Curing, defects and mechanical performance of fiber-reinforced composites

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(1)UNIVERSIDAD POLITÉCNICA DE MADRID ESCUELA TÉCNICA SUPERIOR DE INGENIEROS DE CAMINOS, CANALES Y PUERTOS. Curing, Defects and Mechanical Performance of Fiber-Reinforced Composites. TESIS DOCTORAL. SILVIA HERNÁNDEZ RUEDA Ingeniera de Materiales Licenciada en Fı́sica. 2013.

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(3) Departamento de Ciencia de Materiales Escuela Técnica Superior de Ingenieros de Caminos, Canales y Puertos Universidad Politécnica de Madrid. Curing, Defects and Mechanical Performance of Fiber-Reinforced Composites. TESIS DOCTORAL. Silvia Hernández Rueda Ingeniera de Materiales Licenciada en Fı́sica. Directores de Tesis Carlos Daniel González Martı́nez Dr. Ingeniero de Caminos, Canales y Puertos Javier Llorca Dr. Ingeniero de Caminos, Canales y Puertos 2013.

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(5) A mi familia.

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(7) Agradecimientos En primer lugar, deseo expresar mi sincero agradecimiento a mis directores de tesis, Carlos Gonzalez y Javier Llorca, por su dedicación y ayuda durante la realización de este trabajo, por compartir su experiencia y por todo el conocimiento transmitido. Este agradecimiento se hace extensivo a mis compañeros del Instituto IMDEA Materiales por su ayuda y animos en numerosas ocasiones y por los buenos momentos compartidos durante estos años. En especial, a Vane, Katia y Natha por su apoyo, amistad y sobre todo por su cariño. También quiero expresar mi agradecimiento a Jon Molina y Federico Sket del Instituto IMDEA Materiales por su ayuda durante el proyecto DEFCOM (6o Programa Marco) y su dedicación, tiempo y ayuda con el tomógrafo. Agradecimiento que hago extensivo a la Technical University of Vienna (Austria) y FHOÖ Forschungs and Entwicklungs (Austria) por su colaboración, disposición y ayuda en el ámbito del proyecto DEFCOM. En especial a Marta Rodriguez Hortalá y Dietmar Salaberger por su interés, ayuda y ganas para sacar adelante el proyecto. Agradezco al Departamento de Ciencia de Materiales la colaraboración y facilidades recibidas durante la realización de la tesis. Me gustarı́a agradecer especialmente a mi familia su apoyo, paciencia y compresión durante esta etapa..

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(9) Resumen Tradicionalmente, la fabricación de materiales compuestos de altas prestaciones se lleva a cabo en autoclave mediante la consolidación de preimpregnados a través de la aplicación simultánea de altas presiones y temperatura. Las elevadas presiones empleadas en autoclave reducen la porosidad de los componentes garantizando unas buenas propiedades mecánicas. Sin embargo, este sistema de fabricación conlleva tiempos de producción largos y grandes inversiones en equipamiento lo que restringe su aplicación a otros sectores alejados del sector aeronáutico. Este hecho ha generado una creciente demanda de sistemas de fabricación alternativos al autoclave. Aunque estos sistemas son capaces de reducir los tiempos de producción y el gasto energético, por lo general, dan lugar a materiales con menores prestaciones mecánicas debido a que se reduce la compactación del material al aplicar presiones mas bajas y, por tanto, la fracción volúmétrica de fibras, y disminuye el control de la porosidad durante el proceso. Los modelos numéricos existentes permiten conocer los fundamentos de los mecanismos de crecimiento de poros durante la fabricación de materiales compuestos de matriz polimérica mediante autoclave. Dichos modelos analizan el comportamiento de pequeños poros esféricos embebidos en una resina viscosa. Su validez no ha sido probada, sin embargo, para la morfolologı́a tı́pica observada en materiales compuestos fabricados fuera de autoclave, consistente en poros cilı́ndricos y alargados embebidos en resina y rodeados de fibras continuas. Por otro lado, aunque existe una clara evidencia experimental del efecto pernicioso de la porosidad en las prestaciones mecánicas de los materiales compuestos, no existe información detallada sobre la influencia de las condiciones de procesado en la forma, fracción volumétrica y distribución espacial de los poros en los materiales compuestos. Las técnicas de análisis convencionales para la caracterización microestructural de los materiales compuestos proporcionan información en dos dimensiones (2D) (microscopı́a óptica y electrónica, radiografı́a de rayos X, ultrasonidos, emisión acústica) y sólo algunas son adecuadas para el análisis de la porosidad. En esta tésis, se ha analizado el efecto de ciclo de curado en el desarrollo de los poros durante la consolidación de preimpregnados Hexply AS4/8552 a bajas presiones mediante moldeo por compresión, en paneles unidireccionales y multiaxiales utilizando tres ciclos de.

(10) curado diferentes. Dichos ciclos fueron cuidadosamente diseñados de acuerdo a la caracterización térmica y reológica de los preimpregnados. La fracción vloúmetrica de poros, su forma y distribución espacial se analizaron en detalle mediante tomografı́a de rayos X. Esta técnica no destructiva ha demostrado su capacidad para nalizar la microestructura de materiales compuestos. Se observó, que la porosidad depende en gran medida de la evolución de la viscosidad dinámica a lo largo del ciclo y que la mayorı́a de la porosidad inicial procedı́a del aire atrapado durante el apilamiento de las láminas de preimpregnado. En el caso de los laminados multiaxiales, la porosidad también se vio afectada por la secuencia de apilamiento. En general, los poros tenı́an forma cilı́ndrica y se estaban orientados en la dirección de las fibras. Además, la proyección de la población de poros a lo largo de la dirección de la fibra reveló la existencia de una estructura celular de un diámetro aproximado de 1 mm. Las paredes de las celdas correspondı́an con regiones con mayor densidad de fibra mientras que los poros se concentraban en el interior de las celdas. Esta distribución de la porosidad es el resultado de una consolidación no homogénea. Toda esta información es crı́tica a la hora de optimizar las condiciones de procesado y proporcionar datos de partida para desarrollar herramientas de simulación de los procesos de fabricación de materiales compuestos fuera de autoclave. Adicionalmente, se determinarón ciertas propiedades mecánicas dependientes de la matriz termoestable con objeto de establecer la relación entre condiciones de procesado y las prestaciones mecánicas. En el caso de los laminados unidireccionales, la resistencia interlaminar depende de la porosidad para fracciones volumétricas de poros superiores 1%. Las mismas tendencias se observaron en el caso de GIIc mientras GIc no se vio afectada por la porosidad. En el caso de los laminados multiaxiales se evaluó la influencia de la porosidad en la resistencia a compresión, la resistencia a impacto a baja velocidad y la resistencia a copresión después de impacto. La resistencia a compresión se redujo con el contenido en poros, pero éste no influyó significativamente en la resistencia a compresión despues de impacto ya que quedó enmascarada por otros factores como la secuencia de apilamiento o la magnitud del daño generado tras el impacto. Finalmente, el efecto de las condiciones de fabricación en el proceso de compactación mediante moldeo por compresión en laminados unidireccionales fue simulado mediante el método de los elementos finitos en una primera aproximación para simular la fabricación de materiales compuestos fuera de autoclave. Los parámetros del modelo se obtuvieron.

(11) mediante experimentos térmicos y reológicos del preimpregnado Hexply AS4/8552. Los resultados obtenidos en la predicción de la reducción de espesor durante el proceso de consolidación concordaron razonablemente con los resultados experimentales..

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(13) Abstract Manufacturing of high performance polymer-matrix composites is normally carried out by means of autoclave using prepreg tapes stacked and consolidated under the simultaneous application of pressure and temperature. High autoclave pressures reduce the porosity in the laminate and ensure excellent mechanical properties. However, this manufacturing route is expensive in terms of capital investment and processing time, hindering its application in many industrial sectors. This fact has driven the demand of alternative outof-autoclave processing routes. These techniques claim to produce composite parts faster and at lower cost but the mechanical performance is also reduced due to the lower fiber content and to the higher porosity. Corrient numerical models are able to simulate the mechanisms of void growth in polymer-matrix composites processed in autoclave. However these models are restricted to small spherical voids surrounded by a viscous resin. Their validity is not proved for long cylindrical voids in a viscous matrix surrounded by aligned fibers, the standard morphology observed in out-of-autoclave composites. In addition, there is an experimental evidence of the detrimental effect of voids on the mechanical performance of composites but, there is detailed information regarding the influence of curing conditions on the actual volume fraction, shape and spatial distribution of voids within the laminate. The standard techniques of microstructural characterization of composites (optical or electron microscopy, X-ray radiography, ultrasonics) provide information in two dimensions and are not always suitable to determine the porosity or void population. Moreover, they can not provide 3D information. The effect of curing cycle on the development of voids during consolidation of AS4/8552 prepregs at low pressure by compression molding was studied in unidirectional and multiaxial panels. They were manufactured using three different curing cycles carefully designed following the rheological and thermal analysis of the raw prepregs. The void volume fraction, shape and spatial distribution were analyzed in detail by means of X-ray computed microtomography, which has demonstrated its potential for analyzing the microstructural features of composites. It was demonstrated that the final void volume fraction depended on the evolution of the dynamic viscosity throughout the cycle. Most of the initial voids.

(14) were the result of air entrapment and wrinkles created during lay-up. Differences in the final void volume fraction depended on the processing conditions for unidirectional and multiaxial panels. Voids were rod-like shaped and were oriented parallel to the fibers and concentrated in channels along the fiber orientation. X-ray computer tomography analysis of voids along the fiber direction showed a cellular structure with an approximate cell diameter of ≈ 1 mm. The cell walls were fiber-rich regions and porosity was localized at the center of the cells. This porosity distribution within the laminate was the result of inhomogeneous consolidation. This information is critical to optimize processing parameters and to provide inputs for virtual testing and virtual processing tools. In addition, the matrix-controlled mechanical properties of the panels were measured in order to establish the relationship between processing conditions and mechanical performance. The interlaminar shear strength (ILSS) and the interlaminar toughness (GIc and GIIc ) were selected to evaluate the effect of porosity on the mechanical performance of unidirectional panels. The ILSS was strongly affected by the porosity when the void contents was higher than 1%. The same trends were observed in the case of GIIc while GIc was insensitive to the void volume fraction. Additionally, the mechanical performance of multiaxial panels in compression, low velocity impact and compression after impact (CAI) was measured to address the effect of processing conditions. The compressive strength decreased with porosity and ply-clustering. However, the porosity did not influence the impact resistance and the coompression after impact strength because the effect of porosity was masked by other factors as the damage due to impact or the laminate lay-up. Finally, the effect of the processing conditions on the compaction behavior of unidirectional AS4/8552 panels manufactured by compression moulding was simulated using the finite element method, as a first approximation to more complex and accurate models for out-of autoclave curing and consolidation of composite laminates. The model parameters were obtained from rheological and thermo-mechanical experiments carried out in raw prepreg samples. The predictions of the thickness change during consolidation were in reasonable agreement with the experimental results..

(15) Contents. List of Figures. V. List of Tables. XIII. 1 Introduction. 1. 1.1. Fiber-reinforced Polymer Composites . . . . . . . . . . . . . . . . . . . . .. 1. 1.2. Manufacturing Defects in Composite Laminates . . . . . . . . . . . . . . .. 2. 1.3. Effect of Defects on Mechanical Performance . . . . . . . . . . . . . . . . .. 4. 1.4. Objectives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 9. 2 Consolidation and Curing of Thermoset Fiber-Reinforced Composites. 11. 2.1. Experimental Evidences . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 11. 2.2. Governing Equations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 14. 2.2.1. Resin Cure Kinetics . . . . . . . . . . . . . . . . . . . . . . . . . . .. 15. 2.2.2. Resin Viscosity . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 18. 2.2.3. Fiber Bed Permeability and Elasticity . . . . . . . . . . . . . . . .. 20. Flow-compaction modeling . . . . . . . . . . . . . . . . . . . . . . . . . . .. 26. 2.3. 3 Materials and Cure Cycle Definition. 35. 3.1. AS4/8552 prepreg system . . . . . . . . . . . . . . . . . . . . . . . . . . .. 35. 3.2. Cure Cycles Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 36. I.

(16) Contents. 3.3. 3.2.1. Rheology. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 36. 3.2.2. Isothermal Viscosity Profiles . . . . . . . . . . . . . . . . . . . . . .. 37. 3.2.3. Dynamic Viscosity Profiles . . . . . . . . . . . . . . . . . . . . . . .. 42. 3.2.4. Definition of Cure Cycles . . . . . . . . . . . . . . . . . . . . . . . .. 44. 3.2.5. Thermal Characterization . . . . . . . . . . . . . . . . . . . . . . .. 48. Manufacturing of Composite Laminates . . . . . . . . . . . . . . . . . . . .. 53. 3.3.1. 54. Thermogravimetric Measurements . . . . . . . . . . . . . . . . . . .. 4 Simulation of the Compaction Process 4.1. 57. Bidimensional Finite Element Model . . . . . . . . . . . . . . . . . . . . .. 57. 4.1.1. Fiber bed constitutive equation . . . . . . . . . . . . . . . . . . . .. 60. 4.1.2. Fiber Bed Permeability . . . . . . . . . . . . . . . . . . . . . . . . .. 63. 4.1.3. Effect of the Temperature Cycle on the Compaction . . . . . . . . .. 69. 5 X-ray Computed Tomography Characterization of Defects. 77. 5.1. Non-Destructive Evaluation Techniques . . . . . . . . . . . . . . . . . . . .. 77. 5.2. X-ray Computed Tomography Fundamentals . . . . . . . . . . . . . . . . .. 79. 5.3. Characterization of Void Population . . . . . . . . . . . . . . . . . . . . . .. 82. 5.3.1. Unidirectional Laminates . . . . . . . . . . . . . . . . . . . . . . . .. 83. 5.3.2. Multiaxial Laminates . . . . . . . . . . . . . . . . . . . . . . . . . .. 93. 6 Mechanical Behavior 6.1. 6.2. 109. Unidirectional Laminates . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110 6.1.1. Interlaminar Shear Strength (ILSS) . . . . . . . . . . . . . . . . . . 110. 6.1.2. Mode I and II Interlaminar Toughness . . . . . . . . . . . . . . . . 118. Multiaxial Laminates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 125 6.2.1. Plain Compression . . . . . . . . . . . . . . . . . . . . . . . . . . . 125 II.

(17) Contents 6.2.2. Low Velocity Impact . . . . . . . . . . . . . . . . . . . . . . . . . . 130. 6.2.3. Compression After Impact (CAI) . . . . . . . . . . . . . . . . . . . 139. 7 Conclusions and Future Work. 143. 7.1. Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143. 7.2. Future Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145. Appendices. 146. A Mathematica Input for Unidimensional Compaction. 149. B Abaqus Input for Unidimensional Compaction. 155. Bibliography. 161. III.

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(19) List of Figures. 1.1. Interlaminar shear strength as a function of void content for carbon fabric/epoxy laminates Costa et al. (2001). . . . . . . . . . . . . . . . . . . . .. 1.2. Interlaminar shear strength as a function of void content for carbon fabric/bismaleimide laminates Costa et al. (2001). . . . . . . . . . . . . . . . .. 1.3. 5. 6. Influence of the void content on the (a) longitudinal and (b) transverse tensile strength for [0]16 unidirectional carbon/epoxy composites T2H 132 300 EH (A) (Hexcel) and R922 12K (Ciba) (B) Olivier et al. (1995). . . . .. 1.4. Effect of vacuum pressure on void volume fraction and fatigue life at σmax = 0.8 Chambers et al. (2006). . . . . . . . . . . . . . . . . . . . . . . . . . . .. 2.1. 13. Thickness of individual plies of AS4/3501-6 laminates after autoclave curing. nc stands for the final number of compacted plies Cambell et al. (1985). . .. 2.4. 12. Compaction processes. a) resin flow normal to the laminate. b) Resin flow parallel to the plies. c) Mixed resin flow normal and parallel to the plies. .. 2.3. 9. Possible resin flow patterns, Dusi et al. (1987): a) Normal to the laminate, b) Parallel to the plies, c) Mixed flow. . . . . . . . . . . . . . . . . . . . . .. 2.2. 7. 14. Representative curing time-temperature-transformation diagram of a thermoset polymer, Berglund & J.M. Kenny (1991). . . . . . . . . . . . . . . .. 16. 2.5. Evolution of viscosity as a function of α and temperature . . . . . . . . . .. 19. 2.6. Pressure carried by the fibers as a function of the fiber volume fraction, Gutowski et al. (1986). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . V. 23.

(20) LIST OF FIGURES 2.7. 2. Normalized effective stress σ 0 /[A0 /( π163 βE )] (Equation 2.19) vs. fiber volume fraction for different maximum fiber volume fraction, Va , Gutowski et al. (1986). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 2.8. Load-displacement curve for the load-hold test method for AS4/3501-6 composite prepreg, Hubert & Poursartip (2001). . . . . . . . . . . . . . . . . .. 2.9. 25. 26. Schematic showing the geometry and the deforming coordinate system of a control volume. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 28. 2.10 Boundary conditions and material properties inputs for run R1 . . . . . . .. 30. 2.11 Simulation results for (a) resin pressure evolution and (b) fiber effective stress as a function of consolidation time. . . . . . . . . . . . . . . . . . . .. 32. 2.12 Evolution of compaction displacement as a function of time. . . . . . . . .. 33. 3.1. Gel point of the AS4/8552 prepreg under isothermal conditions at (a) 110◦ C, (b) 120◦ C, (c) 140◦ C, (d) 160◦ C, (e) 170◦ C and (f) 180◦ C. . . . . . . . . .. 40. 3.2. Storage (G0 ) and loss moduli (G00 ) of AS4/8552 prepreg at 120◦ C. . . . . .. 41. 3.3. ∗ Minimum complex viscosity, ηmin , and gel time, tgel , under isothermal con-. ditions for the AS4/8552 prepreg. . . . . . . . . . . . . . . . . . . . . . . .. 41. 3.4. Isothermal viscosity profiles of AS4/8552 prepreg. . . . . . . . . . . . . . .. 42. 3.5. Dynamic complex viscosity profiles of the AS4/8552 prepregs at different heating rates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 3.6. Viscosity measurements of 8552 epoxy resin and of S2/8552 prepregs Boswell (2000). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 3.7. 43. 44. Temperature profile of the cure cycles used to process AS4/8552 composite prepregs and the corresponding evolution of the complex viscosity, η ∗ , during the (a) cycle C-1, (b) cycle C-2 and (c) cycle C-3. . . . . . . . . . . . . . .. 3.8. 3.9. 46. Gel point of the AS4/8552 prepreg subjected to different cure cycles (a) cycle C-1, (b) cycle C-2 and (c) cycle C-3. . . . . . . . . . . . . . . . . . .. 48. MDSC Q200 (TA Instruments). . . . . . . . . . . . . . . . . . . . . . . . .. 49. VI.

(21) LIST OF FIGURES 3.10 Heat flow of the AS4/8552 prepreg as a function of temperature and heating rate (5, 8 and 10◦ C/min). . . . . . . . . . . . . . . . . . . . . . . . . . . .. 50. 3.11 Residual reaction heat of AS4/8552 prepreg after curing cycles C-1, C-2, C-3. 50 3.12 Glass transition temperature of AS4/8552 prepreg after curing cycles C-1, C-2, C-3 at onset. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 51. 3.13 Evolution of the degree of cure of the AS4/8552 prepreg. Predictions from Williams and Hubert model for curing cycles C-1, C-2, C-3 and experimental results of curing cycle C-1. . . . . . . . . . . . . . . . . . . . . . . . . . . .. 53. 3.14 Staking and packing process of unidirectional laminates. . . . . . . . . . .. 54. 3.15 Mass loss of AS4/8552 unidirectional laminates a function of temperature.. 55. 3.16 Mass loss of AS4/8552 multiaxial clustered laminates as a function of temperature (a) dispersed laminate [45o /0o /-45o /90o ]3s and (b) clustered lamio. o. o. o. nate [453 /03 /-453 /903 ]s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. 56. Resin bleeding during compression molding of a unidirectional panel. Fibers run horizontally and resin bleeding only occurred along the borders perpendicular to the fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 58. 4.2. Sketch and representative section of the panel for the finite element model.. 59. 4.3. a) Testing rig used for the compaction tests, b) Evolution of the laminate temperature. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 62. 4.4. Estimated compaction curve for the AS4/8552 prepreg. . . . . . . . . . . .. 63. 4.5. Linear fit according to Equation 4.6 of the logarithmic viscosity vs. 1/T at 8◦ C/min and 10◦ C/min. . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 4.6. Non-linear fit according to Equation 2.12 of the viscosity vs. degree of cure α at 130 and 160◦ C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 4.7. 4.8. 66. 66. Evolution of the complex viscosity with cure time: a) 120◦ C, b) 140◦ C, c) 160◦ C and d) 180◦ C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 68. Viscosity profiles for curing cycle (a) C-1, (b) C-2 and (c) C-3. . . . . . . .. 71. VII.

(22) LIST OF FIGURES 4.9. Numerical simulation of compaction strain as a function of the curing time for curing cycles C-1, C-2 and C-3 . . . . . . . . . . . . . . . . . . . . . . .. 71. 4.10 Numerical predictions of the evolution of the hydraulic conductivity as a function of the curing time for curing cycles (a) C-1, (b) C-2 and (c) C-3, element 501 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 74. 4.11 Evolution of the (a) pore pressure (Pr ) and (b) effective stress (σ 0 ) along width of the laminate for curing cycle C-2 . . . . . . . . . . . . . . . . . .. 75. 5.1. Schematic of a X-ray tomography system. . . . . . . . . . . . . . . . . . .. 80. 5.2. Principle of tomography and illustration of the Fourier slice theorem. . . .. 81. 5.3. Nanotom 160NF tomograph. . . . . . . . . . . . . . . . . . . . . . . . . . .. 82. 5.4. X-ray microtomography cross-section of the raw prepreg perpendicular to the fiber tows. Matrix appears as light gray regions, fibers tows as dark gray regions and pores are black. . . . . . . . . . . . . . . . . . . . . . . .. 5.5. 83. (a) OM montage of a cross-section of the composite panel manufactured with cure cycle C-1. (b) XCT slice of the same cross-section with 4 µm/voxel resolution. (c) Idem as (b) with 11 µm/voxel resolution. (d) Average of all the slices along the fiber direction with 4 µm/voxel resolution. (e) Idem as (d) with 11 µm/voxel resolution. Regions with a large volume fraction of interply voids are marked with an ellipse. . . . . . . . . . . . . . . . . . . .. 5.6. 84. (a) X-ray microtomography of void spatial distribution in the uniaxial composite panels manufactured according to the curing cycles C-1, C-2 and C-3. (b) Typical rod-like void together with its equivalent cylinder. . . . . . . .. 86. 5.7. Definition of the elongation factor and flatness ratio of individual voids. . .. 87. 5.8. Elongation factor of individual voids for the different cure cycles. . . . . . .. 88. 5.9. Dynamic evolution of the complex viscosity, η ∗ , of unidirectional AS4/8552 composite prepreg at the processing window region. . . . . . . . . . . . . .. 89. 5.10 Distribution of porosity across the width (Y axis) of the AS4/8552 unidirectional laminates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . VIII. 91.

(23) LIST OF FIGURES 5.11 Averaging gray values of X-ray absorption of the composite panel along the fiber axis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 91. 5.12 Average X-ray absorption of composite panel along the fiber (Z axis). Black zones stand for low density sections (pores), while white zones represent high density sections (fibers). Gray zones stand for matrix-rich regions. . .. 92. 5.13 Void distribution through the thickness of the laminate (X axis). . . . . . .. 93. 5.14 X-ray microtomography of void spatial distribution in the quasi-isotropic [453 /03 /-453 /903 ]s composite panel manufactured following the curing cycle C-3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 94. 5.15 Void density (expressed as the number of voids per mm3 ) as a function of the orientation of the major axis of the equivalent ellipsoid for the (a) [45/0/-45/90]3s dispersed quasi-isotropic laminates and (b) [453 /03 /-453 /903 ]s clustered quasi-isotropic laminates processed with different cure cycles. . .. 96. 5.16 Distribution of porosity along the width (Y axis) for AS4/8552 multiaxial panels (a) dispersed ([45/0/-45/90]3s ), (b) clustered ([453 /03 /-453 /903 ]s ). .. 97. 5.17 (a) Distribution of porosity along the width (Y axis) in a single cluster of three plies with fibers parallel to Z direction in the [453 /03 /-453 /903 ]s laminate manufactured according curing cycle C-3. (b) Average X-ray absorption of composite panel along the fiber (Z axis) of a single cluster of plies. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 99. 5.18 Void distribution through the thickness of the multiaxial panels (a) dispersed ([45/0/-45/90]3s ), (b) clustered ([453 /03 /-453 /903 ]s ). . . . . . . . . . . . . . 100 5.19 Dimensions of (a) major axis, (b) medium axis and (c) minor axis of individual voids for dispersed panels [45/0/-45/90]3s manufactured with curing cycles C-1, C-2 and C-3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102 5.20 Dimensions of (a) major axis, (b) medium axis and (c) minor axis of individual voids for clustered panels [453 /03 /-453 /903 ]s manufactured with curing cycles C-1, C-2 and C-3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104 5.21 Flatness ratio as a function of the void volume for different laminate ply clustering stacking sequences cured using cycle C-3. . . . . . . . . . . . . . 105 IX.

(24) LIST OF FIGURES 5.22 Elongation factor as a function of the void volume for different laminate ply clustering stacking sequences cured using cycle C-3. . . . . . . . . . . . . . 105 5.23 (a) Major axis, (b) medium axis and (c) minor axis dimensions of individual voids for panels manufactured with curing cycle C-3 and different laminate lay-ups: multiaxial dispersed ([45/0/-45/90]3s ), multiaxial clustered ([453 /03 /-453 /903 ]s ) and unidirectional ([0]10 ). . . . . . . . . . . . . . . . . 107 6.1. Three point bending fixture. . . . . . . . . . . . . . . . . . . . . . . . . . . 110. 6.2. ILSS load-displacement curves. . . . . . . . . . . . . . . . . . . . . . . . . 111. 6.3. Interlaminar shear strength of the unidirectional AS4/8552 composite laminates as a function of void content. . . . . . . . . . . . . . . . . . . . . . . 112. 6.4. Scanning electron micrograph of the fracture surface of a coupon tested to measure the ILSS; showing serrated feet for the laimate cured using cycle C-3.113. 6.5. Scanning electron micrographs of the fracture surfaces of coupons tested to measure the ILSS. (a) Cure cycle C-2. (b) Cure cycle C-3. . . . . . . . . . 114. 6.6. Cusp formation mechanism Greenhald (2009) . . . . . . . . . . . . . . . . 115. 6.7. (a) Load-indentation depth curves corresponding to pyramidal indentation tests of the resin processed with cure cycles C-2 and C-3, displaying identical behavior. (b) Array of indentations in one of the resin pockets is shown in the 30 × 30 µm SPM image. . . . . . . . . . . . . . . . . . . . . . . . . . . 116. 6.8. (a) Load-fiber displacement curves corresponding to fiber push-in tests in laminates processed with cure cycles C-2 and C-3. The arrow indicates the critical load for interfacial debonding, which was the same in both cases. (b) SPM image showing one fiber debonded from the matrix after the push-in test. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117. 6.9. X-ray tomograms of the cross-section of coupons tested to measure the ILSS for cure cycles C-1, C-2, C-3. . . . . . . . . . . . . . . . . . . . . . . . . . . 118. 6.10 (a) Sketch of the DCB specimens to measure GIc . (b) Typical load-cross head displacement curve. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119 X.

(25) LIST OF FIGURES 6.11 Load-cross head displacement curves for GIc for cure cycle (a) C-1, (b) C-2 and (c) C-3 of the unidirectional AS4/8552 laminates. . . . . . . . . . . . . 121 6.12 Mode I interlaminar fracture toughness, GIc , of the unidirectional [0]10 AS4/8552 laminates as a function of void content. . . . . . . . . . . . . . . 121 6.13 (a) Sketch of the specimens to measure GIIc . (b) Typical load-cross head displacement curve. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122 6.14 Load-cross head displacement curves of the GIIc test of AS4/8552 laminates (a) cycle C-1, (b) cycle C-2 and (c) cycle C-3. . . . . . . . . . . . . . . . . 124 6.15 Interlaminar fracture toughness GIIc of the unidirectional AS4/8552 laminates as a function of void content. . . . . . . . . . . . . . . . . . . . . . . 124 6.16 Scanning electron micrograph of the fracture surface of coupons tested to measure GIIc of unidirectional panels cured following cycle C-3. . . . . . . 125 6.17 Compression IITRI fixture. . . . . . . . . . . . . . . . . . . . . . . . . . . . 126 6.18 Compressive strength of the multiaxial AS4/8552 laminates as a function of void content.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 127. 6.19 (Fracture mechanisms in compression of multiaxial laminates manufactured using curing cycle C-2. a) dispersed stacking sequence [45/0/-45/90]3s (b) clustered stacking sequence [453 /03 /-453 /903 ]s . . . . . . . . . . . . . . . . . 130 6.20 Drop weight apparatus. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131 6.21 Load vs. time curves of multiaxial AS4/8552 laminates subjected to low velocity impact (a) cure cycle C-1, (b) cure cycle C-2 and (c) cure cycle C-3. 133 6.22 Load vs. time curves of multiaxial [45/0/-45/90]3s AS4/8552 laminates subjected to low velocity impact for curing cycles C-1, C-2 and C-3. . . . . . . 133 6.23 Load vs. time curves of multiaxial [453 /03 /-453 /903 ]s AS4/8552 laminates subjected to low velocity impact for curing cycles C-1, C-2 and C-3. . . . . 134 6.24 Results of the C-scan inspections of multiaxial AS4/8552 laminates subjected to low-velocity impact: (a) [453 /03 /-453 /903 ]s , (b) [45/0/-45/90]3s . . 135 XI.

(26) LIST OF FIGURES 6.25 Damage mechanisms of multiaxial laminates subjected to low velocity impact. 3D view of the impacted area (a) stacking sequence [453 /03 /-453 /903 ]s and (b) stacking sequence [45/0/-45/90]3s . . . . . . . . . . . . . . . . . . . 137 6.26 Damage mechanisms of multiaxial laminates subjected to low velocity impact. Cross-section under the impact (a) stacking sequence [453 /03 /-453 /903 ]s and (b) stacking sequence [45/0/-45/90]3s . . . . . . . . . . . . . . . . . . . 138 6.27 Conical distribution of delaminations within the laminate after low-velocity impact (a) stacking sequence [453 /03 /-453 /903 ]s and (b) stacking sequence [45/0/-45/90]3s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 6.28 Compression after impact fixture. . . . . . . . . . . . . . . . . . . . . . . . 140 6.29 Compressive strength after impact of multiaxial AS4/8552 laminates with different stacking sequence. . . . . . . . . . . . . . . . . . . . . . . . . . . . 141. XII.

(27) List of Tables. 1.1. Interlaminar delamination toughness for void free and voided laminates with a volume fraction of voids of 5% Asp & Brandt (1997). . . . . . . . . . . .. 3.1. Gel time of the AS4/8552 prepregs after consolidation following cure cycles C-1, C-2 and C-3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 3.2. 8. 45. Residual heat of reaction, ∆Hres , degree of cure, α, and onset glass transition temperature, Tg , of unidirectional AS4/8552 composite panels manufactured with different curing cycles. . . . . . . . . . . . . . . . . . . . . . . . . . .. 51. 3.3. Constants of the Williams and Hubert kinetic model. . . . . . . . . . . . .. 52. 3.4. Lay up of the manufactured panels. . . . . . . . . . . . . . . . . . . . . . .. 54. 4.1. Final compaction and bleeding strains. . . . . . . . . . . . . . . . . . . . .. 62. 4.2. Paramenters A and B for Kenny’s model. . . . . . . . . . . . . . . . . . . .. 65. 4.3. Compaction strains of unidirectional [0]10 laminates subjected o different curing cycles at 2 bars of pressure.. 4.4. . . . . . . . . . . . . . . . . . . . . . .. 69. Simulation and experimental results of the vertical strain at the end of the curing cycle . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 72. 4.5. Predicted and experimental mass loss for curing cycles C-1, C-2, C-3 . . .. 72. 5.1. Volume fraction of voids, Vf , void flatness ratio, f , and average distance between sections with high porosity along the panel width (Y axis), ∆d, as a function of the cure cycle for AS4/8552 unidirectional laminates. . . . . . XIII. 88.

(28) LIST OF TABLES 5.2. Volume fraction of voids, Vf , as a function of the cure cycle and ply-clustering for AS4/8552 composite panels manufactured with different curing cycles. .. 6.1. 95. Interlaminar shear strength of [0]10 laminates. The average values and standard deviation were obtained from 5 tests for each condition. . . . . . . . . 111. 6.2. Resin hardness, H, and critical load for fiber-matrix interfacial debonding, Pc , as determined from nanoindentation tests. . . . . . . . . . . . . . . . . 116. 6.3. Compressive modulus (Ec ) and compressive strength (σc ) of multiaxial laminates processed using different curing cycles. . . . . . . . . . . . . . . . . . 128. 6.4. Elastic and dissipated energies during low velocity impact of multiaxial AS4/8552 panels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 134. 6.5. Projected delamination areas of multiaxial AS4/8552 panels with different lay-up after low velocity impact. . . . . . . . . . . . . . . . . . . . . . . . . 136. 6.6. Compression after impact strength of multiaxial AS4/8552 laminates with different stacking sequence. . . . . . . . . . . . . . . . . . . . . . . . . . . . 142. XIV.

(29) Chapter. 1. Introduction. 1.1. Fiber-reinforced Polymer Composites. Fiber-reinforced polymer-matrix composites (FRPs) are nowadays extensively used in structural elements due to their high specific stiffness and strength. FRPs are constituted by high performance fibers (carbon, glass, aramid, etc.) embedded in a thermoset matrix and they are normally provided in the form of prepregs semifinishied products. Prepreg sheets consist of a fabric impregnated with resin maintained in a pre-gelled state. Laminates are formed by stacking (manually or automatically) individual prepreg sheets that are consolidated by the simultaneous application of pressure and temperature. The external applied pressure impedes the growth of voids during curing and even leads to the collapse of air bubbles, giving rise to materials with very low porosity and excellent mechanical properties, as required by aerospace and sports industries. Traditionally, the consolidation process is carried out in autoclave which is an expensive manufacturing route in terms of capital investment and processing time, limiting the expansion of composite materials to other industrial sectors. These limitations act as driving forces to look for alternative out-of-autoclave processing routes (OOA), including, among others, the use of prepregs in 1.

(30) 1.2 Manufacturing Defects in Composite Laminates a standard resin transfer molding process (SQRTM), replacing pressurized gas by a fluid with a high thermal inertia to reduce the curing time (QUICKSTEP) Griffiths & Noble (2004), or the development of special low temperature cure prepregs that can be cured in standard ovens. These techniques are able to produce composite parts faster but it should be noted that they do not often obtained the mechanical properties achieved by autoclave processing due to the lower fiber content and to the presence of voids.. 1.2. Manufacturing Defects in Composite Laminates. Defects are introduced in the composite laminates in all manufacturing processes, although the size and frequency of each type depends upon the processing cycle. Typical defects found in thermoset composite parts as a result of the manufacturing conditions are: • Porosity (voids) due to volatile resin components or to trapped air bubbles. • Foreign bodies (for example backing paper of the prepreg sheets). • Incorrect fibre volume fraction due to excess or insufficient resin. Slight local variations of volume fraction are always present but large differences from specifications may be caused by inappropriate processing conditions. • Bonding defects. Components may be bonded together (e.g. panels and stringers) during manufacturing and defects in the bondline occur due to incorrect cure conditions for the adhesive or contamination of the surfaces to be bonded. • Fibre misalignment and fiber waviness. Waviness is detrimental for the mechanical performance of the material (particularly in compression). • Ply misalignment caused b errors during lay-up of the laminate plies. Ply misalignment alters the overall stiffness and strength of the laminate and may cause warping during cure. • Incompletely cured matrix due to incorrect curing cycle or faulty material. • Ply cracking. Thermally-induced cracks occur with certain ply lay-ups due to the differences in the thermal expansion coefficient of the plies. 2.

(31) 1.2 Manufacturing Defects in Composite Laminates • Delaminations are planar defects usually at ply boundaries. They are not typical during manufacturing but may be produced by contamination during lay-up, machining or impact damage (e.g. tool drops). • Fiber defects are one of the ultimate limiting factors in determining strength of composites, and sometimes faulty fibers can be identified as the sites from which damage was initiated. The most important manufacturing defect in composite laminates is porosity. Many of the other defects occur more rarely and always lead to porosity formation. Resin flow governed by pressure gradients in the laminate play a critical role in void formation and migration. Understanding the flow-compaction mechanisms during manufacturing is the key to control the porosity of composite parts. The control of thermoset prepregs manufacturing require the understanding of the resin rheological properties during curing as well as of the cure kinetics. The final quality of the laminate will depend on the competion of resin flow and cure reactions. Low viscosity of the resin is required to impregnate adequately the fiber preform and this is favored by an adequate temperature cycle design. Increasing temperature accelerates the cure reactions and hinders the void migration mechanisms. Rheological analysis has been used to study the cure process of epoxy resins Berglund & J.M. Kenny (1991), Wang et al. (1997) and epoxy prepregs Simon & Gillham (1993) and is essential for the optimization of composite processing. The mechanics of prepreg compaction in autoclave was pioneered by Springer and coworkers Loos & Springer (1983), Tang et al. (1987) starting from the consolidation theory developed for soil mechanics Terzaghi (1943). These authors described the resin flow through the composite following Darcy’s flow theory in a porous medium, and determined the laminate compaction sequence. The external pressure was first supported by the resin and pressure was transferred to the fiber bed as bleeding progressed through the laminate surfaces. This process continued until the composite reached the maximum compaction of the reinforcement and all the resin excess was expelled. Air bubbles are always present in the raw prepreg due to deficient fiber impregnation during prepreg manufacturing. In addition, voids are also introduced during the preparation of the laminate kit. The stability of voids as a function of the temperature and 3.

(32) 1.3 Effect of Defects on Mechanical Performance pressure has been extensively studied by Kardos et al. Kardos et al. (1986), who considered the effects of the resin viscosity and of the resin-void surface tension. They developed a model for void growth which was successfully applied to predict the evolution of voids in thermoset composites Ledru et al. (2010), Grunenfelder & Nutt (2010). Although these models provide the essentials of the mechanics of void growth in polymers, they are restricted to small spherical voids embebed in a viscous resin. Their validity is not proved for long cylindrical voids, the standard morphology observed unidirectional fiber reinforced composites. In addition, although there are many references in the scientific literature regarding the detrimental effect of voids on the mechanical performance of composites Bowles & Frimpong (1992), Costa et al. (2001), Wisnom et al. (1996), there is still a lack of information regarding the influence of curing conditions on the actual volume fraction, shape and spatial distribution of voids within the laminate.. 1.3. Effect of Defects on Mechanical Performance. The effect of voids on the mechanical properties of composites has been the object of many investigations. The results show that fiber-dominated mechanical properties are not significantly influenced by voids Olivier et al. (1995), Bureau & Denault (2004), while matrix-dominated properties are strongly dependent on their presence. Reduction in interlaminar shear strength Olivier et al. (1995), Wisnom et al. (1996), Costa et al. (2001), compressive strength Suarez et al. (1996), Cinquin et al. (2007) tensile transverse strength Olivier et al. (1995), Varna et al. (1995), bending Olivier et al. (1995), Hagstrand et al. (2005), fatigue Bureau & Denault (2004), Almeida & Nogueira Neto (1994), Chambers et al. (2006) and fracture toughness Asp & Brandt (1997), Rizov (2006) have been reported in the literature. The effect of void content on interlaminar shear strength (ILSS) was investigated by Wisnom et al. (1996) using glass/epoxy and carbon/epoxy UD laminates and by Costa et al. (2001) using carbon/epoxy and carbon/bismaleimide woven laminates. Both studies reported a reduction between 8% and 33% depending on the void content ranging from 1.1 to 5.6%, Fig. 1.1 and 1.2. The reduction in the interlaminar shear strength with the void content was justified in both cases on the basis that crack nucleation starts from the voids, according to Scanning Electron Microscopy (SEM) observations on the broken 4.

(33) 1.3 Effect of Defects on Mechanical Performance samples. SEM observations showed that the void location was strongly dependent on the specific matrix system. In the case of epoxy resin, the voids were preferentially located at the crossing point of the woven fibre tows, while they were typically found at interface of woven fibre tows in the carbon/bismaleimide laminates.. 8 0 7 5. IL S S (M P a ). 7 0 6 5 6 0 5 5 5 0 1. 2. 3. 4. 5. 6. V o id s c o n te n t ( % ) Figure 1.1: Interlaminar shear strength as a function of void content for carbon fabric/epoxy laminates Costa et al. (2001).. Olivier et al. (1995) analyzed the effect of curing cycle pressure on the porosity of carbon/epoxy UD laminates and reported a similar reduction in the ILSS with void contents in the range 0.3 and 10%. The effect of voids on the longitudinal and transverse tensile properties were also investigated by these authors. They noticed that the longitudinal modulus as well as the longitudinal tensile strength (fiber-dominated properties) were not affected by the porosity, Fig. 1.3. However, the transverse properties (which are matrixcontrolled) were found to be extremely sensitive to the presence of defects with a reduction of 10% and 30%, respectively, for a void contents of 0.3 and 10% respectively, Fig. 1.3. The shape and size of the voids was characterized by means of optical microscopy and image analysis for different curing cycles. Void shape was assessed from three different sections of the same void obtained from at least three parallel cut planes spaced ≈ 10 µm apart. For a given void content, the specimens with the largest voids showed a reduction in the bending 5.

(34) 1.3 Effect of Defects on Mechanical Performance. 8 0. IL S S (M P a ). 7 5. 7 0. 6 5. 6 0. 5 5 1 .0. 1 .5. 2 .0. 2 .5. 3 .0. 3 .5. V o id s c o n te n t ( % ) Figure 1.2: Interlaminar shear strength as a function of void content for carbon fabric/bismaleimide laminates Costa et al. (2001).. modulus three times larger (15%) than those with small defects (5%). The influence of void content on the bending properties was also investigated by Hagstrand et al. (2005) for UD glass fibre reinforced polypropylene composites. As in case of Olivier et al. (1995), they found a reduction in both flexural modulus and bending strength of 20% and 28%, respectively, for a void content of 14%. Suarez et al. (1996) investigated the effect of void content on the compressive strength of UD carbon/epoxy laminates. They found a roughly linear correlation between void content and compressive strength, with a reduction of about 40% for a volume fraction of 4% of voids. Lower influence of porosity was found by Cinquin et al. (2007) for quasi isotropic carbon/epoxy laminates with a the reduction in the compressive strength of 14% for a void content of 11%. Asp & Brandt (1997) investigated the effects of pores and voids on the interlaminar delamination toughness of carbon/epoxy laminates, by means of static Mode I, Mode II and mixed mode fracture tests, Table 1.1. The results were inconclusive due to the large scatter. 6.

(35) 1.3 Effect of Defects on Mechanical Performance. L o n g itu d in a l te n s ile s tr e n g th ( % ). (a ) C o m p o s ite A C o m p o s ite B. 1 0 0. 9 5. 9 0. 8 5 0. 2. 4. 6. 8. 1 0. 1 2. V o id c o n te n t ( % ) (b ). T r a n s v e r s e te n s ile s tr e n g th ( % ). 1 0 0. C o m p o s ite A C o m p o s ite B. 9 0. 8 0. 7 0 0. 2. 4. 6. 8. 1 0. 1 2. V o id c o n te n t ( % ). Figure 1.3: Influence of the void content on the (a) longitudinal and (b) transverse tensile strength for [0]16 unidirectional carbon/epoxy composites T2H 132 300 EH (A) (Hexcel) and R922 12K (Ciba) (B) Olivier et al. (1995).. 7.

(36) 1.3 Effect of Defects on Mechanical Performance GC (J/m2 ) Test method Void free laminate. GC (J/m2 ) Voided laminate. Mode I. 229.0 ± 17.8. 239.8 ± 9.7. Mode II. 883.1 ± 117.5. 811.3 ± 57. Mixed Mode. 478.8 ± 43.7. 454.5 ± 70.3. Table 1.1: Interlaminar delamination toughness for void free and voided laminates with a volume fraction of voids of 5% Asp & Brandt (1997).. Olivier et al. (1995) found that Mode I fracture toughness depend very much on the void volume fraction. They reported a reduction of 22% in GIC for a void content of 5%.. Fatigue properties were in general more affected by the void content than static properties. Almeida & Nogueira Neto (1994) carried out four-point bending tests on [0/90]12 carbon/epoxy laminates and found that the static strength was not influenced by a void content of 3% but had a detrimental effect on fatigue strength. Cyclic bending results Bureau & Denault (2004) for continuous glass fibre/polypropylene woven composites showed that different void contents led to a shift of the S-N curves without changing their slope, indicating a reduction of fatigue life with increasing void content. The damage evolution under bending fatigue was also investigated by Chambers et al. (2006) for UD carbon fibre composites. They noticed that the fatigue life changed from 2000 to 106 cycles by varying the void content from 1.6% to 3.1%, Fig. 1.4. The authors concluded that the voids played a fundamental role in the fatigue life when they were located in the inter-ply region where delamination occurred.. Rizov (2006) investigated the influence of voids on the Mode I fatigue behavior of glass fiber reinforced polypropylene plates manufactured by injection molding. An increase in the void content resulted in higher crack propagation rates. A limited influence was, however, reported for void volume fractions below 1%, whereas higher void contents (up to 7%) induced significant reductions in the fatigue crack propagation threshold and fatigue crack growth resistance. 8.

(37) 1.4 Objectives. Figure 1.4: Effect of vacuum pressure on void volume fraction and fatigue life at σmax = 0.8 Chambers et al. (2006).. 1.4. Objectives. Traditionally, the manufacturing of high performance polymer-matrix composites is carried out by means of autoclave systems using prepreg tapes stacked and consolidated under the simultaneous application of pressure and temperature. High autoclave pressures prevent the growth of bubbles and promote the collapse of air entrapments in the laminate, controlling the final porosity and ensuring the high performance of the components. However this manufacturing method is expensive in terms of capital investment and processing time and hence is not cost-effective for other industries. This fact has driven an increasing demand of alternative out-of-autoclave processing routes (OOA). However, these techniques usually are not able to produce composite parts with equivalent mechanical properties in comparison with components manufactured using autoclave due to the lower fiber volume fraction and higher porosity contents. A deeper understanding of the effect of the processing conditions (pressure and temperature) on prepreg consolidation would allow to improve the quality of the components manufactured by means of out-ofautoclave processing routes. According to this, the main goal of this thesis was to assess the effect of the curing cycle on the development of voids during consolidation of prepreg at low pressure. The proper design of the temperature curing cycle, based on the rheological 9.

(38) 1.4 Objectives and thermal characterization of the prepregs, led to the manufacture of unidirectional and multiaxial panels with controlled void content. The void volume fraction, shape and spatial distribution were also analyzed in detail by means of X-ray computed microtomography and the results were discussed in the light of the processing conditions. This information is critical to optimize processing parameters and to provide inputs for virtual testing and virtual processing tools. In addition, the matrix-controlled mechanical properties of the panels were measured in order to establish the effect of the voids on the mechanical performance of the laminates. Finally, the effect of the processing conditions on the compaction behavior of unidirectional AS4/8552 panels manufactured by compression molding was simulated using the finite element, as a first approximation to more complex and accurate models for out-of-autoclave curing and consolidation of composite laminates.. 10.

(39) Chapter. 2. Consolidation and Curing of Thermoset Fiber-Reinforced Composites. 2.1. Experimental Evidences. Processing of thermoset composites takes place by the simultaneous exposition of the material to heat and pressure for a given period of time. The resulting cure cycle is therefore a combination of temperature and pressure profiles. The temperature leads to the initiation of the crosslinking chemical reactions. It also reduces the viscosity of the resin favoring the impregnation of the fibers while the excess of resin and vapor bubbles are squeezed out from the material. Pressure and temperature are the driving forces to bleed the laminate, consolidate individual plies and reduce the void content. During consolidation of prepreg materials, resin flow can be dominant in the direction perpendicular to the laminate (Fig. 2.1.a), parallel to the laminate (Fig. 2.1.b) or in both directions (Fig. 2.1.c). Depending on the width to thickness ratio, the first case is repre11.

(40) 2.1 Experimental Evidences. Figure 2.1: Possible resin flow patterns, Dusi et al. (1987): a) Normal to the laminate, b) Parallel to the plies, c) Mixed flow.. sentative of the compaction process under hydrostatic pressure which occurs in autoclave consolidation while the second type of flow is representative of the behavior under hot press conditions. The typical compaction mechanisms in porous media are percolation and shear flow. The resin flows through the pores between the fibers when percolation flow is dominant and the resin excess is squeezed-out allowing the compaction of the material which attains the maximum fiber volume fraction. The percolation mechanism is typically used to describe resin flow within the laminates in thermoset matrix composites. Under shear flow, fiber and matrix experience a homogeneous solid-like deformation and the material behavior under compaction is similar to that of a soft solid. Shear flow is usually observed in thermoplastic matrix composites in which the high viscosity of polymer prevents percolation flow. The compaction and the arrangement of individual plies during consolidation are controlled by the prevailing flow patterns associated to each manufacturing route (Fig. 2.2). Under through-thickness flow, compaction occurs sequentially and the thickness of each individual ply decreases gradually from the top of the tool surface to the bottom, (Fig. 2.2). The resin is squeezed out from the gap between the first and second ply and then it is again squeezed-out from the second gap due to the movement of the two first layers. This process is repeated up to the final compaction of the laminate. However, in case of resin flow parallel to the laminate, the thickness reduction of all plies is simultaneous. Compaction came out as a result of both mechanisms in the case of mixed flow (parallel and perpendicular to the plies). The mechanics of autoclave prepreg compaction was pioneered by Springer and coworkers, Springer (1982), starting from the consolidation theory developed for soil mechan12.

(41) 2.1 Experimental Evidences. Figure 2.2: Compaction processes. a) resin flow normal to the laminate. b) Resin flow parallel to the plies. c) Mixed resin flow normal and parallel to the plies.. ics. Their experiments verified the wavelike nature of the compaction process described above by analyzing the relative motion of a suspension of a rod network in a viscous liquid. Subsequently, Cambell et al. (1985) found an analogous mechanism of compaction in thick graphite-epoxy laminates. In this case, the thickness of each layer was measured using photomicrographs of laminates cured at different pressures in autoclave. The results are shown in Fig. 2.3. As expected, the compacted layers were located at the top of the laminate (bag surface) and the number of fully compacted plies increased with the applied 13.

(42) 2.2 Governing Equations pressure. Resin flow occurs only through regions with pressure gradients and ends when they are relieved during the process.. Figure 2.3: Thickness of individual plies of AS4/3501-6 laminates after autoclave curing. nc stands for the final number of compacted plies Cambell et al. (1985).. 2.2. Governing Equations. The optimum curing conditions for a given composite system can be determined once the fundamental physical and chemical mechanisms involved are well understood. Obviously, the optimum cure cycle can be established empirically by means of expensive trial and error experimental campaigns, but the whole approach can be more efficient by means of mathematical models representing the underlying physics of the compaction phenomena. A suitable model for simulating the curing process should be supported by a set of submodels based on the governing equations describing the physico-chemical phenomena occurring during processing (i.e. cure kinetics, resin flow, ply compaction, heat transfer, residual stresses, etc.). Such kind of approaches could considerably reduce the number of experimental trials to reach an optimum cure cycle. The following sections are devoted 14.

(43) 2.2 Governing Equations to summarize the governing equations controlling the compaction phenomena in standard thermoset prepreg manufacturing.. 2.2.1. Resin Cure Kinetics. Cure of thermoset resins occurs via the incorporation of curing agents that trigger the curing reactions -by addition or condensation chemical mechanisms- leading to a three dimensional cross-linked network of polymeric chains. During the process, the resin experiences a number of changes which depend on time and temperature: gelation, vitrification and cure. These phenomena are usually represented in time-temperature-transformation diagrams (Fig. 2.4). Regions in the diagram represent different physical states of a given thermoset polymer: liquid, gel-gummy, gel-glassy and vitreous-ungelled. The gel point is defined as the instant at which the three-dimensional network reaches an infinite molecular weight due to an irreversible process. The initial stages of the resin curing will be more likely dominated by purely viscous effects but the resin will behave increasingly as a viscoelastic solid as the crosslinking reactions progress, and particularly near to the gel point. Above the gel point, the polymer behaves as a solid and the resin no longer needs the mold or the die to maintain its final shape so the part can be demolded. After gelification, vitrification may occur when the curing process takes place under non-isothermal conditions if the glass transition temperature, Tg , rises the cure temperature leading to a drastic reduction of the cure rate due to the restriction of mobility between neighbor polymeric chains. The reduction in the cure rate at vitrification is believed to be caused by a shift in the rate-controlling mechanism from kinetics (dependent on temperature and reactants concentration) to diffusion as a result of the reduction in the resin free-volume and the molecular mobility that accompanies this transition, Montserrat (1992) and Berglund & J.M. Kenny (1991). Three critical temperatures are highlighted on the temperature axis: Tg0 , the glass transition temperature for completely uncured resin,. gel Tg ,. the temperature at which vit-. rification and gelation occurs simultaneously, and Tg∞ , is the glass transition temperature of the fully-cured material. In order to understand in more detail the cure reactive process, it is necessary to examine the reaction kinetics for a given temperature-time profile. The curing kinetics can be 15.

(44) 2.2 Governing Equations. Figure 2.4: Representative curing time-temperature-transformation diagram of a thermoset polymer, Berglund & J.M. Kenny (1991).. analyzed from a double perspective: the microscopic models based on mechanistic kinetic approaches and the macroscopic phenomenological models, Van Overbeke et al. (2001). The former analyze the kinetic mechanisms associated with each of the reactions involved in the process resulting in very complex models coupled with sophisticated experimental techniques for measuring the concentration of all chemical species. The latter phenomenological methods analyze the overall process from a single reaction which is selected to represents the global curing process. Such models are semi-empirical and do not provide a clear description of the individual chemical reactions involved in the process but they do not require very sophisticated experimental techniques for parameter identification and can provide very accurate results. The phenomenological models are developed from the concept of the reaction rate inferred from the heat generated during the crosslinking reaction. For instance, let us assume two reactive groups A and B present in a given resin system whose initial concentrations, CA0 and CB0 , are known. The reaction rate of components A and B depends on the curing temperature and concentration of the reactants (kinetic control): the higher the concen16.

(45) 2.2 Governing Equations tration of A and B, the higher the reaction probability between them. The reaction rate, vreac , is defined as the time derivative of the variation of concentration of the reagent, dα dt. vreac =. (2.1). where α is the degree of cure defined as,. α=. CA0 − CA CA0. (2.2). where CA is the concentration of the component A at time t and CA0 the initial concentration. The value of α ranges from 0 at the initial stage to 1 when the material is fully cured. The general kinetic equation expressing the variation of the cure rate with temperature T and concentration of the reactants is expressed mathematically by Van Overbeke et al. (2001). dα = κ(T )f (α) dt. (2.3). where f (α) is a function which depends on the current reactant concentration and κ(T ) is a thermally-activated rate constant defined by an Arrhenius-type equation as, κ(T ) = A exp. −Ea RT. (2.4). where A is a proportionality constant, Ea the activation energy, and R the ideal gas constant. Substituting Equation 2.4 into Equation 2.3 yields the time derivative of the degree of cure as, dα = A exp dt. . −Ea f (α) RT. (2.5). Several expressions for f (α) have been proposed in the past to fit experimental results, Keenan (1987), Mijovic et al. (1984), Moroni et al. (1986) and Dusi et al. (1987). Most epoxy-amine systems show an autocathalytic behavior during the cure and, the term f (α) can be expressed in such cases as Yang et al. (1999), 17.

(46) 2.2 Governing Equations. f (α) = αm (1 − α)n. (2.6). where m and n are the orders of cure reaction. Notice that f (0) = 0 and f (tf ullcure ) = 1. Substituting Equation 2.7 into Equation 2.3 and rearranging terms yields the typical expression of the autocatalytic model for the dynamic curing process without diffusion, also known as Borchardt and Daniels equation, Borchardt & Daniels (1956), dα = A exp dt. . −Ea αm (1 − α)n RT. (2.7). This expression is strictly valid up to the point where the reaction becomes kinetically controlled. While this is usually true for early stages of the cure process, other factors may come into play as reactants are consumed and a macromolecular polymer network is formed. Borchardt and Daniels approach was modified by Johnston and Hubert (1995), Hubert et al. (1995), to take into account reduction in the curing rate at the last stages of cure as a result of the change in mechanism from kinetics to diffusion when Tg reaches the cure temperature. Mathematically, the cure rate can be expressed as follows, dα A exp [−Ea /RT ]αm (1 − α)n = dt 1 + exp [C(α − (αC0 + αCT T ))]. (2.8). where m, n, A, C, αC0 and αCT stand for model constants to be determined experimentally. The diffusion mechanisms are included by adding the term [1/[1+exp [C(α − (αC0 + αCT T ))]]] to Equation 2.8.. 2.2.2. Resin Viscosity. For a Newtonian fluid, the applied shear stress necessary to deform a fluid, τ , is proportional to the shear velocity gradient, γ̇, according to. τ = η γ̇. (2.9). where η is the viscosity of the resin and γ̇ stands for the velocity gradient perpendicular to the fluid motion. It is given by ∆v/h, where ∆v is the velocity difference (relative velocity) and h is the distance between adjacent layers. 18.

(47) 2.2 Governing Equations Matrix resin flow during prepreg compaction is induced by the pressure gradient necessary to remove the excess of resin from the laminate, promote adequate bonding between plies, and collapse most of voids within the laminate. The rheological behavior of thermoset resins is governed by two main physical mechanisms. On the one hand, the viscosity decreases with temperature as a result of the higher mobility of the polymer chains. On the other hand, cross-linking reactions, which are thermally activated, lead to an increase of viscosity. The resin viscosity can be modeled also by empirical equations assuming that the temperature, T , and degree of cure, α(t), are known at any time during the curing process. Several approaches can be found in the literature assuming uncoupled effects of temperature and degree of cure (i.e. Lee et al. (1982), Dusi et al. (1987), Ciriscioli et al. (1992) and Kenny (1992)). Such approaches are based on a general constitutive equation in which both phenomena are described separately as, Flory (1953) (Fig. 2.5).. η(T, α) = Φ(T )χ(α). (2.10). where η is the resin viscosity and Φ(T ) and χ(α) functions of the temperature and the degree of cure, respectively. This expression leads to a minimum of viscosity over the time that can be used to define the processing window of the material.. Figure 2.5: Evolution of viscosity as a function of α and temperature. Stolin et al. (1979) and lately Lee et al. (1982), Dusi et al. (1987) and Ciriscioli et al. (1992) proposed uncoupled models based on the following equation, 19.

(48) 2.2 Governing Equations. η(T, α) = η0 exp. −U RT. + κα. (2.11). where U is the activation energy associated with viscous flow and κ a constant accounting for the effect of the chemical reaction on the resin viscosity. It was also assumed that U is independent of α and therefore it only leads to a constant shift in the viscosity under isothermal conditions. Alternative models were developed by Kenny (1992), Kim & Kim (1994) in an attempt to predict the rheological behavior of the resin more accurately for cure degrees close to the gelification point by incorporating the parameter αg , which is related with the degree of cure at gel point as, η(T, α) = Aη exp. −Eη RT. . αg (αg − α). (a+bα) (2.12). where Aη , Eη , a, b and αg are model parameters. Equation 2.12 is also an Arrhenius-type relation in which temperature and degree of cure are uncoupled.. 2.2.3. Fiber Bed Permeability and Elasticity. The consolidation of thermoset FRPs is a complex process involving coupled mechanisms such as the resin rheology and cure behavior. Other mechanisms controlling the compaction phenomena are related with the fiber bed architecture, namely permeability and elasticity. • Fiber bed permeability Permeability characterizes the permeation of a fluid through a porous medium. As discussed previously, the resin has to be squeezed-out during the composite consolidation to achieve the maximum fiber volume fraction, to remove voids and to favor fiber impregnation. The resin flow through the channels of the fiber preform can be easily described by Darcy’s law by means of the permeability parameter which establishes the relationship between the flow rate and the pressure gradient necessary to drive the flow. This law was originally developed for the flow of Newtonian fluids through porous media made up of granular particles. Gebart (1992) validated Darcy’s law for low flow rate processes as 20.

(49) 2.2 Governing Equations the one ocuring during composite compaction. The generalized form of Darcy’s law is expressed as,. ~u = −. K gradP η. (2.13). where u is the volume averaged flow velocity, η the viscosity of the fluid, gradP the pressure gradient, and K the permeability tensor of the fiber preform. The three-dimensional form of Darcy’s law can be expressed in matrix form as,       u x .     ∂P    Kxx Kxy Kxz    ∂x   1 ∂P uy = − Kyx Kyy Kyz ∂y      η     ∂P   u     K K K z. zx. zy. zz. (2.14). ∂z. where ux , uy , uz are the resin velocity components and Kij the permeability tensor components. Kij = 0 for i 6= j when the orthotropic axes of the fabric are used as the reference frame. The motion of the resin through the fiber preform is usually modeled as the flow through a porous medium constituted by the fiber network. As a result, the resin mobility decreases as the fiber volume fraction increases during compaction. Other factors, such as the fiber architecture and sizing of the fibers, can also affect the permeability of the fiber bed. Hence, the permeability factors of the fiber bed should be determined experimentally from the relation between the pressure drop and the flow rate through the fiber network or estimated by means of empirical models. Many empirical models based on the physics of lubrication flow or flow through capillarity tubes have been developed to describe this relationship. The Carman-Kozeny equation (Equation 2.15), developed by adopting the capillary models from the soils mechanics literature, is one of the most widely accepted for calculating the permeability of fiber beds. It considers the porous medium as a system of parallel capillaries with diameters estimated in terms of the hydraulic radius of the system, rf2 (1 − Vf )3 KK = 4kK Vf 21. (2.15).

(50) 2.2 Governing Equations where rf is the fiber radius, Vf is the fiber volume fraction, KK the permeability in the flow direction and kK the Kozeny constant which has to be determined experimentally. However, several shortcomings of this model should be indicated. The resin flow and therefore the value of Kozeny constant, kK , will depend on fluid type, fiber packing and porosity. Experiments carried out by Lam & Kardos (1989) indicated that the permeability was dependent on the permeating liquid due to deviations from Newtonian fluid behavior and on fiber arrangements due to deviations from the assumed regular patterns. Moreover, the experimental values of kK determined for high porosities cannot be used to describe low porosity scenarios. In conclusion, Equation 2.15 fails to predict the permeability over the total porosity range of the fiber beds, Skartsis et al. (1992), Gebart (1992) and Aström et al. (1992). Gutowski and coworkers, Gutowski et al. (1987b), found that Equation 2.15 could give a good fit to the axial permeability of unidirectional reinforcements, while there were certain discrepancies for the transverse permeability as the model was not able to capture the absence of flow when the fibers touch each other blocking any transverse flow. Despite all the above limitations, the Carman-Kozeny equation seems to be valid for slow Newtonian flow through porous media over moderated porosity ranges. For these situations, the proportionality between flow rate and pressure drop is retained and Darcy’s law is valid Carman (1956), Durst et al. (1987), Gebart (1992), Dullien (1979). Alternative constitutive equations were developed to overcome these shortcomings Gebart (1992), Bruschke & Advani (1993) and Gutowski et al. (1987b) but there is no universal able to predict accurately the permeability of fiber beds within the whole porosity range. The particular feature of each composite manufacturing process should be taken into account when dealing with physically-based simulations. • Fiber bed elasticity For technologically relevant materials, the fiber volume fraction is within the range 50-70% and therefore inter-fiber spacing becomes very small, of the order of microns or smaller, leading to multiple fiber-to-fiber contacts when consolidation forces are applied during processing. The external pressure is initially supported by the resin and, as bleeding progresses, pressure is transferred to the fiber bed. This process continues until the composite reaches the maximum compaction of the fibers for the applied external pressure 22.

(51) 2.2 Governing Equations and no more resin can be squeezed-out. The load carried by the fibers becomes appreciable for fiber volume fractions in the range 60 to 70%, Gutowski et al. (1986) (Fig. 2.6).. Figure 2.6: Pressure carried by the fibers as a function of the fiber volume fraction, Gutowski et al. (1986).. The effective stress theory, originally developed to study soil consolidation, Terzaghi (1943) and Biot (1941), was applied by Gutowski et al. (1987a,b) and Davé et al. (1987) to study the compaction of fiber beds. The partition of the stress tensor between fibers and resin can be expressed as,. σij = σij0 − Pr δij. (2.16). where σij is the total Cauchy stress tensor, σij0 stands for the effective stress carried by the fiber bed, Pr is the resin pressure and δij the Kronecker delta (δij = 1 for i = j and δij = 0 for i 6= j). The fiber bed effective stress tensor can be related to the strain tensor through the following elastic constitutive equation,. ij = Sij σij0 23. (2.17).