readable files if use of an IBM analysis package is desired. Sample Calculations The surface free energies reported in this work were calculated using the tabulated Bashforth Adams shape factors presented in Physical Chemistry of Surfms by Adamson. This calculation is for PR10X2 PPS at 300°C. The calibration is listed on the raw data diskette as cgpps300.01c. The final drop image is listed as cgpps300.01. First, the calibration grid image was loaded into the paint program. Horizontally, 7 grid divisions (equivalent to 0.35 cm) took up 557 - 48 pixels. This makes the horizontal calibration: H = (557-48)pixels/0.35 cm = 1454.29 pixels/cm. Vertically, 5 grid divisions(equivalent to 0.25 cm) took up 379 - 63 pixels. This makes the vertical calibration: V = (379-63)pixels/0.25 cm = 1264 pixels/cm. 113 Next, the drop image file is loaded into the paint program. A schematic of a drop image is presented in Figure A.1. First, the equatorial diameter (dc , the widest portion of the drop) is measured in terms of pixels. dc = (357-18)pixels/(1264pixels/cm) = 0.2681962 cm The other diameter of interest is d, , the diameter a distance (I, up from the diameter d,. d, = (334-42)pixels/(1264 pixels/cm) = 0.2310127 cm and S = d,/de = 0.2310127cm/0.2681962 = 0.8613569. The shape factor, HR is listed in the table as a function of S. Linear interpolation gives l/H = 0.4649361. The densities of molten PPS as a function of temperature as determined by Craig Bero at the University of Pittsburgh are in a folder near the apparatus. At 300°C, the density of molten PPS is 1.185 g/cm’. The surface free energy can now be calculated as: ‘y = (1/I~I)Apgd,2 = (0.464936l)(980 g°cm/sec2)(0.2681962 cm)2 = 38.8 dyne/cm. Drops were taken to be at equilibrium when 3 or more consecutive measurements 5 minutes apart yielded values of 7 that were constant. Drops with air bubbles can not be used for analysis even if they appear to be at equilibrium. 114 Figure A.l Schematic of Pendant Drop. 115 PPS Pendant Drop Raw Data Files The pendant drop raw data files used for this work are stored on the diskette labelled ”PPS RAW DATA FILES". The composite grade PPS files are named cgpps300.01, .02, etc. for the 300°C values and cgpps330.01, .02 etc. for the 330°C values. The amine terminated files are named amine330.01, .02, etc. The PR09 files are named ryton330.01, .02, etc. Each file has a companion calibration grid file with an identical name but with the letter "c" at the end of the extension. Checking Technique Before going to the trouble of testing polymer melts, it is important that the technique is validated by each user of the system. This is best and most easily done by measuring the surface tension of water or glycerin at room temperature and comparing the results to known literature values.