TECHNISCHE UNIVERSITÄT MÜNCHENLehrstuhl für Carbon Composites

Gradual Impregnation during theProduction of Thermoplastic Composites

Veronika Anna Bühler

Vollständiger Abdruck der von der Fakultät für Maschinenwesen der Technischen Univer-sität München zur Erlangung des akademischen Grades eines

Doktor-Ingenieurs

genehmigten Dissertation.

Vorsitzender: Prof. Dr.-Ing. Veit Senner

Prüfer der Dissertation: Prof. Dr.-Ing. Klaus Drechsler

Prof. Paul Compston, Ph.D.

Die Dissertation wurde am 12.04.2017 bei der Technischen Universität München eingere-icht und durch die Fakultät für Maschinenwesen am 21.06.2017 angenommen.

Technische Universität MünchenFakultät für MaschinenwesenLehrstuhl für Carbon CompositesBoltzmannstraße 15D-85748 Garching bei München

Tel.:+ 49 (0) 89 / 289 - 15092Fax: + 49 (0) 89 / 289 - 15097Email: info@lcc.mw.tum.deWeb: www.lcc.mw.tum.de

FürArco Oma & OpaSteiner Oma & Opa

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DeclarationIch erkläre hiermit ehrenwörtlich, dass ich die vorliegende Arbeit selbstständig undohne Benutzung anderer als der angegebenen Hilfsmittel angefertigt habe; die ausfremden Quellen (einschließlich elektronischer Quellen) direkt oder indirekt über-nommenen Gedanken sind ausnahmslos als solche kenntlich gemacht.

Die Arbeit wurde in gleicher oder ähnlicher Form noch keiner anderen Prüfungs-behörde vorgelegt.

...................................Ort, Datum

...................................Unterschrift

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AcknowledgementI would like to express my gratitude to my supervisor Prof. Klaus Drechsler forgiving me the opportunity to write a PhD thesis at the Chair of Carbon Compos-ites (LCC). Being a part of the institute’s Material Behavior and Testing group Iwas guided by Dr. Hannes Körber. I want to thank him and the deputy head ofthe institute, Dr. mont. Elisabeth Ladstätter, for their constant as well as indis-pensable support and for giving me sufficient confidence to work independently.I express my deepest thanks to my co-supervisor Prof. Paul Compston for the fruit-ful discussions we had, his valuable technical input to my thesis and for makingthis long journey from Australia to take part at my examination. He also enableda two-month research period at the Australian National University in Canberrawhere I got to know Sherman Wang, an expert in nano-indentation, who greatlysupported me and gave valuable input to my thesis. I also want to express mygratitude to Dr. Christopher Stokes-Griffin for valuable discussions, great supportand especially for the fantastic time when we shared an office. Thank you!My PhD project was made possible by the financial support from SGL CarbonGmbH. I am very thankful for their generosity and highly appreciate the close col-laboration, the many valuable discussions and the great support by Dr. AndreasErber, Patrik-Vincent Brudzinski, Dr. Steffen Janetzko, Dr. Christian Stang, Dr.Oswin Öttinger and Veronika Hirschinger.In addition, my great time at the LCC owes to my colleagues. I specially want tothank Cigdem Filker for her perpetual support in administrative matters and forbeing not only a wonderful office colleague but also a valued friend. I also highlyappreciate the technical discussions, support, friendship, chats and coffees with JanKrollmann, Luciano Avila Gray, Rhena Helmus, Philipp Hörmann, Philipp Bruck-bauer, Alex Schwingenschlögel, Stefan Ehard, Andreas Kollmannsberger, ThorstenHans, Philipp Picard, Marina Plöckl, Peter Kuhn, Ludwig Eberl, Philipp Fahr,Christoph Ebel and Swen Zaremba. I also would like to thank the workshop forenabling almost impossible things.It was also a great pleasure to work with excellent students during my time at theLCC. Thank you very much Christian Heckel, Heiko Baumann, Patrick Consul,Miriam Ernst, Adrián García López and Stefan Ender.Moreover, I’d like to thank my family and friends for their indispensable encour-agement and great support not only during my PhD but all my life. Thank you somuch, Mama, Papa, Steffie and Sebastian for everything! Thank you, Joe, for yourloving support, incredible encouragement and unbowed belief in me!

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AbstractCommonly, pre-formed or pre-impregnated intermediates are used to produce car-bon fiber reinforced thermoplastics (CFRTP). However, impregnation of the fiberwith the polymer is a time-consuming process, due to the high melt viscosity of ther-moplastics leading to high manufacturing costs. Existing approaches to overcomethe challenges during impregnation deal with advancements in the impregnationtechnology, improvements in process control or the use of low viscous prepolymers.In this work, partially impregnated tapes are developed which can be manufacturedwith increased production rates to reduce the costs of intermediates. The partiallyimpregnated tapes are intended to completely impregnate throughout subsequentheating and consolidation processes, required to produce a final component madefrom CFRTP.To begin with, the fiber-matrix compatibility between various polyamide types andcarbon fibers with different sizings was investigated to produce high-performancecomposites from powder-coated tows. Based on results from macro- and micro-mechanical tests, the most suitable material combinations were selected for subse-quent investigations.To enable the comparability of production methods that are characterized by signif-icantly different cooling rates, the crystallization kinetics of the selected polyamideswas studied. This study on neat polymers, as well as in presence of differently sizedcarbon fibers, was conducted by using the differential scanning calorimetry. Usinga new method to determine the crystallized fraction in fiber reinforced polymers,mechanical properties were correlated to different cooling rates used for the pro-duction of test panels.Based on literature review, the transverse impregnation was modeled using Darcy’slaw. To verify the derived model experimentally, an impregnation study was con-ducted to produce test panels from powder-coated tows by varying the three mainprocess parameters that drive impregnation: time, temperature and pressure. Theexperimental design followed the design of experiments to consider extreme processsettings, suitable for model verification. By post-processing micrographs from thetest panels, an experimental procedure was developed to determine the degree ofimpregnation. Comparing the experimental results obtained from the impregnationstudy with the values predicted by the model, a good correlation was found for theselected material combinations.To enable gradual impregnation throughout component production, process-relatedeffects on the polymer influencing the impregnation behavior were investigated.Temperature profiles with different dwell times were derived from a typical CFRTPproduction process. The selected polyamides were subjected to these profiles and

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analyzed, using differential scanning calorimetry, thermogravimetric analysis, rheom-etry and gel permeation chromatography. The exposure of the selected polyamidesto the temperature profiles yielded significant increases in the melt viscosity dueto thermo-oxidative degradation. A processing window of 5 minutes was identifiedfor the various process steps where gradual impregnation is enabled. Beyond thisprocessing window, the observed increase in viscosity can prevent complete impreg-nation during the actual process step or in the following steps.By adding a suitable antioxidant, the extent of the degradation reactions was re-duced and substantial increases in viscosity were limited. Additional modificationof the selected polyamides with a lubricant further decreased the viscosity leadingto reduced impregnation times.Eventually, partially impregnated tapes were produced from powder-coated towsin a double-belt press with increased production rates. With a completely im-pregnated tape as a reference, the partially impregnated tapes were processed bypress forming and thermoforming with different dwell times to simulate a typicalproduction of a CFRTP component. The impregnation of partially impregnatedintermediates was found to be completed after the repeated heating processes dur-ing CFRTP production. Studying the flexural properties of test panels producedfrom the partially impregnated tapes, comparable values to completely impreg-nated tapes were achieved upon press forming for 5 to 10 minutes. The analysis ofthe manufacturing costs yielded cost savings between 40 to 60 % as the partiallyimpregnated tapes can be produced with double or quadruple production rates,compared to completely impregnated tapes.

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KurzfassungZur Herstellung von carbonfaserverstärkten Thermoplasten (CFTP) werden übli-cherweise thermoplastische Halbzeuge verwendet, wobei die Fasern meist bereitsvollständig mit der thermoplastischen Kunststoffmatrix imprägniert sind. Die Im-prägnierung der Fasern mit der Matrix ist aufgrund der hohen Schmelzviskositätvon Thermoplasten jedoch ein zeit- und somit kostenintensiver Prozess. BisherigeAnsätze zur Bewältigung dieser Herausforderung bestehen in der Weiterentwick-lung der Imprägnierungstechnologien, der Verbesserung der Prozessführung oderder Verwendung von niederviskosen Vorpolymeren.In der vorliegenden Arbeit wird die Entwicklung von teilimprägnierten Tapes vor-gestellt. Derartige Tapes können mit höheren Prozessgeschwindigkeiten produziertwerden, um so die Herstellungskosten zu senken. Diese teilimprägnierten Tapessollen im Laufe der anschließenden, zur Bauteilherstellung nötigen Aufheiz- sowieKonsolidierungsprozesse vollständig imprägniert werden.Zunächst wurde die Faser-Matrix-Kompatibilität zwischen verschiedenen Polyamid-typen und Carbonfasern mit unterschiedlichen Schlichten untersucht, um hochleis-tungsfähige Verbundwerkstoffe aus pulverbeschichteten Fasern herzustellen. Basie-rend auf den Ergebnissen aus makro- sowie mikromechanischen Materialprüfungenwurden die Materialkombinationen mit der höchsten Kompatibilität von Faser,Schlichte und Matrix abgeleitet.Um Produktionsprozesse mit signifikant unterschiedlichen Kühlraten miteinandervergleichen zu können, wurde die Kristallisationskinetik der ausgewählten Polyami-de bestimmt. Die Untersuchung mittels der Differenzkalorimetrie erfolgte sowohlan reinen Polymeren als auch an den hergestellten Verbundwerkstoffen. Durch dieVerwendung einer neuen Methode zur Bestimmung des kristallinen Anteils in fa-serverstärkten Kunststoffen konnten die mechanischen Kennwerte mit den Kris-tallisationsgraden, die sich während der Herstellung von Prüfplatten aufgrund derverschiedenen Kühlraten einstellen, korreliert werden.Basierend auf einer Literaturübersicht wurde die Imprägnierung in Dickenrichtungmodellhaft mit dem Gesetz von Darcy beschrieben und eine Imprägnierstudie zurexperimentellen Verifikation durchgeführt. Dazu wurden in einer statischen Pres-se Prüfplatten aus pulverbeschichteten Fasern mit verschiedenen Kombinationender Prozessparameter Zeit, Druck und Temperatur, die die Imprägnierung maß-geblich beeinflussen, hergestellt. Zur Versuchsplanung wurde Design of Experi-ments verwendet, um auch extreme Prozessparametersätze zu berücksichtigen. DieEntwicklung einer automatisierten Nachbearbeitung von Schliffbildern der in derImprägnierstudie hergestellten Prüfplatten erlaubte die Bestimmung des sich ein-stellenden Imprägnierungsgrades. Der Abgleich der Ergebnisse für den Imprägnie-

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rungsgrad aus der Imprägnierstudie mit den errechneten Werten ergab eine guteÜbereinstimmung zwischen Modell und Realität.Um eine graduelle Imprägnierung während der Bauteilherstellung zu ermöglichen,wurden prozessinduzierte Effekte auf die Matrixeigenschaften untersucht, welchedas Imprägnierungsverhalten beeinflussen. Dazu wurden zunächst Temperaturpro-file aus typischen Produktionsprozessen für thermoplastische Verbundwerkstoffeabgeleitet. Die Auswirkungen der Temperaturprofile auf die Polymereigenschaf-ten wurden mithilfe der Differenzkalorimetrie, der thermogravimetrischen Analyse,der Rheometrie sowie der Gelpermeationschromatographie untersucht. Die thermi-sche Belastung aufgrund der angelegten Temperaturprofile verursachte einen star-ken Anstieg der Viskosität der untersuchten Polyamidtypen. Für die verschiedenenProzessschritte wurde ein Prozessfenster von 5 Minuten identifiziert, in welchemdie graduelle Imprägnierung ungehindert stattfinden kann. Außerhalb dieses Zeit-fensters kommt es zum beobachteten Viskositätsanstieg, welcher eine vollständigeImprägnierung im vorliegenden oder nachfolgenden Prozessschritt erschweren kann.Durch die Beigabe eines geeigneten Antioxidanten wurde das Ausmaß der Degrada-tionsreaktionen und der Viskositätsanstieg begrenzt. Eine zusätzliche Modifikationder ausgewählten Polyamidtypen mit einem Fließmittel führte zu einer weiterenSenkung der Viskosität und ermöglicht somit eine erhebliche Verkürzung der Im-prägnierungszeit.Schlussendlich wurden teilimprägnierte Tapes in einer Doppelbandpresse mit er-höhten Produktionsgeschwindigkeiten hergestellt. Diese sowie vollkommen imprä-gnierte Tapes wurden in Press- und Thermoformverfahren weiterverarbeitet, umeine typische Bauteilherstellung zu simulieren. Nach den wiederholten Aufheiz- undKonsolidierungsprozessen zur Bauteilherstellung wurde eine vollständige Tränkungder zunächst teilimprägnierten Tapes erreicht. Verglichen mit vollständig imprä-gnierten Tapes wurden nach dem Pressen für 5 bis 10 Minuten vergleichbare Bie-geeigenschaften für die teilimprägnierten Tapes erzielt. Eine anschließende Kosten-analyse ergab ein potentielles Einsparpotential von 40 bis 60 % für die Herstellungvon teilimprägnierten Tapes, da diese mit der doppelten beziehungsweise vierfachenProzessgeschwindigkeit produziert werden können.

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ContentsContents xviii

List of Figures xxv

List of Tables xxxi

1 Introduction 11.1 Motivation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11.2 Thermoplastic composites . . . . . . . . . . . . . . . . . . . . . . . 2

1.2.1 Intermediate materials . . . . . . . . . . . . . . . . . . . . . 31.2.2 Production methods . . . . . . . . . . . . . . . . . . . . . . 6

1.3 State-of-the-art . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81.4 Objectives and outline of the thesis . . . . . . . . . . . . . . . . . . 9

2 Fiber-matrix compatibility 132.1 Investigated materials . . . . . . . . . . . . . . . . . . . . . . . . . 13

2.1.1 Carbon fibers . . . . . . . . . . . . . . . . . . . . . . . . . . 132.1.2 Polyamides . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

2.2 Experimental methods . . . . . . . . . . . . . . . . . . . . . . . . . 162.2.1 Four-point bend test . . . . . . . . . . . . . . . . . . . . . . 162.2.2 Double-cantilever beam test . . . . . . . . . . . . . . . . . . 172.2.3 Statistics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 182.2.4 Nano-indentation . . . . . . . . . . . . . . . . . . . . . . . . 182.2.5 Scanning electron microscopy . . . . . . . . . . . . . . . . . 20

2.3 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . 202.3.1 Intermediate production . . . . . . . . . . . . . . . . . . . . 202.3.2 Test panel production . . . . . . . . . . . . . . . . . . . . . 212.3.3 Test specimen preparation . . . . . . . . . . . . . . . . . . . 222.3.4 Micrographs . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

2.4 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 232.4.1 Influence of the sizing . . . . . . . . . . . . . . . . . . . . . 232.4.2 Influence of matrix ductility . . . . . . . . . . . . . . . . . . 282.4.3 Development of an interphase . . . . . . . . . . . . . . . . . 31

2.5 Selection of compatible material combinations . . . . . . . . . . . . 33

3 Crystallization of polyamides 353.1 Theory . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

3.1.1 Crystallization and nucleation . . . . . . . . . . . . . . . . . 353.1.2 Crystallization kinetics . . . . . . . . . . . . . . . . . . . . . 37

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3.2 Experimental methods . . . . . . . . . . . . . . . . . . . . . . . . . 383.2.1 Differential scanning calorimetry . . . . . . . . . . . . . . . 393.2.2 Crystallinity ratio . . . . . . . . . . . . . . . . . . . . . . . . 403.2.3 Four-point bend test . . . . . . . . . . . . . . . . . . . . . . 413.2.4 Visualization of crystals . . . . . . . . . . . . . . . . . . . . 41

3.3 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 423.3.1 Preliminary tests . . . . . . . . . . . . . . . . . . . . . . . . 423.3.2 Neat polymers . . . . . . . . . . . . . . . . . . . . . . . . . . 443.3.3 Influence of carbon fibers on crystallization . . . . . . . . . . 513.3.4 Relation between mechanical properties and crystallinity ratio 53

3.4 Conclusion and implications . . . . . . . . . . . . . . . . . . . . . . 55

4 Impregnation model 574.1 Transverse resin flow . . . . . . . . . . . . . . . . . . . . . . . . . . 574.2 Processing phenomena of individual constituents . . . . . . . . . . . 59

4.2.1 Fiber bed properties . . . . . . . . . . . . . . . . . . . . . . 594.2.2 Matrix . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

4.3 1D through thickness model . . . . . . . . . . . . . . . . . . . . . . 644.3.1 Assumptions . . . . . . . . . . . . . . . . . . . . . . . . . . . 644.3.2 Model derivation . . . . . . . . . . . . . . . . . . . . . . . . 64

4.4 Experimental work . . . . . . . . . . . . . . . . . . . . . . . . . . . 664.4.1 Rheology . . . . . . . . . . . . . . . . . . . . . . . . . . . . 664.4.2 Experimental determination of degree of impregnation . . . 684.4.3 Design of experiments . . . . . . . . . . . . . . . . . . . . . 694.4.4 Interlaminar shear test . . . . . . . . . . . . . . . . . . . . . 71

4.5 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 714.5.1 Viscosity data . . . . . . . . . . . . . . . . . . . . . . . . . . 724.5.2 Influence of processing on impregnation progress . . . . . . . 744.5.3 Model calibration . . . . . . . . . . . . . . . . . . . . . . . . 784.5.4 Influence of degree of impregnation on interlaminar shear

strength . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 804.6 Conclusion and implications . . . . . . . . . . . . . . . . . . . . . . 80

5 Degradation of polyamides 835.1 Literature review . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83

5.1.1 Thermal degradation . . . . . . . . . . . . . . . . . . . . . . 835.1.2 Thermo-oxidative degradation . . . . . . . . . . . . . . . . . 835.1.3 Post-condensation . . . . . . . . . . . . . . . . . . . . . . . . 855.1.4 Influence of degradation on processing . . . . . . . . . . . . 86

5.2 Temperature profiles . . . . . . . . . . . . . . . . . . . . . . . . . . 87

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5.3 Experimental methods . . . . . . . . . . . . . . . . . . . . . . . . . 885.3.1 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . 885.3.2 Differential scanning calorimetry . . . . . . . . . . . . . . . 885.3.3 Thermogravimetric analysis . . . . . . . . . . . . . . . . . . 895.3.4 Rheometry . . . . . . . . . . . . . . . . . . . . . . . . . . . . 905.3.5 Gel permeation chromatography . . . . . . . . . . . . . . . . 90

5.4 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 925.4.1 Influence on melting temperature . . . . . . . . . . . . . . . 925.4.2 Effect on mass loss . . . . . . . . . . . . . . . . . . . . . . . 935.4.3 Impact on complex viscosity . . . . . . . . . . . . . . . . . . 945.4.4 Effect on molecular composition . . . . . . . . . . . . . . . . 965.4.5 Processing window . . . . . . . . . . . . . . . . . . . . . . . 975.4.6 Conclusion and implications . . . . . . . . . . . . . . . . . . 98

6 Thermal stabilization and flow promotion of polyamides 1016.1 Thermal stabilization . . . . . . . . . . . . . . . . . . . . . . . . . . 101

6.1.1 Chain breaking antioxidants . . . . . . . . . . . . . . . . . . 1016.1.2 Radical scavengers . . . . . . . . . . . . . . . . . . . . . . . 1036.1.3 Preventive antioxidants . . . . . . . . . . . . . . . . . . . . . 1036.1.4 Investigated antioxidants . . . . . . . . . . . . . . . . . . . . 104

6.2 Flow promotion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1066.2.1 Internal and external lubricants . . . . . . . . . . . . . . . . 1066.2.2 Investigated lubricants . . . . . . . . . . . . . . . . . . . . . 108

6.3 Experimental methods . . . . . . . . . . . . . . . . . . . . . . . . . 1096.3.1 Polymer samples . . . . . . . . . . . . . . . . . . . . . . . . 1096.3.2 Composite samples . . . . . . . . . . . . . . . . . . . . . . . 110

6.4 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1126.4.1 Effectiveness of additives . . . . . . . . . . . . . . . . . . . . 1126.4.2 Effects of combined use of additives . . . . . . . . . . . . . . 1166.4.3 Influence of matrix modification on impregnation . . . . . . 121

6.5 Conclusion and implications . . . . . . . . . . . . . . . . . . . . . . 122

7 Gradual impregnation during production 1257.1 Manufacture of differently impregnated tapes . . . . . . . . . . . . . 1257.2 Manufacture of thermoplastic composites . . . . . . . . . . . . . . . 128

7.2.1 Prediction of dwell times . . . . . . . . . . . . . . . . . . . . 1287.2.2 Test panel production . . . . . . . . . . . . . . . . . . . . . 1297.2.3 Four-point bend test . . . . . . . . . . . . . . . . . . . . . . 129

7.3 Influence of initial degree of impregnation on mechanical properties 1307.3.1 Final degree of impregnation after production . . . . . . . . 130

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7.3.2 Flexural properties . . . . . . . . . . . . . . . . . . . . . . . 1327.4 Cost analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 134

7.4.1 Procedure . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1357.4.2 Data collection . . . . . . . . . . . . . . . . . . . . . . . . . 1377.4.3 Monetary effect of partially impregnated tapes . . . . . . . . 139

7.5 Correlation of mechanical properties and manufacturing costs . . . 1407.6 Conclusion and implications . . . . . . . . . . . . . . . . . . . . . . 141

8 Summary and outlook 1438.1 Summary and conclusion . . . . . . . . . . . . . . . . . . . . . . . . 1438.2 Future work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 148

Bibliography 151

A Appendix 169A.1 to Section 4.4.3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 169A.2 to Section 6.4 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171A.3 to Section 7.3.1 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171

B Publications 173

C Supervised student theses 175

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NomenclatureSymbol Unit Meaning

a [mm] Delamination lengtha0 [mm] Initial delamination lengthA [nm2] Project area of Berkovich indenterA1 [-] Constant for modified compliance calibrationAs [Pa] Empirical constantaT [-] Temperature shift factorC [mm/N] Compliance orC [°C] Cooling rateC1 [-] Fiber packing constantCR [-] Crystallinity ratioDOI [%] Degree of impregnationDOIf [%] Final degree of impregnationDOIi [%] Initial degree of impregnationE [GPa] Elastic modulusEa [kJ/mol] Activation energyEf1 [GPa] Longitudinal flexural modulusEf2 [GPa] Transverse flexural modulusG′ [MPa] Storage modulusG′′ [MPa] Loss modulusGI [J/m2] Mode I interlaminar fracture toughnessGIc [J/m2] Opening mode I interlaminar fracture toughnessH [GPa] HardnessΔHc,∞ [J/g] Enthalpy of crystallization at the end of the crys-

tallization processΔHc [J/g] Enthalpy of crystallizationΔHcc [J/g] Enthalpy of cold crystallizationΔHf [J/g] Enthalpy of fusionΔH0

f [J/g] Enthalpy of fusion of a 100% crystalline polymerhp [nm] Plastic depth of penetrationILS [MPa] Interlaminar shear strengthk [10−2min−n] Crystallization rate constant according to AvramiK [m2] PermeabilityK(T ) [-] Cooling functionKzz [m2] Permeability in z-directionkzz [-] Kozeny constantk′

zz [-] Kozeny constant, modified by Gutowski

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L [mm] Support spanL′ [mm] Load spanl [mm] Specimen lengthm [-] Ozawa exponent orm [-] Population meanMn [g/mol] Molar massMw [-] Molar weightMWD [-] Molecular weight distributionn [-] Number of measurements orn [-] Avrami exponentnd [-] Avrami exponent for crystal dimensionnn [-] Avrami exponent for crystallization typeOIT [min] Oxidation induction timep,P [bar] PressurePmax [N] Maximum forceR [J/molK3] Universal gas constantrf [m] Fiber radiuss [mm] Deflection ors [-] Standard deviationT [°C] Temperaturet [s] or [min] Time ort [mm] Specimen thicknesst0.925 [-] Confidence level of 95%Tg [°C] Glass transition temperatureTm [°C] Melting temperatureTp [°C] Peak crystallization temperatureuf [mm/s] Fiber bed velocity vectorum [mm/s] Matrix velocity vectorV0 [%] Initial fiber volume contentVa [%] Maximum possible fiber volume contentV ′

a [%] Maximum available fiber volume contentVc [Vol.-%] Volume fraction of crystalline phaseVf [%] Fiber volume fractionVf ,max [%] Maximum fiber volume contentw [mm] Specimen widthWc [wt%] Crystallized mass fractionx [-] Arithmetic mean of n measurementsX [%] Relative degree of crystallinityα, λ, n [-] Fitting parameters Carreau-Yasuda modelδ [mm] Load point deflection or

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δ [-] Phase lagεf [%] Strain of the outer fiberη [Pa s] Viscosityη′ [Pa s] In-phase component of complex viscosityη′′ [Pa s] Out-of-phase component of complex viscosityη∗ [Pa s] Complex viscosityη0 [Pa s] Zero-shear viscosityγ [1/s] Shear rateγ0 [-] Strain amplitudeφ [wt%] Fiber weight fractionρ [g/cm3] Densityρa [g/cm3] Density of amorphous polymer fractionρc [g/cm3] Density of crystalline polymer fractionσ [MPa] Stressσf [MPa] Flexural stressσf1 [MPa] Longitudinal flexural strengthσf2 [MPa] Transverse flexural strengthσyield [MPa] Yield stressτ , τ0 [Pa] Shear stress, shear stress amplitudeξ [°] Angle of twist

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List of Acronyms

Abbreviation Meaning

AFP Automated fiber placementAMMRF Australian Microscopy and Microanalysis Research Fa-

cilityANU Australian National University, CanberraASTM American Society for Testing and MaterialsATL Automated tape layingB3L Toughened polyamide 6 grade from BASFB3S Low-flow polyamide 6 grade from BASFB40 Polyamide 6 grade with high molecular weight from

BASFC2000 Semi-aromatic co-polyamide PA10T/X from EvonikCAM Centre for Advanced MicroscopyCB Chain breaking; group of antioxidantsCB-A Chain breaking acceptorCB-D Chain breaking donorCF-EPY SIGRAFIL C T50-4.0/240-E100 carbon fibers from SGL

GroupCFRP Carbon fiber reinforced plasticsCFRTP Carbon fiber reinforced thermoplasticsCF-TP SIGRAFIL C T50-4.0/240-T140 carbon fibers from SGL

GroupDCB Double-cantilever beamDIN Deutsches Institut für Normung e.V.DOE Design of experimentDSC Differential scanning calorimetryEBS Bis-stearyl ethylenediamine / ethylenbisstearamideEC European CommunityFIT Fibre impregnée thermoplastiqueGMT Glass-mat reinforced thermoplasticsGPC/SEC Gel-permeation chromatography/ Size-exclusion chro-

matographyLFT Long-fiber reinforced thermoplasticsLPL Laboratory prepreg lineMCC Modified compliance calibrationNCF Non-crimped fabricOoA Out-of-Autoclave

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PA10T/X Co-polyamide of the group of polyphthalamidesPA6 Polyamide 6PEEK PolyetheretherketonePEI PolyetherimidePES Polyether sulfonePPA PolyphthalamidePPS Polyphenylene sulfidePVC Polyvinyl chlorideRFI Resin film infusionRIM Reaction injection moldingRTM Resin transfer moldingSCB Side-clamped beamSEM Scanning electron microscopyTGA Thermogravimetric analysisT-RTM Thermoplastic-resin transfer moldingUD Unidirectional

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List of Figures1-1 Classification of thermoplastics [8]. . . . . . . . . . . . . . . . . . . 31-2 Overview of potential manufacturing routes originating from the pro-

duction of intermediates for thermoplastic composites, based on [7,10, 11]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

1-3 Principle process steps during thermoplastic composite productionalong with the governing process parameters time, pressure and tem-perature; based on [29]. . . . . . . . . . . . . . . . . . . . . . . . . . 6

1-4 Principle production steps to manufacture thermoplastic composites;pictures provided as courtesy by SGL Group and from Celanese [44]as well as from the Institut für Verbundwerkstoffe (IVW) Kaiser-slautern [45], as indicated. . . . . . . . . . . . . . . . . . . . . . . . 10

1-5 Schematic structure of the present thesis. . . . . . . . . . . . . . . . 11

2-1 Chemical structure of aliphatic PA6. . . . . . . . . . . . . . . . . . 152-2 Chemical structure of semi-aromatic co-polyamide (PA10T). . . . . 152-3 Four-point bend test setup with support span L and a load span

L’=L/3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 162-4 a) DCB test specimen with initial delamination length a0 from load

line to end of insert and b) test specimen clamped to the SCB testfixture mounted to a Hegewald & Peschke 100 kN universal testingmachine. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

2-5 a) Principle of nano-indentation on carbon fibers surrounded by ma-trix; b) SEM image of Berkovich indenter tip [72]. . . . . . . . . . . 19

2-6 Schematic of the prepreg line used to produce powder-coated towson a laboratory scale. . . . . . . . . . . . . . . . . . . . . . . . . . . 21

2-7 a) Stacking, b) processing in a static press and c) produced test panel. 222-8 a) Stress-strain curves of CF-EPY/B3S and b) CF-TP/B3S. . . . . 242-9 R curves for a) CF-EPY/B3S and b) CF-TP/B3S. . . . . . . . . . . 242-10 a) Mean transverse flexural strength σf2 and b) mean mode I inter-

laminar fracture toughness GIc of CF-EPY/B3S in comparison toCF-TP/B3S. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

2-11 Fracture surface analysis of tested DCB specimens made of a) CF-EPY/B3S in comparison to b) CF-TP/B3S. . . . . . . . . . . . . . 25

2-12 Stress-strain curves for a) CF-EPY/C2000 and b) CF-TP/C2000. . 262-13 a) R curves of CF-EPY/C2000 and b) CF-TP/C2000. . . . . . . . . 262-14 a) Mean transverse flexural strength σf2 and b) mean mode I inter-

laminar fracture toughness GIc of CF-EPY/C2000 in comparison toCF-TP/C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

xxv

2-15 Fracture surface analysis of tested flexural specimens made of a)CF-EPY/C2000 compared to b) CF-TP/C2000. . . . . . . . . . . . 27

2-16 Micrographs of four-point bend test panels made of a) CF-EPY/C2000and to b) CF-TP/C2000. . . . . . . . . . . . . . . . . . . . . . . . . 28

2-17 Fracture surface analysis of tested DCB specimens made of a) CF-EPY/C2000 compared to b) CF-TP/C2000. . . . . . . . . . . . . . 28

2-18 Stress-strain curves of a) CF-TP/B3L in comparison to b) CF-TP/B40. 292-19 R curves for a) CF-TP/B3L and b) CF-TP/B40. . . . . . . . . . . 292-20 a) Mean transverse flexural strength σf2 and b) mean mode I inter-

laminar fracture toughness GIc of CF-TP/B3L and CF-TP/B40 incomparison to CF-TP/B3S. . . . . . . . . . . . . . . . . . . . . . . 30

2-21 Fracture surface analysis of tested DCB specimens made of a) CF-TP/B3S, b) CF-TP/B3L and c) CF-TP/B40. . . . . . . . . . . . . 30

2-22 Nano-indentation on a) CF-TP only and b) B3S only. . . . . . . . . 312-23 Nano-indentation on carbon fibers coated with a) epoxy-based and

b) thermoplastic-based sizing, surrounded by B3S; blue arrow indi-cates indenting direction and covered area. . . . . . . . . . . . . . . 32

2-24 Close-up view of a) CF-EPY fibers with highlighted gap betweenfiber and matrix, b) CF-TP fibers, surrounded by B3S, under theSEM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

3-1 Time-delayed development of nucleation and growth rate duringcrystallization as a function of temperature, redrawn from [75]. . . . 36

3-2 Rucks thermoforming unit and used aluminum tool to perform ther-moforming of flat test panels. . . . . . . . . . . . . . . . . . . . . . 42

3-3 Influence of pre-drying process on crystallization behavior of a) un-dried and b) pre-dried B3S. . . . . . . . . . . . . . . . . . . . . . . 43

3-4 Subjection of C2000 to ten heating and cooling cycles. . . . . . . . 443-5 Development of the relative degree of crystallinity X(t) for a) neat

B3S and b) C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . 453-6 Plot of log

{− ln[1−X(t)]

}versus logt for the isothermal crystalliza-

tion of a) B3S with ρa = 1.08 g/cm3 and ρc = 1.24 g/cm3 b) C2000,based on Wc according to [98]. . . . . . . . . . . . . . . . . . . . . . 46

3-7 Development of the relative degree of crystallinity X(T ) for a) neatB3S and b) C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

3-8 Plots of log{

− ln[1−X(T )]}

versus ln|(dT/dt)−1| for a) neat B3S at205, 200, 195, 190 and 185 °C; b) neat C2000 at 235, 230, 225, 215and 205 °C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

3-9 Plots of cooling rate versus enthalpy of crystallization ΔHc for B3Sand C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

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3-10 Influence of carbon fiber sizing on ΔHc for tapes made of a) B3Sand b) C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51

3-11 Close-up view of B3S specimens with a) epoxy-sized fibers and b)polyamide-sized fibers that were etched by oxygen plasma to visual-ize crystalline structures . . . . . . . . . . . . . . . . . . . . . . . . 52

3-12 Influence of CR on σf2 and Ef2 of a), c) CF-TP/B3S and b), d)CF-TP/C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54

4-1 Idealized arrangements of fiber packings in a composite yieldingmaximum fiber volume contents Vf,max of 78.5 % (quadratic) and90.7 % (hexagonal). . . . . . . . . . . . . . . . . . . . . . . . . . . . 60

4-2 Comparison of production processes for thermoplastics with regardto shear rates; redrawn from [133]. . . . . . . . . . . . . . . . . . . 63

4-3 Schematic of the flow front progression according to the derived 1Dthrough thickness model. . . . . . . . . . . . . . . . . . . . . . . . . 65

4-4 a) Parallel-plate and b) cone-plate fixtures for use in rotationalrheometers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

4-5 a) Oscillatory measurement and b) time-delayed shift of stress re-sponse compared to applied strain rate; redrawn from [137]. . . . . 66

4-6 Schematic of the degree of impregnation of thermoplastic interme-diates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68

4-7 Micrographs of cross-sections of tapes with a) highlighted entire fiberbundle area, b) highlighted non-impregnated area and binary pic-tures of c) the entire fiber bundle area and d) the non-impregnatedarea within the fiber bundles. . . . . . . . . . . . . . . . . . . . . . 69

4-8 Graphical representation of a Box-Behnken Design with three factorshaving two extreme factor levels (−1 (minimum) and 1 (maximum))and a center point (0); redrawn from [142]. . . . . . . . . . . . . . . 70

4-9 Viscosity curves at different temperatures for a) B3S, B3L, B40 andb) C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4-10 Plots of temperature shift factor aT versus 1/RT for a) B3S, B3L,B40 and b) C2000; the slope of the lines represents the activationenergy Ea. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73

4-11 Viscosity curve of B3S calculated according to Arrhenius for a typicaltemperature profile during thermoforming. . . . . . . . . . . . . . . 73

4-12 Micrographs of CF-TP/B3S test panels after a press time of a) 1minute revealing non-impregnated fiber bundles and b) 10 minutesshowing complete impregnation. . . . . . . . . . . . . . . . . . . . . 74

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4-13 Micrographs of CF-TP/B3S test panels pressed with a) 260 °C show-ing non-impregnated fiber bundles and b) 280 °C revealing an ad-vanced impregnation progress. . . . . . . . . . . . . . . . . . . . . . 75

4-14 Micrographs of CF-TP/B3S test panels pressed with a) 5 bar re-vealing non-impregnated fiber bundles and b) 30 bar with increasedDOI. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

4-15 Main effect plot for DOI of CF-TP/B3S. . . . . . . . . . . . . . . . 764-16 Interaction plot for DOI of CF-TP/B3S. . . . . . . . . . . . . . . . 764-17 Main effect plot for DOI of CF-TP/C2000. . . . . . . . . . . . . . . 774-18 Interaction plot for DOI of CF-TP/C2000. . . . . . . . . . . . . . . 774-19 Development of the experimentally determined Vf (◦) as a function

of applied pressure for a) CF-TP/B3S and b) CF-TP/C2000. . . . . 784-20 Comparison of experimentally determined (◦) and calculated DOI

for various temperatures at 17.5 bar for a-c) CF-TP/B3S and d-f)CF-TP/C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79

4-21 ILSS and the DOI for a) B3S at constant time, b) C2000 at constanttime, c) B3S at constant temperature, d) C2000 at constant temper-ature, e) B3S at constant pressure, f) C2000 at constant pressure. . 81

5-1 Principle scheme of thermal decomposition of PA6 [145]. . . . . . . 845-2 Basic mechanism for chain scission during oxidation of aliphatic

polyamides [156]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 855-3 Considered CFRTP production process to derive temperature profiles. 875-4 Temperature profiles P1 and P2 for a) B3S and ) C2000 derived from

a CFRTP production process. . . . . . . . . . . . . . . . . . . . . . 885-5 Determination of Tm for bimodal melt peaks as present for C2000. . 895-6 Procedure to determine the mass loss that has occurred during every

process step of P1 and P2 for B3S as an example. . . . . . . . . . . 905-7 Size separation and detection of dissolved molecules by GPC; re-

drawn from [173]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 915-8 DSC thermograms for a) B3S and b) C2000 subjected to temperature

profile P2 in air. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 925-9 Development of Tm of a) B3S and b) C2000 under air and nitrogen

atmosphere when subjected to temperature profiles P1 and P2. . . . 925-10 Mass loss of a) B3S and b) C2000 samples subjected to temperature

profile P1 and P2 under air and nitrogen gas atmosphere in TGA. . 945-11 Development of the complex viscosity η∗ of a) B3S and b) C2000

during subjection to temperature profile P2 in air and nitrogen gasatmosphere. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 94

xxviii

5-12 Minimum complex viscosity η∗ of a) B3S and b) C2000 subjected totemperature profiles P1 and P2 in air and nitrogen atmosphere. . . 95

5-13 MWD of a) B3S and b) C2000 as-received and after exposure totemperature profile P2 under air and nitrogen gas atmosphere. . . . 96

5-14 Development of the DOI of CF-TP/B3S as a function of a) constantviscosity and b) viscosity development as measured for temperatureprofile P2 during the laminate production step. . . . . . . . . . . . 97

5-15 Development of the DOI of CF-TP/C2000 as a function of a) con-stant viscosity and b) viscosity development as measured for tem-perature profile P2 during the laminate production step. . . . . . . 98

6-1 Stabilization reaction using sterically hindered phenols [176]; R1,R2,and R3 denote moiety. . . . . . . . . . . . . . . . . . . . . . . . . . 102

6-2 Stabilization reaction of aromatic amines [176]. . . . . . . . . . . . 1026-3 Effect of lubricants as a function of solubility in the host polymer;

redrawn from [192]. . . . . . . . . . . . . . . . . . . . . . . . . . . . 1076-4 Temperature profiles used for compounding additives to B3S and

C2000 in a twin-screw extruder. . . . . . . . . . . . . . . . . . . . . 1096-5 The OIT of B3S samples at 320 °C and C2000 samples at 340 °C,

both neat and modified with antioxidants. . . . . . . . . . . . . . . 1126-6 Mass loss of neat and single-modified a) B3S samples and b) C2000

samples during temperature profile P1. . . . . . . . . . . . . . . . . 1136-7 Complex viscosity of neat and modified a) B3S and b) C2000 sub-

jected to temperature profile P2 under oxidative and inert atmosphere.1146-8 Complex viscosity of neat and modified C2000 subjected to temper-

ature profile P2 under oxidative and inert atmosphere. . . . . . . . 1166-9 Complex viscosity of neat and multi-functionalized a) B3S and b)

C2000 subjected to temperature profile P2 under oxidative and inertatmosphere. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 118

6-10 Results from four-point bend testing a) in fiber direction and b)perpendicular to fiber direction of test panels produced from non-modified and multi-functionalized polymers at different dwell timesin a press. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119

6-11 MWD of as-received without processing, single-modified and multi-functionalized polymers after subjection to temperature profile P2in air of a) B3S and b) C2000; MWD of non-modified and multi-functionalized samples extracted from four-point bend test panels ofc) B3S and d) C2000. . . . . . . . . . . . . . . . . . . . . . . . . . . 120

xxix

6-12 Development of the DOI of CF-TP/B3S as a function of a) constantviscosity and viscosity development as measured for temperatureprofile P2 during the laminate production step of b) non-modifiedB3S and c) multi-functionalized B3S. . . . . . . . . . . . . . . . . . 121

6-13 Development of the DOI of CF-TP/C2000 as a function of a) con-stant viscosity and viscosity development as measured for temper-ature profile P2 during the laminate production step of b) non-modified C2000 and c) multi-functionalized C2000. . . . . . . . . . 122

7-1 Double-belt press with seven heated sections used to produce CF-TP/B3S tapes with different DOIi; modified from [201]. . . . . . . . 126

7-2 Differently impregnated CF-TP/B3S tapes with highlighted non-impregnated areas produced in a double-belt press with a) 2 m/minat 40 bar b) 4 m/min at 5 bar and c) 8 m/min at 5 bar. . . . . . . 127

7-3 Predicted final degree of impregnation (DOIf ) after press formingtapes with a DOIi of a),c) 80 % and b),d) 90 % at a dwell time of90 s and 300 s. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129

7-4 Micrographs of laminates produced with varying dwell times basedon differently impregnated tapes with highlighted non-impregnatedareas if applicable. . . . . . . . . . . . . . . . . . . . . . . . . . . . 130

7-5 a) Longitudinal flexural strength σf1 and b) longitudinal flexuralmodulus Ef1 of test panels made from CF-TP/B3S tapes with dif-ferent DOIi that were processed with varying dwell times in a staticpress or thermoformed (+TF). . . . . . . . . . . . . . . . . . . . . . 132

7-6 a) Transverse flexural strength σf2 and b) transverse flexural mod-ulus Ef2 of laminates made of tapes with different DOIi that wereprocessed with varying dwell times in a static press or thermoformed(+TF). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133

7-7 Correlation of costs to mechanical performance for a) σf1, b) σf2,c) Ef1 and d) Ef2 compared to reference values obtained from com-pletely impregnated tapes pressed for 1200 s. . . . . . . . . . . . . . 140

xxx

List of Tables1-1 Comparison of intermediate materials; based on [7]. . . . . . . . . . 6

2-1 Comparison of tensile strength σ, tensile modulus E and sizing typeof investigated carbon fibers. . . . . . . . . . . . . . . . . . . . . . . 14

2-2 Properties of investigated polyamides including density ρ, yield stressσyield (dried), tensile modulus E and Tm [54–58]. . . . . . . . . . . . 15

2-3 Investigated material combinations produced by the powder-coatingtechnique along with required lay-up to obtain desired test panelthickness for four-point bend and DCB testing. . . . . . . . . . . . 22

2-4 Summary of the mechanical properties for the investigated materialcombinations of differently sized carbon fibers and polyamides. . . . 33

3-1 Avrami exponent n, nucleation mode, crystal growth shape accord-ing to [76, 80]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

3-2 Kinetic parameters for the isothermal crystallization of neat B3S. . 463-3 Kinetic parameters for the isothermal crystallization of neat C2000. 473-4 Effect of cooling rate on crystallization of neat B3S. . . . . . . . . . 493-5 Effect of cooling rate on crystallization of neat C2000. . . . . . . . . 503-6 Overview of produced test panels along with matrix mass fraction

determined by acid digestion. . . . . . . . . . . . . . . . . . . . . . 53

4-1 Comparison of boundary conditions as present in Darcy’s Law andthermoplastic matrix flow through carbon fiber bed [121]. . . . . . . 58

4-2 Experimental design with three factors - time, temperature and pres-sure - for CF-TP/B3S and CF-TP/C2000. . . . . . . . . . . . . . . 70

4-3 Zero-shear viscosity data for B3S used as input parameter. . . . . . 744-4 Zero-shear viscosity data for C2000 used as input parameter. . . . . 744-5 Input parameters for model calibration. . . . . . . . . . . . . . . . . 78

6-1 Selected additives for thermal stabilization of B3S and C2000. . . . 1056-2 Selected additives to increase flowability. . . . . . . . . . . . . . . . 1086-3 Effectiveness of different lubricants on B3S and C2000, compared

across the first three process steps (powder-coating until laminateproduction) and the final step (thermoforming) of temperature pro-file P2 in an oxidative and inert atmosphere. . . . . . . . . . . . . . 116

6-4 Mass loss of neat and multi-functionalized (MF) B3S and C2000recorded during subjection to temperature profile P1 under an ox-idative atmosphere. . . . . . . . . . . . . . . . . . . . . . . . . . . . 117

xxxi

7-1 Double-belt press settings and yielded different DOIi of UD tapesdetermined from micrographs. . . . . . . . . . . . . . . . . . . . . . 127

7-2 Input parameters for the prediction of dwell times to completelyimpregnate partially impregnated UD tapes during press forming. . 128

7-3 Dwell times used to process differently impregnated tapes in a staticpress and via thermoforming along with the final DOIf of all testpanels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131

7-4 Assumptions made for cost analysis based on forecast for 2020. . . . 1357-5 Collected data used for the cost analysis. . . . . . . . . . . . . . . . 1387-6 Calculation of the machine hour rate for varying operating speeds

of a double-belt press. . . . . . . . . . . . . . . . . . . . . . . . . . 139

A-1 Three-factor Box-Behnken Design for CF-TP/B3S . . . . . . . . . . 169A-2 Three-factor Box-Behnken Design for CF-TP/C2000 . . . . . . . . 170A-3 FVC averaged over three samples of test panels produced from non-

modified and multi-functionalized B3S and C2000 reinforced by CF-TP carbon fibers . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171

A-4 FVC averaged over three samples of test panels produced with var-ious dwell times from differently impregnated tapes . . . . . . . . . 171

xxxii

1 Introduction

1.1 Motivation

Fiber reinforced composites have established in the aerospace industry over decadesdue to their specific properties. More than 50 % of the primary structures of the787 Dreamliner launched by Boeing are made of composites enabling fuel savingsup to 20 % [1]. Every reduction in weight and corresponding fuel savings are crucialduring the design due to the long service life of aircraft. The automotive industryhas a strong interest in weight reduction of cars as alternative propulsion conceptssuch as electric vehicles with heavy battery packs gain in importance, too. Byintroducing the i model series, BMW undertook a major step towards the massproduction of carbon fiber reinforced plastics (CFRP). There is a high demand forcost-efficient production of CFRP since these costs can comprise 50 % [2] of thetotal CFRP component cost.The main driver in reduction of process costs is automation. Automated fiber place-ment (AFP), automated tape laying (ATL) or press forming techniques (e.g. di-aphragm forming, thermoforming) are examples for highly advanced techniquesthat can reduce processing costs by their high level of automation involving repro-ducibility and accuracy.Another possibility for cost reduction lies in the use of suitable materials. By intro-ducing intermediate materials such as carbon fibers pre-impregnated with matrix(prepregs) ready for the use in automated processes, the time-consuming infiltra-tion or impregnation step of carbon fibers with matrix is mostly finished before theactual component production starts.In case of CFRP, two polymer groups are commercially used: Thermosets and ther-moplastics. In 2014, thermosets represented the most common (49 %) group usedfor CFRP amongst other matrix materials such as metals, ceramics, hybrids, car-bon or thermoplastics [3]. Established production processes and the initial chemicalconstitution of thermosets are responsible for their wide use in industry. Thermosetsinitially consist of two or three low-molecular constituents (resin, hardener and cat-alyst) with very low viscosities. During curing, these components start crosslinkingand form non-fusible polymers. The uncured thermosets facilitate wetting and in-filtration of thin reinforcing fibers due to the low viscosity.However, the eminent problems in processing epoxy-based thermosets were alreadyidentified in 1980: high brittleness, absorption of water/moisture and long manu-facturing times due to crosslinking [4]. In contrast, thermoplastic matrix systemsare already completely polymerized making them fusible, shapeable and more suit-able for repair and recycling processes. In addition, they are typically more ductile

1

2 Introduction

and tougher than thermosets. With having the impregnation process completed,thermoplastics can be processed with short forming and consolidation times of sev-eral minutes as no crosslinking reactions of up to several hours are required. Theshort processing times of thermoplastic composites can meet the low cycle timesthat govern the automotive industry.Besides economic advantages, thermoplastics meet also ecologic demands due totheir recycling potential. The last two decades were dominated by the increasingimportance of environmental impact. The European Community (EC) guideline [5]for reuse or disposal of 95% of the automobile weight came into effect in January2015. Other investments by the EC such as CleanSky worth several billions em-phasize the strong need for more environmentally-friendly air transportation.However, the long and branched molecular chains of polymerized thermoplasticsare the cause of melt viscosities that are 100-1000 times higher than for thermosetseven at processing temperatures well above the melting point [6]. The high melt vis-cosity of thermoplastics complicates wetting during processing. Especially carbonfibers with diameters of about 7 μm are difficult to impregnate. Thus, the produc-tion of intermediate products is costly due the time-consuming impregnation thatrequires high temperatures as well as pressures explaining the modest industrialuse of thermoplastics as matrix materials in continuously reinforced composites [3].In order to expand the use of thermoplastics as matrix materials, either the costsof the intermediate materials need to be reduced or the production of carbon fiberreinforced thermoplastics (CFRTP) needs to become more efficient. A combinedapproach may be most successful and will lead to an increased acceptance of ther-moplastic composites.

1.2 Thermoplastic composites

The component production of continuously carbon fiber reinforced thermoplastics(CFRTP) follows three essential steps: impregnation, consolidation and solidifica-tion. The impregnation step is usually carried out in a separate intermediate stepbefore the actual component production starts. In contrast to thermoset-basedCFRP, fibers and matrix are combined to obtain partially or completely impreg-nated intermediate materials that form the raw materials for CFRTP production.Thus, the time-consuming impregnation step is separated from the actual compo-nent production to make use of the potentially short forming and consolidationtimes of thermoplastics.The quantitatively most used intermediate materials with thermoplastic matrix areglass-mat reinforced thermoplastics (GMT) and long-fiber reinforced thermoplas-tics (LFT) [7] suitable for mass production of automotive thermoplastic compos-

Introduction 3

ite components with medium strength. For high-performance components made ofthermoplastic composites, intermediate materials with continuous carbon fibers arerequired. These intermediate materials are the focus of this thesis.

1.2.1 Intermediate materials

The production of intermediate materials to combine fibers and matrix follows thesame principal steps from impregnation, consolidation to solidification as knownfrom the CFRTP component manufacture. The carbon fiber reinforcement for inter-mediate materials covers the whole range of available fiber architectures includingspread unidirectional (UD) tows, woven fabrics, non-crimped fabrics (NCF), knit-ted or braided preforms.Amorphous or semi-crystalline thermoplastic matrix systems for intermediate ma-terials are selected from all application areas that are depicted in Figure 1-1.

PEEKPEKKPEK

PCPMMA PET

PSU

PA6

PPS

PPA

PBT POM

PE

PPHD-PE

Standard thermoplastics

Engineering thermoplastics

High-performance thermoplastics

Pric

e, p

erfo

rman

ce

PS

ABS

PVC

PEI

PA66SAN

Semi-crystalline

Amorphous

LD-PE

Figure 1-1 Classification of thermoplastics [8].

In general, thermoplastics are used in form of pellets, ground powder, suspension(with water) or solution (dissolved polymer) to produce intermediate materials.Amorphous thermoplastics such as polyetherimide (PEI) or polyether sulfone (PES)with high viscosities and no melting point are often processed as powder, solu-tion or suspension [7]. Semi-crystalline thermoplastics such as polyetheretherketone(PEEK), polyphenylene sulfide (PPS) or polyamides (PA6, PA66, PA10T) cannotbe dissolved properly due to their high chemical resistance against most solvents.

4 Introduction

They are usually processed via melt impregnation or powder-coating [9].Over the last decades, new types of thermoplastic intermediates appeared simulta-neously to the invention of new manufacturing techniques for thermoplastic com-posites [10]. Considering various forms of intermediates, the production may bedivided in the following principal process steps with regard to the thermoplasticmatrix [7]:

• Matrix application (suspension, solution, melt, film, powder)

• Heating (oven, infra-red source, calender, nozzle, double-belt press)

• Cooling/calibration (calender, double-belt press)

Due to this large variety, potential manufacturing routes for thermoplastic com-posites are presented in Figure 1-2, starting from reinforcing fibers and matrix toa final part.

Pre-forming

UD tapes, tows, textile prepregs

Powder-coated tows, commingled

yarns, FlT* bundles

Pre-consolidated sheetsWoven fabrics, braids

Manufacturing technique for shaping and formingAutoclave, vacuum consolidation, AFP, ATL, filament winding, pultrusion, press forming, thermoforming

PART*FIT: Fibres Imprégnées de Thermoplastique

Reinforcement(e.g. unidirectional (UD) spread tows,

woven fabrics, NCF, braids)

Thermoplastic matrix(e.g. powder, fiber, film, suspension)

Pre-impregnation

Techniques:• Film-stacking• Powder-coating• Fiber hybridization

Techniques:• Melt impregnation• Solution

impregnation

Figure 1-2 Overview of potential manufacturing routes originating from the production of in-termediates for thermoplastic composites, based on [7, 10, 11].

The production of intermediates can be divided into two principal techniques:pre-impregnation and pre-forming [10]. Pre-impregnation is commonly reached bymelt [12–18] or solution impregnation [19–21], and other exotic techniques such asimpregnation by using aqueous suspensions [22, 23].The pre-forming technique brings reinforcement and matrix together in a definedway without impregnation. Here, the impregnation takes place during part manu-

Introduction 5

facture by forming or shaping. Using such intermediates, reinforcement and matrixreveal a weak link due to the non-impregnated state leading to a high degree ofdrapeability that is maintained during lay-up. At the same time, a large flow dis-tance has to be covered in pre-formed intermediates. This makes the wetting aswell as the impregnation to the critical phase during component production [10].In the case of commingled yarns, the matrix is mixed in the form of thermoplasticfibers to reinforcing fibers producing a hybrid yarn that becomes rigid after consol-idation [24]. Another example for pre-formed intermediates is represented by thefilm-stacking method where thin polymer films and reinforcement layers (fabrics,NCF, spread tows etc.) [10] are consecutively stacked and consolidated in a double-belt press [25].Using the powder-coating method, polymer powder with particles in the range of5 to 200 μm is deposited on the reinforcement [10]. The powder deposition cantake place in an impregnation bath by using a fluidized bed or a fine suspension ofpowder particles in a liquid. The powder may also be directly applied by a needleroller or electrostatic deposition. To avoid loss of powder, the fabric, NCF or spreadtows coated with powder subsequently pass a heat system. By means of an oven,calender or heater, the powder is surface-fused ensuring sufficient adhesion to thereinforcement without impregnation [7].In 1983, Ganga [26] patented a special type of powder-coated intermediates: FibreImpregnée Thermoplastique (FIT). Here, the tows are powder-coated and enclosedby a flexible sheath made of preferably the same thermoplastic as the powder parti-cles. Thus, the powder maintains its position while these intermediates are furtherprocessed [27]. In general, FIT and commingled yarns are usually further processedto more complex preforms such as braids, knits, fabrics or three-dimensional pre-forms and consolidated afterwards. Film-stacked and powder-coated intermediatescan transform into pre-impregnated intermediates after passing a heat system un-der pressure e.g. in a double-belt press as indicated in Figure 1-2 [7, 28].The introduced intermediates vary with regard to the degree of impregnation (DOI)and the remaining flow path. Both characteristics determine the type of productiontechnology that can be used for the subsequent processing to a CFRTP compo-nent. Table 1-1 compares the introduced intermediate materials with regard to theinitial degree of impregnation (DOIi) before component production along with theremaining flow path, production rate, flexibility in relation to the material avail-ability and equipment costs. Here, the expression “hybrid” designates commingledyarns or FIT bundles. As not every thermoplastic polymer can be spun into a fiberor dissolved the flexibility of hybrid and solvent-impregnated intermediates is con-sidered to be low. Intermediates produced via melt impregnation show the highestDOIi with the lowest remaining flow path.

6 Introduction

Table 1-1 Comparison of intermediate materials; based on [7].

Intermediate Pre-formed Pre-impregnatedtypes Powder Film Hybrid Melt SolutionDOIi low medium medium high mediumRemaining flow path high medium medium low mediumProduction rate high medium medium medium mediumFlexibility medium medium low high lowEquipment cost medium high low-high medium high

1.2.2 Production methods

According to the classification of the previously introduced intermediates into pre-formed and pre-impregnated forms, different production methods are required. Pre-impregnated materials undergo three principle process steps: heating/melting aboveglass transition temperature Tg or melting temperature Tm, consolidation and cool-ing/solidification below Tg [10, 29] as depicted in Figure 1-3.

Pre

ssur

e [M

Pa]

Time [s]

Heating/Melting Consolidation Cooling/Solidification

Tem

pera

ture

[°C

]

Figure 1-3 Principle process steps during thermoplastic composite production along with thegoverning process parameters time, pressure and temperature; based on [29].

Pre-impregnated intermediates are typically completely impregnated and enablethe use of production techniques such as laser-assisted AFP or filament windingthat are aimed at in-situ consolidation. ATL and other automated placement tech-nologies such as FiberForge [30] were designed for quick and automated lay-up oflaminates without in-situ consolidation between plies. AFP, filament winding orthe FiberForge process technology are also applicable to pre-impregnated interme-

Introduction 7

diates with low to medium DOI along with subsequent consolidation in a press orin an autoclave. In addition, pre-impregnated intermediates are suitable for pul-trusion to profiles.Pre-formed intermediates such as commingled yarns or FIT generally involve an-other textile processing step to form a complex structure before the actual compo-nent production starts. Film-stacked or powder-coated reinforcements may be fur-ther processed in a calender, double-belt press or pultrusion equipment to obtaincompletely impregnated prepregs for subsequent use in automated production tech-nologies. In addition, they may be directly processed to pre-consolidated sheets in astatic press or in an autoclave under vacuum. By feeding alternately reinforcementlayers and polymer films into a double-belt press according to the film-stackingmethod, pre-consolidated sheets with multiple plies are produced [7, 10]. The di-rect production of multi-ply laminates in a double-belt press can also be achieved byusing powder-coated tows. In this way, a previous manufacture of pre-impregnatedmaterials can be eliminated and shorten the overall process chain.Pre-consolidated sheets, also referred to as multi-ply laminates or “organo-sheets”,are commonly shaped into complex structures by using different techniques ofpress forming or thermoforming. The actual forming and consolidation of the pre-consolidated laminates can occur in a matched-metal mold (matched-die molding),in a rigid female mold with flexible male mold (rubber forming, rubber-pad form-ing, hydroforming) or in between two membranes (double-diaphragm forming) [10].To heat or melt the pre-consolidated laminates, the used molds or press plates ofthe press equipment are heated and cooled during the press forming process. Inthe thermoforming process, pre-consolidated sheets are pre-heated externally byan infra-red source or heater and then transferred to the not-heated mold or pressplates. As mold and/or press plates are not heated and cooled, thermoforming al-lows short processing times involving high cooling rates.In more recent approaches such as SpriForm [31], injection molding and thermo-forming are combined into a single process. Here, multi-ply laminates are shapedin the cavity of an injection molding machine while more complex structures suchas ribs are injection-molded by using neat, short-fiber or long-fiber reinforced ther-moplastics. This process is also referred to as back-injection molding.Direct production techniques such as vacuum consolidation on a heated plate, astatic press or in an autoclave are independent of the used intermediate materialsand their DOIi. However, partially impregnated prepregs are not suitable for pro-cessing in AFP or filament winding as long as in-situ consolidation and completionof the impregnation progress is required. The pressure applied during AFP andfilament winding is insufficient and the processing time too short to completelyimpregnate and consolidate prepregs with remaining flow path [32].

8 Introduction

1.3 State-of-the-art

Within the last decades, several approaches were developed to overcome the dif-ficulty of impregnating fibrous reinforcements with high viscous thermoplastics.There are several paths to achieve a more efficient production of intermediates orthermoplastic composites which are summarized in the following:

• Optimization of the impregnation technology

• Integration of an online-impregnation device into processing technologies

• Reduction of the impregnation time by matrix modification and/or suitableprocess settings

By using pre-formed materials such as commingled yarns or powder-coated towsthe initial flow path is reduced and the actual impregnation occurs during partprocessing. By feeding commingled yarns into a pultrusion device, a good impreg-nation level is reported for high pultrusion speeds up to 10 m/min [33]. However,commingled yarns are constrained in width enabling the production of rods or pro-files but are less suitable for sheet production.Considering typically completely impregnated prepregs, the throughput can be in-creased by improving the impregnation technology. Marissen et al. [34] used specif-ically designed bars/pins which reduce the viscous drag and hence the relativespeed between spreader bars and fiber bundles to enhance the throughput of thethermoplastic pultrusion process. Weustink [35] refined the impregnation devicedeveloped by Marissen et al. by identifying some important key aspects for fixedand driven pins. The device itself including the pins shall be heated to achieve thelowest possible viscosity of the thermoplastic. Additionally, fiber bundles whichwere produced by using the impregnation device were directly fed into a filamentwinding machine.This combination of placement or winding technologies with online-impregnationwas also investigated by several other researchers. The operating efficiency of on-line melt-impregnated fiber bundles with direct processing in a filament windingdevice was found to be capable of competing to the use of pre-impregnated in-termediates [36]. In other approaches, powder-coated fiber bundles [37] or FITbundles [38] are fed into filament winding equipment.On a more comprehensive level, the overall process time to produce thermoplasticcomponents can be reduced by integrating different material forms into one pro-cess. Instead of consecutively proceeding processes, commingled yarns and powder-coated tows may be placed locally onto an injection-molded polymer part in anintegrated manufacturing cell [39]. This enables the production of complex partswith short cycle times.

Introduction 9

Another path to produce sufficiently impregnated intermediates or components ateconomically attractive production rates is the use of low viscous thermoplastics.This can be achieved by impregnation with thermoplastics previously dissolved ina suitable solvent [40] or by using prepolymers. The polymeric precursors enablethe use of production techniques generally designed for thermosets such as resintransfer molding (RTM), resin film infusion (RFI) or reaction injection molding(RIM) [41]. During thermoplastic RTM (T-RTM), the low molecular prepolymersfacilitate impregnation of fibrous reinforcement and form the thermoplastic matrixafter impregnation by in-situ polymerization [41, 42].The previous studies showed that focus is put on efficient production of either in-termediates or components. A more comprehensive approach lies in the productionof intermediate materials at enhanced production rates that are intended to fur-ther impregnate during the processing steps to obtain a final CFRTP component.Hayashi et al. [43] investigated the effects of different pressures and processing tem-peratures during thermoforming of completely and semi-impregnated (semi-preg)materials on their mechanical properties. The semi-pregs could be produced two orfour times faster than the fully impregnated prepreg. Mechanical properties of semi-pregs were found to be decreased compared to completely impregnated prepregs asit was not aimed at complete impregnation when processing intermediates to com-ponents. However, the DOIi of these semi-pregs was not evaluated and thus couldnot be correlated to the obtained mechanical properties. In addition, a relation be-tween mechanical properties to the monetary effect due to increased productivityduring semi-preg manufacture was not investigated.

1.4 Objectives and outline of the thesis

Currently, mostly completely impregnated intermediate materials are used for theproduction of thermoplastic composites. Considering a principal production processfor thermoplastic composites, the manufacturing steps require repeated heating ofthe polymer above Tg or Tm from intermediate production to tape or laminateconsolidation until forming to a component as presented in Figure 1-4.The overall objective of the present work is the production of cost-efficient interme-diates by increasing the operational throughput. Therefore, partially impregnatedtapes are developed that can be produced with enhanced production rates inde-pendent of the used manufacturing technology. The idea behind the use of partiallyimpregnated tapes lies in utilizing the repeated heating cycles during the produc-tion of thermoplastic components to complete impregnation. However, the remain-ing flow path must be adjusted such that gradual and complete impregnation isenabled throughout the subsequent process steps.

10 Introduction

Intermediate Products

[SG

L G

roup

]

Raw Materials

[SG

L G

roup

]

[IVW

]

Final componentPreforms, multi-ply laminates

[Cel

anes

e]

Heating Heating Heating

[SG

L G

roup

]

[SG

L G

roup

]

Figure 1-4 Principle production steps to manufacture thermoplastic composites; pictures pro-vided as courtesy by SGL Group and from Celanese [44] as well as from the Institutfür Verbundwerkstoffe (IVW) Kaiserslautern [45], as indicated.

The major research objectives of this work can be summarized as follows:

1. Characterization of suitable fiber-matrix combinations on microscopic andmacroscopic level as a function of the fiber sizing.

2. Study of the crystallization behavior to compare production processes thatinvolve different cooling rates.

3. Modeling the thermoplastic impregnation of fiber bundles to determine theinfluence of process parameters on the DOIi of intermediates as well as com-ponents.

4. Evaluation of viscosity changes induced by degradation reactions which de-velop during repetitive heating processes.

5. Prevention of viscosity changes due to degradation by thermal stabilizationand viscosity reduction by adding suitable lubricants.

6. Investigation of the gradual impregnation of partially impregnated tapesthroughout component production and evaluation of intermediate cost.

Based on these objectives, the structure of the present thesis was developed and isschematically shown in Figure 1-5.In Chapter 2, the compatibility of various carbon fiber sizings to several polyamidesis studied. The transverse four-point bend and the double-cantilever beam test areselected to characterize the adhesion between fiber and matrix. By using scanningelectron microscopy, the adhesion behavior is analyzed qualitatively. The nano-indentation technique serves the investigation of a three-dimensional phase that

Introduction 11

Chapter 6:Thermal stabilization and flow promotion

Viscosity reduction

Chapter 4:Impregnation model

• Exp. determination of DOI,• Process parameter prediction

Chapter 5:Degradation of polyamides

• Influence of thermal cycling• Processing window

Chapter ResultSelection of fibers and matrix

Chapter 2:Fiber-matrix compatibility

Chapter 3:Crystallization of polyamides

Comparability of production processes

Chapter 7:Gradual impregnation during production

Rapid production of partially impregnated tapes

Process-induced effects on impregnation

Enhancement of throughput during intermediate production

Reductions of Intermediate Cost

Aim

Objects of investigation

Figure 1-5 Schematic structure of the present thesis.

develops around the carbon fibers and is compared for various carbon fiber sizings.Based on the results from mechanical testing, the most suitable carbon fiber-matrixcombinations are selected for the subsequent investigations.To enable the comparison of differently processed intermediates and laminatesin terms of mechanical properties, the isothermal and non-isothermal crystalliza-tion behavior of the selected polyamide types is investigated in Chapter 3. Non-isothermal crystallization is studied on intermediate and composite level, too. Theintroduction of a new method to analyze the amount of crystallinity in fiber re-inforced thermoplastics enables full comparability. Thus, the dependency of themechanical properties on various cooling rates can be evaluated.In Chapter 4, the literature is reviewed to determine the phenomena that gov-ern the impregnation of carbon fibers with thermoplastics. Based on the literaturereview, a 1D model for the impregnation of powder-coated tows is derived. An ex-perimental technique is developed to determine the degree of impregnation (DOI)

12 Introduction

of powder-coated tows as well as composite parts. Impregnation experiments areconducted to evaluate the influence of process parameters on the DOI of the pro-cessed powder-coated tows.Long exposure of the used polymers to elevated temperatures during CFRTP pro-duction may lead to degradation reactions that involve considerable viscosity in-creases. Using powder-coated tows that are intended to completely impregnatethroughout processing, significant rises in viscosity may prevent the impregnationprogress. Thus, Chapter 5 concentrates on thermo-oxidative and thermal degrada-tion reactions of polyamides induced by repetitive heating processes from powder-coating tows to consolidation of tapes or laminates until thermoforming to a fi-nal component. Derived from this production process for thermoplastic compos-ites, the repetitive heating processes are represented by two temperature profiles.The temperature profiles are generated and analyzed by using differential scanningcalorimetry (DSC), thermogravimetric analysis (TGA), rheometry and gel perme-ation chromatography (GPC). By varying the dwell times of the single processsteps and the surrounding atmosphere, the initiation of degradation reactions isidentified to define a processing window for the investigated polyamides. In addi-tion, the influence of viscosity changes resulting from degradation reactions on theimpregnation time is evaluated by using the developed impregnation model.Several antioxidants are selected and mixed to both polyamides in Chapter 6 tolimit the extent of degradation reactions. Besides the suppression of degradationreactions, flow promotion is desired and realized by selecting different lubricants.The effectiveness of the antioxidant is evaluated by measuring the oxidation induc-tion time in the DSC. By subjecting polyamides modified with lubricants to thetemperature profiles derived from a CFRTP production process, the most efficientlubricant is chosen. Compounding the most efficient lubricant and antioxidant topolyamides, multi-functionalized polymers are produced. Potential interactions be-tween both additives are investigated by using DSC, TGA and rheometry. Then,powder-coated tows are produced from the multi-functionalized polymers and pro-cessed in a static press to determine influences on mechanical properties inducedby additives.Chapter 7 describes the gradual impregnation of partially-impregnated tapes dur-ing further processing steps. Tapes with different initial degree of impregnation(DOIi) are manufactured in a double-belt press with varying pressure and oper-ating speed. After processing in a static press with different dwell times and in athermoforming unit, the final degree of impregnation (DOIf ) and mechanical prop-erties are determined. Thus, the completion of impregnation is monitored and theeffect of the DOIi of tapes on mechanical properties is evaluated. The monetaryeffect arising from enhanced operating speeds in a double-belt press is estimatedby process cost analysis and related to the obtained mechanical properties.

2 Fiber-matrix compatibilityThe aim of this work lies in the cost-efficient production of intermediate materialsto produce high-performance CFRTP components. A sufficient fiber-matrix adhe-sion is crucial to the production of composites with high performance. Therefore,the compatibility of various polyamides to carbon fibers with different sizings isinvestigated within this chapter.The different combinations of carbon fibers and thermoplastics are produced frompowder-coated tows that are further processed in a static press. The various fiber-matrix combinations are evaluated by the transverse four-point bend and thedouble-cantilever beam test supported by fracture surface analysis via scanningelectron microscopy (SEM). In addition, nano-indentation tests are conducted toinvestigate the influence of different sizings on a three-dimensional interphase thatdevelops around the fibers. The fiber-matrix combinations that are found mostcompatible are used for the following investigations with regard to gradual impreg-nation during the production of CFRTP.

2.1 Investigated materials

2.1.1 Carbon fibers

Continuous carbon fibers are selected as reinforcement within this work because oftheir excellent mechanical properties. However, composites with high performancecan only be obtained when an ideal load transfer between fibers and matrix is en-abled [6]. Using a suitable sizing to match with e.g. engineering thermoplastics suchas polyamide 6 (PA6), a strong interface between fibers and matrix is establishedto enable load sharing among the fibers leading to excellent mechanical properties.The dominant share of carbon fiber reinforced epoxies in composite applications isresponsible for the large availability of carbon fibers compatible to epoxy matrixsystems. Typically, carbon fiber sizings consist of a mixture of adhesion promoters,film former, coupling agents, diluted solution of the matrix polymer or pre-polymersto enhance the adhesion between carbon fiber and host matrix [6, 7, 46]. The sizingneeds to fulfill multiple tasks such as improvement of handling and protection dur-ing textile processing as well as enhancement of the adhesion between fibers andmatrix [47]. In addition, a suitable chemical composition of the sizing can decreasethe contact angle by increasing the surface energy of the carbon fibers. This canimprove the wetting performance of the matrix enabling faster impregnation duringproduction of CFRTP [48].

13

14 Fiber-matrix compatibility

Sizings that are compatible to epoxy resins contain partially crosslinked polymers orresin components without crosslinking agent. Latter are designed to crosslink withthe host matrix during the composite production establishing a strong chemicalbond between fiber and matrix [49]. The use of epoxy-based coatings for thermo-plastic matrix materials is not only prevented by the lack of chemical compatibilitybut also by the decomposition temperature of epoxy groups at 250 °C [47]. Thistemperature represents the lower processing bound for most engineering and highperformance thermoplastics.Some researchers investigated surface-treated carbon fibers without finish/sizingand reported positive effects by removing the inert surface layer, forming oxygengroups and/or enhancing the surface roughness [46]. However, Drzal and Raghaven-dran [50] found that increasing oxygen groups on the fiber surface and enhancedsurface roughness are insufficient for good adhesion between fibers and matrix dueto a lack of strong covalent bonds. As thermoplastics are increasingly used as ma-trix materials for CFRP, there is a rising demand for new sizings. These have to bespecifically designed to chemically interact with thermoplastics and to withstandthe high processing temperatures during production.The properties of the used carbon fibers with different sizings supplied by SGLGroup are summarized in Table 2-1.

Table 2-1 Comparison of tensile strength σ, tensile modulus E and sizing type of investigatedcarbon fibers.

Carbon fiber σ [GPa] E [GPa] Sizing typeSIGRAFIL C T50-4.0/240-E100(CF-EPY) 4.0 240 epoxy-compatible

SIGRAFIL C T50-4.0/240-T140(CF-TP) 4.0 240 polyamide-compatible

Starting from precursor spinning to carbonization and surface treatment via an-odization, both carbon fiber types with 50,000 filaments (50k) were processed inan identical manner. They distinguish from each other by the application of differ-ent sizing after surface treatment. SIGRAFIL C T50-4.0/240-E100 carbon fibers(CF-EPY) are sized with an epoxy-compatible suspension and SIGRAFIL C T50-4.0/240-T140 carbon fibers (CF-TP) are coated with a sizing that mainly containspolyamides enabling compatibility to various polyamides [51].

2.1.2 Polyamides

Cost-efficient production of tapes goes along with the selection of cost-efficientthermoplastic matrix types with moderate mechanical properties. Therefore, sev-

Fiber-matrix compatibility 15

eral polyamide types were selected from the group of engineering thermoplas-tics (compare Figure 1-1) that are well-known for good mechanical performance,high chemical resistance and good damping behavior at moderate cost between2 to 10e per kg [52, 53]. Polyamides are macromolecular polymers that are linkedby acid amide groups (-CONH-). Between these links are hydrocarbons that can bealiphatic, aromatic or semi-aromatic. In general, there are two types of polyamides:polymers that are produced from two monomers (type AA/BB) and polymers pro-duced from bifunctional polymers such as ε-caprolactam (type AB). Table 2-2 sum-marizes the properties of the polyamides that were investigated within this work.

Table 2-2 Properties of investigated polyamides including density ρ, yield stress σyield (dried),tensile modulus E and Tm [54–58].

Polyamide type ρ [g/cm3] σyield [MPa] E [MPa] Tm [ °C]PA6 BASF B3Ssw 1.13 90 3500 220PA6 BASF B3L 1.10 70 2800 220PA6 BASF B40 1.13 79 2750 220PA10T/X Evonik C2000nc 1.11 62 2100 265

B3S, B40 and B3L belong to the class of aliphatic polyamides. The chemical struc-ture of typical aliphatic nylons is schematically shown in Figure 2-1. B3S is a⎡

⎣NH O

⎤⎦

n

Figure 2-1 Chemical structure of aliphatic PA6.

polyamide grade specifically developed for processes that require low viscosities.The affix ‘sw’ indicates that this polyamide grade is pigmented with carbon black.B3L is an aliphatic polyamide 6 (PA6) with elastomeric particles for tougheningand within the same viscosity range as B3S. The aliphatic B40 reveals a highermolecular weight than B3S and B3L (Mn = 33.000 g/mol). PA10T/X belongs tothe group of polphthalamides (PPA) and is a semi-aromatic co-polyamide poly-merized from 1,10 diaminodecan and terephthalic acid. The constitutional formulaof PA10T is depicted in Figure 2-2.

CO

CO

NH (CH2)10 NH

Figure 2-2 Chemical structure of semi-aromatic co-polyamide (PA10T).

The incorporation of terephthalic acid raises both Tg and Tm leading to an improvedheat resistance compared to aliphatic PA6. In addition, PPAs are characterized bydecreased moisture absorption in comparison to PA6 [59].

16 Fiber-matrix compatibility

2.2 Experimental methods

The experimental methods selected within this section aim at the evaluation ofthe fiber-matrix interface of the different material combinations. The transversebending test and the mode I interlaminar fracture toughness test were chosen asthey have been proven to be sufficiently sensitive to changes of the fiber-matrixinterface in CFRP [60, 61]. Besides macro-mechanical tests such as the transversefour-point bend test and the mode I interlaminar fracture toughness test, micro-mechanical nano-indentation tests were conducted. After mechanical testing, visualinspection of the tested specimens was carried out via scanning electron microscopy(SEM).

2.2.1 Four-point bend test

To determine the flexural properties of composites made of different material com-binations, the four-point bend test setup was selected due to the constant bendingmoment without introduction of shear between the loading noses. According to DINEN ISO 14125 B [62], the flexural strength and the flexural modulus transverse tothe fiber direction were determined. The used test setup is presented in Figure 2-3.The test specimens were 15 mm in width w and 60 mm in length l with a thickness t

of 2 mm. The support and loading noses were 2 mm in radius. The support span L

was adjusted to nominally 45 mm and the load span L′ to 15 mm. All tests werecarried out with a crosshead speed of 2.00 mm/min in a displacement-controlledmode on a Hegewald & Peschke 100 kN universal testing machine at room temper-ature. The mid-point deflection was determined by means of a video-extensometerwith telecentric lenses.

l

L

L’

t

Figure 2-3 Four-point bend test setup with support span L and a load span L’=L/3.

The flexural strength transverse to the fiber direction σf2 is determined by dividingthe maximum load Pmax times the support span L by the cross-sectional area ofthe specimen times thickness, according to Equation 2-1:

Fiber-matrix compatibility 17

σf2 = Pmax L

w t2 . (2-1)

The flexural modulus Ef2 is calculated as follows

Ef2 = 0.21 L3

w t3ΔP

Δs. (2-2)

Δs denotes the difference in deflection s′′ − s′ where s′ and s′′ correspond to thestrain of the outer fiber of ε′

f = 0.0005 and ε′′f = 0.0025. The difference in force ΔP

corresponds to Δs.

2.2.2 Double-cantilever beam test

The double-cantilever beam (DCB) test method based on ASTM D 5528 [63] wasapplied to determine the mode I interlaminar fracture toughness. For all DCB tests,a test fixture according to the side clamped beam (SCB) hinge system developed byRenart et al. [64] was used. In this case, the specimen is clamped from the top andthe bottom by a specifically designed grip zone. By using this test setup, adhesivebonding of loading blocks or piano hinges to the test specimen is eliminated [64].Figure 2-4 shows the specimen geometry (t = 3 mm, w = 25 mm, initial delamina-tion length a0 of 63 mm and l = 125 mm) and a DCB test specimen clamped bythe SCB fixture that is installed in a Hegewald & Peschke 100 kN universal testingmachine.

la0

t

Nonadhesive insertLoad line

a) b)

Figure 2-4 a) DCB test specimen with initial delamination length a0 from load line to end ofinsert and b) test specimen clamped to the SCB test fixture mounted to a Hegewald& Peschke 100 kN universal testing machine.

At first, a pre-crack with a delamination length between 3 to 5 mm was created byloading and unloading of each specimen with a crosshead speed of 3 mm/min. Sub-sequently, each specimen was reloaded until fracture at the same crosshead speedas used for the pre-crack. The two loading cycles were recorded with a camera

18 Fiber-matrix compatibility

at 4 frames per second. The opening mode I interlaminar fracture toughness GIc

and the propagation values GI resulting in the delamination resistance curve (Rcurve) were detected by visual observation and calculated by using the modifiedcompliance calibration (MCC) method. According to the MCC method the modeI interlaminar fracture toughness GI is calculated by taking into account the com-pliance C, force P and geometric dimensions such as specimen width w as well asthickness t:

GI = 3 P 2 C2/3

2 A1 w t. (2-3)

Here, the delamination length is first normalized by the specimen thickness (a/t)and plotted against the cube root of compliance including visually observed de-lamination onset values and propagation values. Subsequently, a least square plotis generated where the slope of the line represents A1 [63]. The opening mode Iinterlaminar fracture toughness GIc was determined based on the first visible crackextension from the reloading curve after the pre-crack has generated a defined crackfront.

2.2.3 Statistics

All results from mechanical testing were analyzed by means of a confidence intervalaccording to the test standard ISO 2602 [65]. A two-sided confidence interval wasused for the population mean m at a confidence level of 95 % as the followingequation shows:

x − t0.975√n

s < m < x + t0.975√n

, (2-4)

where x is the arithmetic mean of n results and s is the standard deviation. For n

independent measurements, t is calculated as follows:

t = x − m

s/√

n. (2-5)

The ratio t0.975√n

can be directly gathered from the Student’s t distribution given

in [65].

2.2.4 Nano-indentation

The nano-indentation technique is commonly employed to measure the modulusand hardness of various materials at the nanoscale. In this work, nano-indentationis used to measure the modulus and hardness of the region between matrix andcarbon fiber that is influenced by the use of different sizings. This region is often re-

Fiber-matrix compatibility 19

ferred to as interphase. The interphase is assumed to start somewhere in the fiberuntil the fiber properties convert from bulk to surface properties. At this pointthe actual interface between fiber and matrix is present. The interface connectsthe surface properties from fiber and matrix. Coming from the interface into thematrix, the polymer behavior changes due to the transition from surface to bulkproperties [66].In previous studies, nano-indentation was applied to determine the transverse prop-erties of carbon fibers [67], to characterize the interphase between carbon fibers toepoxy resins [68] or to vinyl ester matrix [69]. In addition, nano-indentation isused to determine the hardness of polymers induced by crystallinity changes afterprocessing with different cooling rates as present for AFP and autoclave consoli-dation [70].Within this work, the nano-indentation technique serves to investigate if an in-terphase establishes between different carbon fiber sizings and polyamides. Theinfluence of the investigated carbon fiber sizings can be determined by analyzingthe width and hardness of the developed interphase.The so-called Berkovich indenter with a tip in the shape of a three-sided pyramidwith a face angle of 65.27◦ [71] was used. The principle of this technique is presentedin Figure 2-5 along with a SEM image of the Berkovich indenter.

Carbon fiber

Matrix

7 μm

Nano-indenter

m

a b

Figure 2-5 a) Principle of nano-indentation on carbon fibers surrounded by matrix; b) SEMimage of Berkovich indenter tip [72].

The Berkovich indenter is ideally suited for polymers as a sharper point results fromthe three-sided pyramid than from other indenter geometries such as the four-sidedVickers tip. This enables a better control of the indentation process [71].Assuming a completely developed plastic area under the indenter tip, the meancontact pressure is defined by dividing the indenter load P by the project area A asEquation 2-6 presents. The mean contact pressure is also described as indentationhardness H [71]:

H = P

A= P

24.5 h2p

. (2-6)

20 Fiber-matrix compatibility

The projected area A is expressed by Equation 2-7 along with the plastic depth ofpenetration hp:

A = 3√

3 h2p tan2θ. (2-7)

The elastic modulus E of the examined materials is calculated from the slopeof the tangent to the initial unloading section of the load-displacement curve asfollows [71]:

E = 12

dP

dh

√π√A

P

24.5 h2p

. (2-8)

A Hysitron Triboindenter was used to conduct the nanoindentation tests. All nano-indentation tests and their evaluation were carried out at the Research Schoolof Physics and Engineering from the Australian National University (ANU) inCanberra in collaboration with Mr. Sherman Wang.

2.2.5 Scanning electron microscopy

SEM was used for visual inspection of the adhesion of differently sized carbon fibersto the selected polyamide grades. Before inspection via SEM, the fractured surfaceof four-point bend and DCB test specimens were sputtered using the sputter coaterBAL-TEC SCD 005 at 40 mA for 40 sec in an evacuated chamber that was flushedwith argon prior to the sputtering step. Thereby, a thin gold film of about 100 nmis created to enhance the electrical conductivity for SEM measurements. The usedSEM was a JSM 6060/6060LV from Jeol. The sputter coating and the image ac-quisition were conducted in collaboration with Mrs. Susanne Schnell-Witteczekfrom the IMETUM Institute of Medical Engineering of the Technical University ofMunich.

2.3 Sample preparation

2.3.1 Intermediate production

The powder-coating technique allows the investigation of different material combi-nations due to its high flexibility and is characterized by full comparability whensimilar particle size distributions of the powder are used. In addition, the DOIi

of 0 % allows the observation of the impregnation progress throughout processing.Therefore, the powder-coating method was selected as production technique forintermediates within this work.A laboratory prepreg line (LPL), located at SGL Carbon GmbH, was used toproduce the various material combinations. The schematic of the LPL used for

Fiber-matrix compatibility 21

production of the powder-coated tows is presented in Figure 2-6. For each of thematerial combination that are summarized in Table 2-3, nine carbon fiber towswere spread simultaneously to a width of nominally 200 mm and fed into the LPL.The powder was dried at 80 °C for at least 4 h before coating to avoid bubble for-

Spread carbon fiber tows

Infrared heat source

Heat

Powder-coated tows

Powder reservoir

Needle roller

Figure 2-6 Schematic of the prepreg line used to produce powder-coated tows on a laboratoryscale.

mation arising from heating up non-dried, hydrophilic powder. To begin with, thespread fibers were coated with 50 g/m2 of polymer powder on one side by meansof a needle roller. Before powder application, the exact width of the spread fiberswas measured and the amount of powder was adjusted if necessary to keep theareal weight constantly at 50 g/m2. The powder was fused by passing an infraredsource incorporated into the LPL. The coated tows were then solidified by runningthrough a pair of calender rolls and wound up. In a second run, the opposite side ofthe spread fibers was coated with 50 g/m2. Therefore, double-sided powder-coatedtows with an areal weight of 235 g/m2 were produced. The velocity of the labora-tory prepreg line was 5 m/min.Epoxy-sized carbon fibers could be spread to a width of 205 mm and CF-TP fibersto a width of 180 mm. The reduction in width resulted in an increase in thicknessof the powder-coated tows. Therefore, powder-coated tows with increased thicknessrequire a lay-up of less plies than tows that could be spread to 205 mm. Differentlay-ups arise from the difference in spreading behavior to generate a comparablethickness of 2±0.2 mm for four-point bend testing and 3±0.2 mm for DCB testingwithin the given tolerances, as presented in Table 2-3.

2.3.2 Test panel production

After manufacturing, the powder-coated tows were cut to the desired length andstacked in accordance to the lay-up stated in Table 2-3. Aluminum foils were usedto produce test panels with varying size as shown in Figure 2-7. During the lay-up

22 Fiber-matrix compatibility

Table 2-3 Investigated material combinations produced by the powder-coating technique alongwith required lay-up to obtain desired test panel thickness for four-point bend andDCB testing.

Material Combination Four-point bend DCBFiber-Sizing/Matrix Lay-up Thickness [mm] Lay-up Thickness [mm]CF-TP/B3S [012] 1.98±0.02 [018] 2.94±0.02CF-EPY/B3S [013] 2.14±0.17 [020] 3.16±0.05CF-TP/B3L [012] 2.05±0.01 [020] 2.86±0.03CF-TP/B40 [013] 2.08±0.01 [020] 3.15±0.05CF-TP/C2000 [011] 1.88±0.02 [020] 3.08±0.03CF-EPY/C2000 [012] 2.17±0.03 [022] 3.27±0.01

of the DCB test panels, a polyimide film (Upilex) with a thickness of 25 μm andcoated with high temperature resistant release agent (Frekote 55 NC) from bothsides was placed at the midplane of the laminate to generate the initial delaminationlength a0.

Aluminum foil

b) Processing in a static pressa) Stacking c) Produced test panel

Polyimide foil

Powder-coatedtows

Figure 2-7 a) Stacking, b) processing in a static press and c) produced test panel.

Subsequently, the stacks of powder-coated tows were dried at 80 °C in a vacuumoven for at least 4 h. After drying, each stack was transferred to a press WKP 1000sfrom Wickert heated to 260 °C (B3S, B3L, B40) and 300 °C (C2000). After thestacks have reached the processing temperature, a pressure of 10 bar was applied.Pressure and temperature were kept constant for 60 min (B3S, B3L, B40) and30 min (C2000) and cooled to 80 °C with a cooling rate of 20 °C/min under pressure.

2.3.3 Test specimen preparation

Test samples were cut to the required dimensions by using a water-cooled circulardiamond saw. The DCB specimens were coated with white water-based fluid and

Fiber-matrix compatibility 23

marked according to ASTM D 5528 [63] to enable observation of the crack propa-gation during testing. After cutting, the test specimens were dried under vacuumat 80 °C (PA6) and 110 °C (PA10T/X), respectively for at least 60 h to ensurecomplete water removal. Storage in hermetically sealed aluminum foils preventedmoisture absorption prior to testing.

2.3.4 Micrographs

Micrographs were prepared by extracting rectangular-shaped specimens from testpanels and potting them in transparent epoxy resin to enable visual inspection ofthe laminate quality. After curing the potting resin, the potted specimens wereground with silicon carbide sand papers with a grain size from 180 to 2400 byusing the polishing machine Struers TegraPol-21. Final surface finish was achievedby polishing the ground specimens with several suspensions and was completedwith a chemical suspension with particles smaller than 1 μm. Micrographs weretaken by using an Olympus BX41-M. Besides visual inspection, potted and polishedspecimens extracted from DCB test panels were used to carry out nano-indentationtests.

2.4 Results

To begin with, the results from macro-mechanical testing are analyzed with regardto the type of sizing. Then, focus is put on different aliphatic polyamide grades toevaluate the effect of increased matrix ductility on carbon fibers with thermoplastic-based sizing. Additionally, the influence of different carbon fiber sizings on thedevelopment of an interphase around the fibers is investigated by nano-indentation.

2.4.1 Influence of the sizing

Aliphatic polyamide

Focusing on the aliphatic polyamide grade B3S, the transverse flexural stress σf2

is plotted against the strain of the outer fiber εf2 for epoxy-sized carbon fibersCF-EPY/B3S and polyamide-based sized fibers CF-TP/B3S in Figure 2-8. Thepolyamide-based sizing seems to promote brittle failure at high stresses. In con-trast, the epoxy-based sizing leads to an early failure at considerably lower stresses.Plotting the opening mode I interlaminar fracture toughness GIc and the propaga-tion values GI against the crack length yields the R curve that is compared for B3Sspecimens with EPY and TP sizing in Figure 2-9. The GIc of B3S combined with

24 Fiber-matrix compatibility

0 1 2 3 4 50

20

40

60

80

100

120

0 1 2 3 4 50

20

40

60

80

100

120σ f2

[MP

a]

εf2 [%]

CF-EPY/B3S_01CF-EPY/B3S_02CF-EPY/B3S_05CF-EPY/B3S_07CF-EPY/B3S_08

a

σ f2[M

Pa]

εf2 [%]

CF-TP/B3S_01CF-TP/B3S_03CF-TP/B3S_04CF-TP/B3S_06CF-TP/B3S_07

b

Figure 2-8 a) Stress-strain curves of CF-EPY/B3S and b) CF-TP/B3S.

epoxy- and polyamide-based sizing are very similar. However, a significant increasein the propagation values GI is observed for CF-TP/B3S that is attributed to fiberbridging and stick-slip behavior noticed during testing.

50 60 70 80 90 100 1100

750

1500

2250

3000

3750

50 60 70 80 90 100 1100

750

1500

2250

3000

3750

GI[J

/m2 ]

crack length [mm]

CF-EPY/B3S_01CF-EPY/B3S_04CF-EPY/B3S_07CF-EPY/B3S_10

a

GI[J

/m2 ]

crack length [mm]

CF-TP/B3S_01CF-TP/B3S_02CF-TP/B3S_07CF-TP/B3S_08

b

Figure 2-9 R curves for a) CF-EPY/B3S and b) CF-TP/B3S.

Figure 2-10 compares the mean transverse flexural strength σf2 along with the meanopening mode I interlaminar fracture toughness GIc of specimens with CF-TP andCF-EPY fibers impregnated with B3S. Comparing the results from mechanical test-ing, the polyamide sizing strongly influences transverse flexural properties. σf2 isincreased by 152 % when polyamide-sized fibers are used and lies above the matrixyield stress indicating a strong adhesion between fibers and B3S. The effect of theTP sizing has a less pronounced effect on GIc revealing an improvement by 11 %.Inspection of the fracture surfaces of tested DCB specimens by SEM (Figure 2-11) reveals that B3S peels off of the epoxy-coated carbon fibers implying a weakadhesion between fibers and matrix. In contrast, B3S almost covers the completecarbon fiber with polyamide-based sizing and shows cohesive failure. In addition,

Fiber-matrix compatibility 25

CF-EPY/B3S CF-TP/B3S0

20

40

60

80

100

120M

ean

σ f2[M

Pa]

a

CF-EPY/B3S CF-TP/B3S0

400

800

1200

1600

2000

Mea

nG

Ic[J

/m2 ]

b

Figure 2-10 a) Mean transverse flexural strength σf2 and b) mean mode I interlaminar fracturetoughness GIc of CF-EPY/B3S in comparison to CF-TP/B3S.

the matrix area around polyamide-sized carbon fibers exhibits a higher degree ofdeformation compared to epoxy-sized fibers.

a) CF-EPY/B3S b) CF-TP/B3S

Figure 2-11 Fracture surface analysis of tested DCB specimens made of a) CF-EPY/B3S incomparison to b) CF-TP/B3S.

Co-polyamide

Although the thermoplastic-based sizing originally has been designed to chemicallyinteract with polyamides, its compatibility to the co-polyamide PA10T/X (C2000)was investigated. The chemical structure of PA10T/X reveals large aliphatic re-gions, segmented by terephthalic acid which is why a sufficient adhesion to CF-TPfibers is expected. Parts of the following experimental work have previously beenpublished in [73].In Figure 2-12, the transverse flexural stress is plotted against the strain of the outerfiber for CF-EPY/C2000 and CF-TP/C2000. In contrast to aliphatic polyamidessuch as B3S, the epoxy-based sizing in combination with C2000 leads to higherstresses at higher strains compared to the polyamide sizing. The flexural stiffness

26 Fiber-matrix compatibility

of the specimens is comparable for both carbon fiber sizings.

0.0 0.2 0.4 0.6 0.8 1.00

20

40

60

80

100

0.0 0.2 0.4 0.6 0.8 1.00

20

40

60

80

100

σ f2[M

Pa]

εf2 [%]

CF-EPY/C2000_01CF-EPY/C2000_02CF-EPY/C2000_03CF-EPY/C2000_04CF-EPY/C2000_12

a

σ f2[M

Pa]

εf2 [%]

CF-TP/C2000_02CF-TP/C2000_03CF-TP/C2000_09CF-TP/C2000_10CF-TP/C2000_13

b

Figure 2-12 Stress-strain curves for a) CF-EPY/C2000 and b) CF-TP/C2000.

The R curves of the C2000 with EPY and TP sizing are presented in Figure 2-13.Lower scatter but also reduced GIc is found for CF-EPY/C2000 specimens com-

40 50 60 70 80 90 100 1100

300

600

900

1200

1500

40 50 60 70 80 90 100 1100

300

600

900

1200

1500

GI[J

/m2 ]

crack length [mm]

CF-EPY/C2000_04CF-EPY/C2000_05CF-EPY/C2000_06CF-EPY/C2000_07CF-EPY/C2000_08

a

GI[J

/m2 ]

crack length [mm]

CF-TP/C2000_02CF-TP/C2000_03CF-TP/C2000_05CF-TP/C2000_07CF-TP/C2000_08

b

Figure 2-13 a) R curves of CF-EPY/C2000 and b) CF-TP/C2000.

pared to CF-TP/C2000 specimens. After crack initiation, the propagation valuesGI of CF-TP/C2000 specimens further increase considerably, compared to CF-EPY/C2000 specimens. This is assumed to arise from fiber bridging noticed duringtesting.Figure 2-14 compares the mean transverse flexural strength σf2 as well as the meanopening mode I interlaminar fracture toughness GIc of specimens with C2000 rein-forced by CF-EPY and CF-TP fibers. In contrast to the aliphatic B3S, the trans-verse flexural strength of specimens with CF-TP fibers was found to be decreasedby 23 % in comparison to CF-EPY/C2000. However, the GIc was increased by 44 %for C2000 reinforced by polyamide-sized fibers.

Fiber-matrix compatibility 27

CF-EPY/C2000

CF-TP/C2000

0

10

20

30

40

50b

Mea

nσ f2

[MP

a]a

CF-EPY/C2000

CF-TP/C2000

0

50

100

150

200

250

300

350

Mea

nG

IC[J

/m2 ]

Figure 2-14 a) Mean transverse flexural strength σf2 and b) mean mode I interlaminar fracturetoughness GIc of CF-EPY/C2000 in comparison to CF-TP/C2000.

In Figure 2-15, the fractured surfaces of flexural test specimens with C2000 as ma-trix are compared with respect to carbon fibers with epoxy- and polyamide-basedsizing.

a) CF-EPY/C2000 b) CF-TP/C2000

Figure 2-15 Fracture surface analysis of tested flexural specimens made of a) CF-EPY/C2000compared to b) CF-TP/C2000.

The fractured surface of representative CF-TP/C2000 specimens show a patternof several broken fiber bundles that may cause the drop to residual stress levelsas observed during testing (see Figure 2-12). When several fiber bundles break atmaximum stress non-broken fiber bundles remain to withstand the bending load.In contrast, a more homogeneous fracture surface is found for CF-EPY/C2000.The epoxy-based sized fibers were spread to a larger width (increased by 14 %)than CF-TP fibers. Cross-sections of the four-point bend test panels presented inFigure 2-16 show insufficiently spread tows leading to fiber agglomerations in caseof polyamide-sized fibers. These areas with high fiber agglomerations are assumedto induce premature failure of the specimens during four-point bend testing. In thiscase, the four-point bend test gives more information about the laminate quality of

28 Fiber-matrix compatibility

a) CF-EPY/C2000 b) CF-TP/C2000

Fiber agglomerations

Homogeneous ply distribution

Figure 2-16 Micrographs of four-point bend test panels made of a) CF-EPY/C2000 and to b)CF-TP/C2000.

CF-TP/C2000 than about the level of adhesion. Thus, the adhesion of differentlysized carbon fibers was further analyzed by comparing the fracture surface of rep-resentative DCB test specimens presented in Figure 2-17.

a) CF-EPY/C2000 b) CF-TP/C2000

Figure 2-17 Fracture surface analysis of tested DCB specimens made of a) CF-EPY/C2000compared to b) CF-TP/C2000.

Improved adhesion between polyamide-sized fibers and C2000 is observed whichis indicated by residual amount of matrix on the surface while epoxy-sized fibersreveal a comparatively clean surface.

2.4.2 Influence of matrix ductility

Besides the influence of the sizing on the interface to the matrix, the compatibilityof the polyamide-sized carbon fibers to toughened (B3L) and high-molecular weight(B40) polyamides is investigated by using the transverse four-point bend as well asthe DCB test method. The large improvement of the transverse flexural strengthfor CF-TP/B3S in comparison to CF-EPY/B3S let expect a similar increase infracture toughness. However, the increase in fracture toughness was more moder-ate when using the polyamide-sized fiber than found for the flexural strength. The

Fiber-matrix compatibility 29

more ductile polyamides B3L and B40 serve to investigate if the strong interfaceestablished between B3S and polyamide-sized fiber can be further utilized as theymay be capable of improved load transfer.In Figure 2-18, the transverse flexural stress is plotted against the strain of theouter fiber for CF-TP/B3L and CF-TP/B40. Similar to B3S, CF-TP/B40 shows

0 1 2 3 4 50

20

40

60

80

100

120

0 1 2 3 4 50

20

40

60

80

100

120

σ f2[M

Pa]

εf2 [%]

CF-TP/B3L_01CF-TP/B3L_02CF-TP/B3L_08CF-TP/B3L_10

a

σ f2[M

Pa]

εf2 [%]

CF-TP/B40_03CF-TP/B40_06CF-TP/B40_09CF-TP/B40_11CF-TP/B40_12

b

Figure 2-18 Stress-strain curves of a) CF-TP/B3L in comparison to b) CF-TP/B40.

brittle failure at low strains. B3L toughened with elastomeric particles reveals amore ductile behavior at lower strains and stresses.The R curves are plotted and compared for CF-TP/B3L and CF-TP/B40 in Fig-ure 2-19. The brittle failure of CF-TP/B40 led to a strong stick-slip behavior dur-

50 60 70 80 90 100 1100

750

1500

2250

3000

3750

50 60 70 80 90 100 1100

750

1500

2250

3000

3750

GI[J

/m2 ]

crack length [mm]

CF-TP/B3L_01CF-TP/B3L_04CF-TP/B3L_05CF-TP/B3L_06

a

GI[J

/m2 ]

crack length [mm]

CF-TP/B40_01CF-TP/B40_06CF-TP/B40_08CF-TP/B40_10

b

Figure 2-19 R curves for a) CF-TP/B3L and b) CF-TP/B40.

ing DCB testing so that the pre-crack could not be initiated properly. Hence, theopening mode I interlaminar fracture toughness values are determined after thefirst rapid crack advancement that may lead to an overestimation of GIc.In Figure 2-20 the mean transverse flexural strength σf2 along with the mean mode

30 Fiber-matrix compatibility

I interlaminar fracture toughness GIc of specimens made of B3S, B3L and B40 re-inforced with CF-TP fibers are compared. The transverse flexural properties arenot affected by increasing matrix ductility. Using B3L even lowers the transverseflexural strength. However, a considerable improvement of GIc is yielded when us-ing the toughened B3L (72 %) and high-molecular weight B40 (55 %) comparedto B3S. This indicates that a strong and brittle interface is established betweenpolyamide-sized fibers to aliphatic PA6. As soon as the matrix provides a higherductility the strong link between fibers and matrix can be exploited to a greaterextent and leads to improved delamination resistance.

CF-TP/B3S

CF-TP/B3L

CF-TP/B40

0

20

40

60

80

100

120

Mea

nσ f2

[MP

a]

a

CF-TP/B3S

CF-TP/B3L

CF-TP/B40

0

400

800

1200

1600

2000M

ean

GIc

[J/m

2 ]b

Figure 2-20 a) Mean transverse flexural strength σf2 and b) mean mode I interlaminar fracturetoughness GIc of CF-TP/B3L and CF-TP/B40 in comparison to CF-TP/B3S.

Fracture analysis (see Figure 2-21) of the tested DCB specimen reveals increasingductility as well as cohesive failure when B3L or B40 are used compared to B3S.The large extent of plastic deformation around CF-TP fibers induced by B3L andB40 may be an indication for improved load transfer from the fibers across thestrong interface to the matrix.

a) CF-TP/B3S b) CF-TP/B3L c) CF-TP/B40

Figure 2-21 Fracture surface analysis of tested DCB specimens made of a) CF-TP/B3S, b)CF-TP/B3L and c) CF-TP/B40.

Fiber-matrix compatibility 31

2.4.3 Development of an interphase

By using carbon fibers with epoxy- as well as polyamide-based sizing surrounded byB3S, nano-indentation tests were conducted to investigate if and how an interphaseis developed between fiber and matrix. Pretrials were conducted to differentiatebetween the hardness of matrix and fibers. For this study, 30 to 32 indents at adistance of 100 nm were carried out on the matrix and the carbon fiber only asdepicted in Figure 2-22.

0 50 100 150 200 2500

50

100

150

200

250

300

350

0 50 100 150 200 2500

50

100

150

200

250

300

350CF-TP #1-#33

Load

[μN

]

Depth [nm]

a B3S #1-#31

Load

[μN

]

Depth [nm]

b

Figure 2-22 Nano-indentation on a) CF-TP only and b) B3S only.

A mean hardness of 6.49 GPa±0.29 is reached for the carbon fiber. When indentingon B3S only, a mean hardness of 0.25 GPa ± 0.05 is reached. Therefore, the fol-lowing classification is made: Indents yielding hardness values of more than 6 GPaare attributed to the carbon fiber and indents that give a hardness of less than0.3 GPa are assigned to the polyamide matrix. Observing the produced indents viamicroscope it was detected that the selected distance of 100 nm between the singleindents was sufficiently large to avoid overlapping of plastically deformed indentzones.In Figure 2-23 the hardness development in the area around CF-EPY and CF-TPfiber (embedded in B3S) is compared when indenting from the matrix to the fiber.For CF-EPY fibers no indents on B3S were found according to the previous hard-ness classification. The decrease in hardness after 500 nm (5 indents) is attributedto less precise indents in the gap (highlighted in Figure 2-24) that has formed be-tween CF-EPY and B3S observed at large magnifications under the SEM by usinga Hitachi 4300 SE/N FESEM.The CF-TP/B3S sample (Figure 2-23b) shows a region with increased hardness(indent #7-#11) of an approximate width of 500 nm between fiber and matrix. Forboth CF-EPY/B3S and CF-TP/B3S, hardened regions were found that are harderthan the neat matrix and softer than carbon fibers. These hardened regions are

32 Fiber-matrix compatibility

Carbon fiber

Hardened region

Carbon fiber

Hardened region

PA6

a b

0 5 10 15 20 25 300123456789

10

Har

dnes

s [G

Pa]

Indent

CF-EPY

0 5 10 15 20 25 300123456789

10

Har

dnes

s [G

Pa]

Indent

CF-TP

Figure 2-23 Nano-indentation on carbon fibers coated with a) epoxy-based and b)thermoplastic-based sizing, surrounded by B3S; blue arrow indicates indenting di-rection and covered area.

5 μm 5 μm

a) CF-EPY/B3S b) CF-TP/B3S

Figure 2-24 Close-up view of a) CF-EPY fibers with highlighted gap between fiber and matrix,b) CF-TP fibers, surrounded by B3S, under the SEM.

considered as interphase between fiber and matrix. The modulus of the hardenedregion is higher for CF-EPY/B3S than for the interphase of CF-TP/B3S. A pos-sible explanation is an influence of the epoxy-based sizing on the crystallizationbehavior in the area around the fibers. Epoxy-based sizing are chemically differentto the host matrix and can act as heterogeneous nuclei as discussed in Chapter 3.The gap between carbon fiber and matrix which was observed for CF-EPY/B3Smay cause the low transverse flexural strength and reasonable mode I interlaminarfracture toughness as slipping of the fibers is facilitated. As the four-point bendtests reveal, a strong interface establishes between CF-TP fibers and B3S indicatedby considerably increased transverse strength higher than the matrix yield stress.As no such gaps were found between CF-TP fibers and B3S as for CF-EPY/B3Sthe fibers may not slip as they are rigidly bond to the matrix. Potential locationsfor failure during four-point bend testing of CF-TP/B3S may be between matrixand interphase. When more ductile PA6 grades are used, the rigid bond between

Fiber-matrix compatibility 33

fibers and matrix can be compensated by increased ductility which leads to animproved load transfer.

2.5 Selection of compatible material combinations

Table 2-4 summarizes the results from four-point bend and DCB testing for theinvestigated material combinations that were produced by the powder-coating tech-nique with subsequent processing in a static press.

Table 2-4 Summary of the mechanical properties for the investigated material combinations ofdifferently sized carbon fibers and polyamides.

σf2 [MPa] GIc [J/m2]CF-EPY/B3S 42.36±0.99 1069.70±93.01CF-TP/B3S 106.92±3.40 1190.90±89.90CF-EPY/C2000 41.78±1.88 187.06±24.00CF-TP/C2000 32.37±1.54 269.78±55.59CF-TP/B3L 83.01±1.71 2047.21±79.61CF-TP/B40 110.65±1.95 1844.34±94.93

By selecting carbon fibers with polyamide-based sizing in combination with thealiphatic PA6 (B3S), a considerable increase by 152 % in transverse flexural strengthand a moderate gain of 11 % in interlaminar fracture toughness were reached com-pared to carbon fibers with epoxy sizing. Consistent with the test results, thefracture surface analysis by using SEM revealed strong adhesion between CF-TPand B3S. A more detailed inspection of CF-EPY/B3S and CF-TP/B3S by nano-indentation revealed the establishment of an interphase with increased hardness.When the co-polyamide C2000 was used as the matrix, a decrease by 23 % in trans-verse flexural strength and an increase by 44 % in interlaminar fracture toughnesswas obtained. The results from four-point bend testing are attributed to signifi-cantly different laminate qualities arising from differences in the spreading behav-ior of epoxy- and polyamide-sized fibers during powder-coating. Advancements inthe chemical composition of the polyamide-based sizing is believed to improve thespreadability. The results from DCB testing as well as the fracture analysis underthe SEM show an improved adhesion between CF-TP fibers and C2000 comparedto epoxy-sized fibers.Both B3L and B40 were investigated to optimize fracture toughness properties. Themoderate gain in interlaminar fracture toughness of CF-TP/B3S in comparison toCF-EPY/B3S could be further increased by 72 % with a toughened PA6 (B3L) andby 55 % using a PA6 with higher molecular weight (B40). However, the transverse

34 Fiber-matrix compatibility

flexural strength was found to be decreased by 23 % for CF-TP/B3L compared toCF-TP/B3S. This is attributed to the incorporated elastomeric particles present inB3L. Using B40 as matrix yielded overall improved flexural and toughness proper-ties compared to CF-TP/B3S but revealed also highly brittle failure behavior. Inaddition, B40 is characterized by a considerably higher melt viscosity compared toB3S as shown in Chapter 4 that is detrimental to gradual impregnation. However,both B3L and B40 may be used when fracture toughness represents the crucialdesign criterion.Supported by most results from mechanical testing and the fracture analysis viaSEM, the material combinations CF-TP/B3S and CF-TP/C2000 are selected forfurther investigations. In general, C2000 shows lower mechanical properties thanB3S but its chemical structure leading to increased heat resistance and low moistureabsorption expands the field of application.

3 Crystallization of polyamidesFor the investigation of gradual impregnation, press forming and thermoforming areconsidered as manufacturing processes. Both processes are characterized by signif-icantly different cooling rates. The selected thermoplastics PA6 (B3S) and the co-polyamide PA10T/X (C2000) are semi-crystalline that develop different amounts ofcrystallized fractions in dependance of the cooling rate. Hence, the non-isothermalcrystallization kinetics needs to be understood to enable comparability of mechan-ical properties across different processes. In addition, the isothermal crystallizationkinetics is investigated to identify the temperature range where the crystallizationproceeds most rapidly and comprehensively. During processing it may be cooled tothese temperatures to maximize crystallization and reduce residual stresses.The isothermal and non-isothermal crystallization kinetics are investigated for bothneat B3S and C2000. Distinctive cooling rates are selected afterwards to conductnon-isothermal experiments on B3S and C2000 combined with CF-TP and CF-EPYto examine the influence of differently sized carbon fibers on the crystallization be-havior. In addition, mechanical properties determined from test panels that wereproduced with the selected cooling rates are correlated with the process-inducedcrystallinity.

3.1 Theory

3.1.1 Crystallization and nucleation

When semi-crystalline thermoplastics are cooled from the molten state, crystal-lization proceeds in two basic phases: primary and secondary crystallization. Theprimary crystallization starts with nucleation and continues with crystal growth.The crystal form depends on the polymer type and the degree of undercooling.When freely formed crystals impinge and limit their further growth, the secondphase, the secondary crystallization, is initiated. In this phase, crystals cannotextend anymore and grow into spaces between other crystals so that the crystalstructure is perfected.The two phases run time-delayed. After a degree of crystallinity of approximately50 % [74], the developed crystals constrain each other and secondary crystallizationstarts. In general, the crystallization rate depends on the nucleation rate and thegrowth rate. The time-delayed course of nucleation and crystal growth is presentedas a function of temperature in Figure 3-1. The nucleation and growth rate de-pend on the cooling rate from melt and crystallization temperature. The larger the

35

36 Crystallization of polyamides

undercooling, the more nuclei are formed. However, high undercooling goes alongwith low crystallization temperatures that restrict the growth rate. Hence, manybut small nuclei are formed at large undercooling [75].

Temperature [°C]

Gro

wth

rate

Undercooled melt

TmTg Tc

Solid MeltCrystallization

Growth rate Nucleation

rate

Figure 3-1 Time-delayed development of nucleation and growth rate during crystallization as afunction of temperature, redrawn from [75].

At the beginning of primary crystallization, nucleation occurs heterogeneously orhomogeneously. Both homogeneous and heterogeneous nucleation refer to the wayhow nucleation is induced. Thermal and athermal nucleation describe at whattime nucleation takes place. Homogenous nucleation is thermal at all times and oc-curs throughout crystallization. During athermal nucleation, simultaneous crystalgrowth occurs at all nuclei. Heterogeneous nucleation is mostly athermal but mightalso be thermal [76].Homogeneous nucleation is believed to take place by the formation of locally alignedchains (“embryos”) due to thermal fluctuations. As soon as the embryo exceeds acritical size a nucleus of a growing crystal is formed [77]. The regularly arrangedcarboxamide groups (-CONH-) of aliphatic polyamides are able to form hydrogenbonds with other carboxamides that benefits nucleation. In addition, the crystal-lization temperature Tc of PA6 is well above Tg that leaves the polymer enoughtime to reach a state required for crystallization as the molecular chains remainmobile. For co-polyamides such as PA10T/X the ability to crystallize is weakenedby disruptive non-aliphatic structures as the terepthalic acid. The terepthalic acidincreases Tg and lowers Tc leaving less time for crystallization [75].Polyamides mainly crystallize heterogeneously at boundary surfaces of materialsdiffering from the polymer. These boundary surfaces can represent additives, im-purities, reinforcement (e.g. carbon fibers), pigments (e.g. carbon black), specialnucleating agents or cavity walls and act as crystal nuclei [75]. Leftovers fromcatalysis, synthesis or first processing represent the most typical foreign nuclei incommercial polymers [74].

Crystallization of polyamides 37

Heterogeneous crystallization is characterized by an immediate start of the crys-tallization at foreign nuclei as soon as the crystallization temperature is reached.Thus, the amount of crystal nuclei becomes independent of the polymer structureand cooling rate. In case of completely heterogeneous nucleation no further nucleiare formed from melt.The entire melt volume crystallizes to a solid compound throughout the first phaseof crystallization. At the beginning of secondary crystallization, crystals grow lin-early away from the nucleus into the three-dimensional space in shape of lamellae.This growth pattern still leaves gaps between the lamellae that are continuouslyfilled by splitting and branching out of the lamellae forming superstructures [74].Spherulithes represent the superstructures of polyamides [75]. As a consequence ofheterogeneous nucleation, many lamellar crystals grow at the same time leadingto a quasi-spherical symmetry. On the contrary, homogeneously nucleated crystalsare characterized by initial anisotropy [77].

3.1.2 Crystallization kinetics

Isothermal Crystallization

The primary crystallization is commonly described isothermally as a function oftime by the general Avrami equation [78, 79], also often referred to as the Kolmogorov-Johnson-Mehl-Avrami equation [74]:

1 − Vc(t) = exp(−ktn) (3-1)

Vc(t) denotes the volumetric fraction of crystalline material and hence the relativedegree of crystallinity as a function of time. n is the Avrami exponent and k thecrystallization rate constant. Equation 3-1 is based on the assumptions of a constantnucleation rate and a constant, linear growth [74]. The Avrami exponent n consistsof two terms:

n = nd + nn (3-2)

The term nd describes the dimension of the growing crystals and integers between1 and 3 refer to one-, two- and three-dimensional growth. nn specifies the time-dependance of the crystallization and ranges between 0 and 1. nn = 0 states thatcrystallization starts immediately at all nuclei (heterogeneous nucleation) and nn =1 means that crystal nuclei are formed from melt (homogeneous nucleation). Hybridforms with values between 0 and 1 may also occur. Hence, the Avrami index rangesbetween 0 and 4. Table 3-1 gives an overview of theoretical Avrami exponents alongwith the corresponding nucleation mode and crystal form.

38 Crystallization of polyamides

Table 3-1 Avrami exponent n, nucleation mode, crystal growth shape according to [76, 80].

Dimension of Avrami exponent ncrystal growth Athermal nucleation Thermal nucleation1D (fibrillar) 1 22D (lamellar) 2 33D (spherulithic) 3 4

Non-isothermal crystallization

Ozawa [81] extended the Avrami equation and further developed the mathematicalexpression by Evans [82], originally derived for cast metals, to non-isothermal mea-surements under the assumption that the polymer is heated or cooled at a constantrate. The Ozawa equation expresses the relative degree of crystallinity X(T ) as afunction of temperature and cooling rate [83]:

1 − X(T ) = exp(−K(T ) 1Cm

) (3-3)

K(T ) is the function describing the cooling process while m is the Ozawa expo-nent in dependance of the dimension of the crystal growth. C denotes the coolingrate (−dT/dt). The Ozawa exponent m has a very similar meaning as the Avramiexponent n and ranges between 0 and 4 as the Avrami index does (see Table 3-1).The Ozawa method represents a complex method which includes that the coolingrate influences both nucleation and crystal growth.

3.2 Experimental methods

An accurate method to study the crystallization kinetics of polymers is the visualobservation of crystal growth of spherulites that develop at a constant tempera-ture using polarized optical microscopy (POM) [84, 85]. However, if the studiedpolymer forms only small or many differently sized spherulites visual observationis limited [74]. Other methods to study the crystallization behavior of polymers areX-ray diffraction [85], wide angle X-Ray diffraction (WAXD) [84, 86, 87], small an-gle X-Ray scattering (SAXS) [87], Raman spectroscopy [77] and Fourier transforminfrared spectroscopy [87]. Over the last decades, the differential scanning calorime-try (DSC) has established as a method to study crystallization kinetics [77, 84–90].The simple sample preparation allows to study various material forms from pow-der, pellets to fiber reinforced test panels. However, the DSC represents an indirectmethod to determine the crystal forms, the crystallization type as well as the

Crystallization of polyamides 39

amount of crystallinity by measuring the enthalpies with subsequent calculationaccording to Avrami (Eq.3-1) or Ozawa (Eq.3-3).For the investigation of isothermal and non-isothermal crystallization behavior ofthe selected polyamides within this work, the DSC is used since powdery samples,pellets as well as reinforced polymers are analyzed. Parts of the following exper-imental work have been conducted within the frame of the student project fromPatrick Consul [91].

3.2.1 Differential scanning calorimetry

Test Set-Up

The measurement principle is based on the release of heat due to physical andchemical transformations of the test sample. Heat which is set free during e.g.crystallization (exothermic) or is consumed during fusion (endothermic) is mea-sured over time (isothermally) or over temperature (non-isothermally) [92, 93].The DSC set up consists of a measuring cell with two heaters controlled by thermo-couples. In the measuring cell, two pans are heated. The material being examinedis placed into a pan (sample pan), the second, typically empty, pan is used as areference [92, 93].Both pans experience the same temperature profile at the same time. However,the rate of temperature changes for a certain amount of heat differs between bothpans. To keep the temperature of both pans constant, the heat for the sample panis varied and results in a different heat output of the two heaters that is recorded.The heat is then recorded over temperature or time yielding the heat flow. Theenthalpy of fusion ΔHf and the enthalpy of crystallization ΔHc is determined byintegrating the endothermic (fusion) and exothermic (crystallization) peak [92, 93].The oven was flushed with nitrogen to avoid chemical reactions with oxygen fromair. The equipment used in this work was a DSC Q200 from TA Instruments.

Sample preparation and test procedure

Neat polymer pellets were selected which were comparable in weight and fit intothe pans in its entirety. To conduct measurements on tapes and carbon fiber re-inforced test panels, samples were carefully cut without additional heat input toavoid changes of the inherent crystallinity. Samples from UD tapes were extractedfrom the center to keep the fiber volume content as constant as possible. The sam-ples from test panels were directly taken from flexural test specimens.The test samples were not dried before conducting the DSC measurement as nodifference was detected due to the drying process in preliminary studies, as dis-cussed in subsection 3.3.1.

40 Crystallization of polyamides

Calibration of the cell constant was conducted with indium. The reference pan wasmaintained for all measurements. The lid of the pans was pressed by means of aconcave indenter to position the material being investigated centrically. The testprocedure followed the test standard prEN ISO 11357-7 [83].

3.2.2 Crystallinity ratio

Commonly, the degree of crystallinity XC is used as an expression for the amount ofcrystallinity in case of fiber reinforced thermoplastics and is determined as follows:

XC = ΔHf + ΔHcc

(1 − φ)ΔH0f

. (3-4)

ΔHcc denotes the enthalpy of cold crystallization, ΔHf is the enthalpy of fusion,ΔH0

f is the enthalpy of fusion of a 100 % crystalline polymer and φ is the fiberweight fraction. If neat polymers are studied [93], Equation 3-4 reduces to

XC = ΔHf + ΔHcc

ΔH0f

. (3-5)

However, this description is prone to errors, especially in the case of reinforced ther-moplastics. Determining the fiber weight fraction very precisely is crucial to relatethe heat flow to the mass of the polymer fraction only. The fiber weight fractioncan be determined by acid digestion and is usually averaged over an intermediatematerial or test panel. However, test samples used for DSC are small in size andthe averaged fiber weight fraction might not reflect the actual fiber weight fraction.This can cause a major error in the determination of XC .In addition, the enthalpy of fusion of a 100 % crystalline polymer ΔH0

f is usu-ally taken from literature and/or measured by X-ray diffraction. Currently, thereis no value in literature for the special chemical constitution PA10T/X of the co-polyamide C2000. For PA6, values for ΔH0

f vary from 190 to 260 J/g [92].Hence, a new expression, the Crystallinity Ratio (CR), is suggested to express theamount of crystallinity in reinforced thermoplastics:

CR = |ΔHf + ΔHcc|ΔHc

, (3-6)

where ΔHc is the enthalpy of crystallization, ΔHcc the enthalpy of cold crystal-lization and ΔHf represents the enthalpy of fusion. CR correlates the heat thatis required to melt the test sample (ΔHf ) to the heat that is released during themaximum crystallization (ΔHc) from the molten state. If cold crystallization isobserved between Tg and Tm, ΔHcc is subtracted from ΔHf (note that by conven-

Crystallization of polyamides 41

tion ΔHcc and ΔHf differ in sign). The maximum crystallization is achieved byselecting a cooling rate which allows the sample to crystallize as much as possible.Therefore, the amount of crystallinity is related to the technically maximum crys-tallization of the investigated polymer and independent of a theoretical value. Asall measurements take place at the same test sample, the fiber weight fraction isnot required to be determined. The analysis according to CR has been previouslypublished in [94] and successfully applied to carbon fiber reinforced PPS.For the present study, the cooling rate that yields maximum crystallization is deter-mined by non-isothermal measurements on neat polymers and subsequently usedto determine the CR of fiber reinforced samples.

3.2.3 Four-point bend test

To evaluate the effect of different cooling rates on crystallinity and on mechanicalproperties, test panels made of polyamide-sized fibers with PA6 (CF-TP/B3S) andPA10T/X (CF-TP/C2000) were produced as described in section 2.3.1 and 2.3.2using two different cooling rates in a static press. In addition, a previously pressedtest panel was additionally thermoformed. The thermoforming unit from Rucksused for this step is presented in Figure 3-2 along with the aluminum tool whichincorporates the test panel to be thermoformed. The aluminum tool is coated blackto reduce reflections during heating by an infra-red source. The tool comprises ofa support fixture to enable installation into the transport frame. When the toolmounted to the transport frame of the thermoforming unit is positioned in betweenthe IR heat sources the test panel is heated to T = Tm+40K. After the temperaturewas reached the transport frame transferred the laminate to unheated press platesof the thermoforming unit to apply a pressure of 10 bar. By this procedure, atypical production process for thermoplastic composites with high cooling rates ofup to 380 °C/min is followed.

3.2.4 Visualization of crystals

For further investigation of the influence of different carbon fiber sizings on thecrystallization behavior of B3S during processing, micrographs were prepared andetched to visualize crystal morphology.Potted and polished specimens made of B3S with different carbon fiber sizings wereprepared from test panels that were manufactured with a cooling rate of 20 °C/min.In addition, the polished specimens were subjected to oxygen plasma treatment.This serves to etch the amorphous regions of the semi-crystalline B3S sparing thecrystalline fractions. A Solarus Gatan Plasma Cleaner was used for the oxygen

42 Crystallization of polyamides

Press field

IR heatsourcesTransport

frame

Lower mold/ support for transport

Test panelUppermold

Figure 3-2 Rucks thermoforming unit and used aluminum tool to perform thermoforming of flattest panels.

plasma treatment. Oxygen plasma was applied for 5 min at 50 W and 40 sccm.Accordingly, the crystals are visualized to be examined via SEM by using a Hitachi4300 SE/N FESEM. Access to the facilities of the Centre for Advanced Microscopy(CAM) with funding through the Australian Microscopy and Microanalysis Re-search Facility (AMMRF) is gratefully acknowledged.

3.3 Results

3.3.1 Preliminary tests

Aliphatic polyamides, in particular PA6, are well-known for their hydrophilic be-havior. Moisture has a softening effect shifting Tg to lower temperatures. Hence,the relaxation temperature to perfect the crystal structure is lower giving the poly-mer more time for secondary crystallization. Moisture influences the primary andespecially the secondary crystallization. The mobility of the molecular chains isenhanced by moisture leading to an increase in growth rate during primary crys-tallization. This does not necessarily result in an increased crystallinity. Duringsecondary crystallization, the moisture intercalated in the amorphous regions in-

Crystallization of polyamides 43

volves also an enhanced molecular mobility and leads to an increase in crystallinityalso because of lowered Tg. This is dependent on the initial degree of crystallinity.The higher the initial degree of crystallinity, the fewer amorphous regions exist thatmay absorb moisture [75, 95].In preliminary tests, the heat flow was recorded and compared for samples withand without pre-drying for 24 h at 80 °C under vacuum as reported in Figure 3-3.

0 50 100 150 200 250

-20

0

20

40

0 50 100 150 200 250

-20

0

20

40

Fusion

Crystallization

Hea

tFlo

w[m

W]

Temperature [°C]

Undried B3S_1st heatingUndried B3S_2nd heating

EndothermalPeak

Crystallization

Fusion

a

Release ofmoisture

Hea

tFlo

w[m

W]

Temperature [°C]

Dried B3S_1st heatingDried B3S_2nd heating

b

Figure 3-3 Influence of pre-drying process on crystallization behavior of a) undried and b) pre-dried B3S.

During the first heating of undried B3S pellets an endothermal peak was evident atapproximately 65 °C. After drying specimens the previous endothermal peak van-ished but another endothermal peak appeared beginning at 100 °C. The latter isattributed to the release of moisture. The endothermal peak at 65 °C is assumed tobe caused by an additive added by the manufacturer to prevent moisture absorp-tion of pellets. The effectiveness of this additive is believed to be eliminated duringboth drying at 80 °C and melting. The second heating is unaffected by the initialdrying process. The crystallization peak is neither influenced by the drying processor number of heating cycles. In addition, moisture was either not completely re-moved by drying for 24 h in the first heating cycle or the short transfer from ovento DSC was sufficient for new moisture absorption. Figure 3-3 clearly shows thatthe crystallization process, which represents the focus of this study, is not affectedby drying. Thus, the test samples were not dried prior to measurement.In first experiments on neat C2000 a continuous decrease of ΔHc was detected thatmay arise from further polymerization of the co-polyamide after synthesis. C2000was therefore subjected to ten subsequent heating and cooling cycles as presentedin Figure 3-4.The baseline remains the same during the ten cycles indicating that no degradationoccurs and the temperature of 310 °C is chosen correctly to eliminate the thermalhistory. The first heating reveals a drastic cold crystallization peak that vanishes

44 Crystallization of polyamides

50 100 150 200 250 300

-20

-10

0

10

20

30

FusionΔHcc: Enthalpy ofcold crystallization

Hea

tFlo

w[m

W]

Temperature [°C]

1st Cycle2nd Cycle10th Cycle

ΔHcc

Crystallization

Figure 3-4 Subjection of C2000 to ten heating and cooling cycles.

for the subsequent cycles. The bimodal melting peak refers to melting of two crys-tal phases as it was also observed for PEEK [96]. After the first melting peak apost-crystallization occurs forming the crystal phase of the second melting peak.ΔHf decreases approximately by 10 % throughout ten cycles with the largest re-duction from first to second heating. The first melting peak remains the same afterthe second heating cycle, the second peak becomes smaller with increasing numberof cycles. The crystallization peak also decreases until the tenth cycle is finished.The considerable cold crystallization peak is assumed to originate from quenchingthe C2000 pellets after synthesis. To eliminate this uncertainty, the isothermal andnon-isothermal data obtained for C2000 is analyzed after the second heating.

3.3.2 Neat polymers

Isothermal crystallization

For isothermal measurements, B3S samples were fused at 250 °C and then cooled toseveral holding temperatures between 201 and 208 °C. Three samples were testedper holding temperature. The temperature was then held constant until the heatflow dropped to zero. C2000 samples were fused at 310 °C and cooled to hold-ing temperatures between 228 to 235 °C. The selected temperatures are within anarrow range. At lower temperatures the polymers started to crystallize alreadyduring cooling and disable the analysis of the isothermal crystallization behavior.All measurements were found valid when the heat flow was 0 before the crystal-lization process started. Based on the DSC measurements, the relative degree ofcrystallinity X(t) was calculated according to the following equation:

X(t) =∫ t

0(dHc/dt)dt∫ ∞0 (dHc/dt)dt

= ΔHc(t)ΔHc,∞(t) = Wc, (3-7)

Crystallization of polyamides 45

where dH/dt expresses the rate of heat evolution, ΔHc(t) represents the heat flowproduced at time t; and ΔHc,∞(t) is the total amount of heat created until the endof the isothermal crystallization process. This ratio can be also expressed as thecrystallized mass fraction Wc.The development of the relative degree of crystallinity X(t) is depicted in Figure 3-5for neat B3S and C2000. Comparing both polymers, the maximum X(t) is obtained

0 500 1000 1500 2000 25000

20

40

60

80

100

0 500 10000

20

40

60

80

100

Rel

ativ

e X(

t) [%

]

Time [s]

B3S_201°C B3S_202°C B3S_203°C B3S_205°C B3S_206°C B3S_208°C

a

Rel

ativ

e X(

t) [%

]

Time [s]

C2000_228°C C2000_229°C C2000_230°C C2000_232°C C2000_233°C C2000_235°C

b

Figure 3-5 Development of the relative degree of crystallinity X(t) for a) neat B3S and b) C2000.

more rapidly for C2000 than for B3S. X(t) determined by DSC represents the crys-tallized mass fraction (Wc). This data is to be converted into volume crystallinity(Vc) as follows to determine isothermal crystallization kinetic parameters:

Vc =

Wc

ρc

Wc

ρc

+ 1 − Wc

ρa

. (3-8)

ρc and ρa refer to the density of the crystalline and amorphous polymer fractions.For the analysis according to Avrami, Equation 3-1 is solved:

log{

− ln[1 − Vc(t)]}

= nlogt + logk. (3-9)

Plotting log{

− ln[1−Vc(t)]}

against log t yields the Avrami plot with straight lineslinearly fitted by using the least square method. The Avrami index n is obtainedfrom the slope of the linear fit whereas the crystallization rate k is determinedexperimentally (kexp) from the intercept (Figure 3-6) for the linear fit. k can alsobe calculated by the following equation [97], based on the half-crystallization timet1/2:

kcalc = ln2(t1/2)n

. (3-10)

46 Crystallization of polyamides

-1.5 -1.0 -0.5 0.0 0.5 1.0 1.5

-8

-6

-4

-2

0

2

-1.5 -1.0 -0.5 0.0 0.5 1.0 1.5

-8

-6

-4

-2

0

2

B3S_201°CB3S_202°CB3S_203°CB3S_205°CB3S_206°CB3S_208°C

log[

-ln(1

-Vc(t

))][-]

log t [min]

a

C2000_228°CC2000_229°CC2000_230°CC2000_232°CC2000_233°CC2000_235°C

log[

-ln(1

-X(t)

)][-]

log t [min]

b

Figure 3-6 Plot of log{ − ln[1 − X(t)]

}versus logt for the isothermal crystallization of a) B3S

with ρa = 1.08 g/cm3 and ρc = 1.24 g/cm3 b) C2000, based on Wc according to [98].

Table 3-2 summarizes the determined Avrami exponent n, the crystallization ratek, along with the time at maximum heat flow tmax, the half-crystallization timet1/2 and the crystallinity at tmax for B3S. The correlation coefficient of all linearfits for the Avrami plots (R2) was R2 > 0.99. The isotherms cannot be super-

Table 3-2 Kinetic parameters for the isothermal crystallization of neat B3S.

Tc [ °C] 201 202 203 205 206 208n [-] 2.4 2.0 2.6 2.7 2.3 2.1kexp[x10−2min−n] 14.16 10.34 3.45 0.71 0.50 0.26kcalc[x10−2min−n] 10.71 10.40 2.62 0.56 0.53 0.30tmax [min] 1.51 1.83 2.73 4.50 6.06 9.12t1/2 [min] 2.19 2.57 3.60 6.17 8.13 13.74X(tmax) [%] 46.21 51.68 45.32 44.53 47.67 39.27

imposed by shifts along the logarithmic time axis. Nucleation mode is thereforethermal and indicates heterogeneous or homogeneous nucleation. Considering thecalculated Avrami exponents, the time-dependance of the crystallization process isindicated by the decimal places (nn) and yield values smaller than 0.5 for four tem-peratures that refer to heterogeneous nucleation. With exception of nn at 205 °Cand 206 °C, the results for nn indicate a heterogeneous nucleation mode. As B3S ispigmented with carbon black it is expected that the nucleation mode is heteroge-neous. The integer nd yielded 2 for all measurements and refers to the formation offibrils according to thermal nucleation. The crystallization rate constant kexp yieldsgood agreement with kcalc and decreases with decreasing undercooling ΔT .

Crystallization of polyamides 47

For neat C2000, the results from the analysis according to Avrami are summarizedin Table 3-3. Interpretation of the yielded Avrami exponents for C2000 concludes

Table 3-3 Kinetic parameters for the isothermal crystallization of neat C2000.

Tc [ °C] 228 229 230 232 233 235n [-] 2.24 2.29 2.11 2.17 2.14 2.14kexp[min−n] 0.71 0.69 0.41 0.25 0.14 0.05kcalc[min−n] 0.60 0.58 0.41 0.24 0.13 0.05tmax [min] 0.87 0.92 1.09 1.41 1.83 2.79t1/2 [min] 1.07 1.08 1.29 1.63 2.17 3.48X(tmax) [%] 23.35 22.16 21.82 19.34 16.89 16.38

athermal nucleation due to isotherms (Figure 3-6) that can be superimposed byshifting along the logarithmic time axis. Hence, the nucleation mode is heteroge-neous, indicated by athermal nucleation and decimal places for nd between 0.1 and0.3, with growth of lamellae (nn = 2). In contrast to B3S, C2000 yields highervalues for crystallization rate k. On the one hand this is contradictory to resultsfrom PA6T where terephthalic acid is known to inhibit crystallization [75]. How-ever, C2000 consist also of other monomers (indicated by affix ‘X’ in PA10T/X),not further specified by the manufacturer, that can act as a foreign nucleus. Thestrong heterogeneous nucleation behavior implied by the superimposable isothermsconfirm this assumption. In this case, the crystallization-inhibiting nature of theterephthalic acid may be overruled by the dominating aliphatic regions with aconsiderable number of monomers acting as nucleating agent.

Non-isothermal crystallization

The non-isothermal measurements on neat polymers were conducted at the coolingrates 2, 5, 10, 20, 35 and 50 K/min. Three samples per cooling rate were investi-gated. The relative degree of crystallinity X(T ) as a function of temperature wascalculated as follows:

X(T ) =∫ T

0 (dHc/dT )dT∫ ∞0 (dHc/dT )dT

= ΔHc(T )ΔHc,∞(T ) , (3-11)

where ΔHc(T ) is the enthalpy of crystallization at a certain temperature dividedby ΔHc,∞(T ), the enthalpy of crystallization when the crystallization process iscompleted. The mean relative degree of crystallinity X(T ) is depicted in Figure 3-7 for neat B3S and C2000.

48 Crystallization of polyamides

100 120 140 160 180 200 2200

20

40

60

80

100

120 140 160 180 200 220 2400

20

40

60

80

100R

elat

ive

X(T

)[%

]

Temperature [°C]

2°C/min5°C/min10°C/min20°C/min35°C/min50°C/min

a

Rel

ativ

eX

(T)[

%]

Temperature [°C]

2°C/min5°C/min10°C/min20°C/min35°C/min50°C/min

b

Figure 3-7 Development of the relative degree of crystallinity X(T ) for a) neat B3S and b)C2000.

The analysis of the non-isothermal measurements was conducted according to theOzawa method by solving Equation 3-3:

log{

− ln[1 − X(T )]}

= logK(T ) − mlogC. (3-12)

Plotting log{

− ln[1 − X(T )]}

against log C shall yield a set of straight lines atconstant temperatures. The Ozawa exponent −m is obtained from the slope ofthe straight lines and the cooling function K(T ) from the intercept of the initiallylinear region with the y-axis. Figure 3-8 shows the Ozawa plot for neat B3S andneat C2000 at constant temperatures. For the investigated cooling rates, the Ozawa

0.0 0.5 1.0 1.5 2.0

-8

-6

-4

-2

0

0.0 0.5 1.0 1.5 2.0

-8

-6

-4

-2

0

log(

-ln(1

-X(T

))

log C

205°C200°C195°C190°C185°C

a

log(

-ln(1

-X(T

))

log C

235°C230°C225°C210°C205°C

b

Figure 3-8 Plots of log{ − ln[1 − X(T )]

}versus ln|(dT/dt)−1| for a) neat B3S at 205, 200, 195,

190 and 185 °C; b) neat C2000 at 235, 230, 225, 215 and 205 °C.

plots yield no straight lines at the majority and disallow the determination of m

and K(T ). This is a common problem in the analysis of non-isothermal crystal-lization data caused by the limited applicability of the Ozawa theory [99]. Ozawa

Crystallization of polyamides 49

made the assumption that the dependance of the nucleation rate on the tempera-ture is not affected by the cooling rate. Thus, the analysis according to Ozawa isrestricted to a narrow range of cooling rates leading to crystallization in resemblingtemperature ranges [87]. The Nakamura theory offers another possibility to analyzenon-isothermal crystallization data and considers non-isothermal measurements asconsecutive isothermal processes. This yields good results only for very slow cool-ing rates when the nucleation is unaffected by the cooling rate. Several authorsintroduced new approaches to handle non-isothermal crystallization data. Liu etal. [100] combined the Avrami (Eq. 3-1) and Ozawa equation (Eq. 3-3) to form anew kinetic equation:

logk + nlogt = logK(T ) − mlogC (3-13)

This method was also investigated for the present non-isothermal results but led tounsatisfactory results. Although this method is used by many researchers, the theo-ries according to Avrami and Ozawa require different assumptions: Avrami assumesthe growth rate and Ozawa the cooling rate to be constant. Hence, combining theAvrami and Ozawa equation into one equation, possesses no theoretical substantia-tion. This also applies to other methods developed by Jeziorny or Kissinger yieldingresults without theoretically proven physical meaning [87].Within this work, the non-isothermal measurements obtained from DSC are notanalyzed further according to another method as the current procedures may betoo simplified to properly account for complex crystallization behaviors. Never-theless, the results from the non-isothermal DSC data relating ΔHc, X(T ), thepeak temperature Tp during crystallization at the maximum rate of heat flow todifferent cooling rates (presented in Table 3-4) generates an understanding of thenon-isothermal crystallization behavior.

Table 3-4 Effect of cooling rate on crystallization of neat B3S.

C [ °C/min] Tp [°C] ΔHc[J/g]2 200.0 77.90±2.225 195.7 78.19±0.3710 191.7 75.90±0.8520 186.0 71.73±1.0235 181.9 63.36±0.6150 174.8 62.79±1.84

Specifically, Tp possesses practical importance since it refers to the maximum rateof crystallization as a function of the cooling rate [76]. This enables a better controlof the final polymer properties during processing. The obtained values for ΔHc are

50 Crystallization of polyamides

within the same range as reported by Fornes and Paul [86] for low to mediummolecular weight nylons and found by Tjong and Bao [101].Table 3-5 shows ΔHc, X(T ) and the peak temperature Tp during crystallization atthe maximum rate of heat flow in dependance of different cooling rates for C2000.

Table 3-5 Effect of cooling rate on crystallization of neat C2000.

C [ °C/min] Tp [°C] ΔHc[J/g]2 232.0 42.53±3.615 226.3 45.54±1.2710 218.6 44.87±0.5420 207.1 46.91±1.5135 194.2 43.59±1.5450 183.9 25.95±6.04

Figure 3-9 compares the enthalpy of crystallization ΔHc as a function of the coolingrate for B3S and C2000. The crystallization process proceeds most rapidly at acooling rate of 2 °C/min for B3S. The maximum ΔHc for C2000 is obtained at acooling rate of 20 °C/min.

0 10 20 30 40 50

20

30

40

50

60

70

80

90 B3SCubic fitC2000Cubic fit

Ent

halp

yof

Cry

stal

lizat

ion

ΔHc

[J/g

]

Cooling rate [°C/min]

Figure 3-9 Plots of cooling rate versus enthalpy of crystallization ΔHc for B3S and C2000.

Crystallization of polyamides 51

3.3.3 Influence of carbon fibers on crystallization

As the previous results obtained from the non-isothermal crystallization measure-ments on neat B3S and C2000 show, significant differences in the degree of crys-tallinity are reached at cooling rates of 2, 10 and 20 °C/min. In case of B3S, themaximum Xc is expected for the cooling rate of 5 °C/min. For C2000, the maxi-mum crystallinity is anticipated at 20 °C/min.The influence of carbon fibers with epoxy- and polyamide-based sizing (see Ta-ble 2-1) is investigated for both B3S and C2000. The material combinations CF-EPY/B3S, CF-TP/B3S and CF-TP/C2000 were present in form of UD tapes sup-plied by SGL Carbon GmbH. The DSC measurements were conducted on a stack ofcircular-shaped stamped tapes. For CF-EPY/C2000 8 plies of powder-coated tapeswere stacked and pressed with 10 bar at 300 °C for 30 min. From this laminate,square-shaped specimens (length, width and thickness of 1 mm) were extractedand investigated in the DSC.Varying matrix mass fractions of the tapes might falsify the results. Hence, thematrix mass fraction was calculated from the mean of three measurements ofthe matrix volume fraction determined by microscopy. The matrix mass fractionswere determined to 26.25±4.04 wt% for CF-EPY/B3S and 39.35±8.01 wt% forCF-TP/B3S. A matrix mass fraction of 24.35±2.34 wt% was obtained for CF-EPY/C2000 and 40.87±8.06 wt% for CF-TP/C2000. The measured heat flows andthe integrated enthalpies were then normalized by the matrix mass fraction.The influence of the carbon fiber sizing on mean values of ΔHc compared to neatB3S and C2000 is presented in Figure 3-10.

2 4 6 8 10 12 14 16 18 20

60

80

100

120

140

160

180CF-EPY/B3SCF-TP/B3SB3S

Cooling rate [°C/min]

ΔHc

[J/g

]

a

0 2 4 6 8 10 12 14 16 18 20 2230

40

50

60

70

80

90CF-EPY/C2000CF-TP/C2000C2000

Cooling rate [°C/min]

ΔHc

[J/g

]

b

Figure 3-10 Influence of carbon fiber sizing on ΔHc for tapes made of a) B3S and b) C2000.

For both polymers, ΔHc is considerably enhanced in presence of carbon fiberswith epoxy-based sizing. The crystallization behavior of C2000 is unaffected by thepolyamide-based sizing referring to comparable values for ΔHc of the neat polymer.

52 Crystallization of polyamides

The influence of the polyamide-based sizing on crystallization of B3S is inexistentfor a cooling rate of 2 °C/min and slightly increases with faster cooling rates.Epoxy-sized carbon fibers however have a significant effect on the non-isothermalcrystallization of B3S. ΔHc of CF-EPY/B3S is increased by 65 % in contrast toneat B3S. ΔHc of neat C2000 is enhanced by 33 % due to the epoxy-based sizing.Thus, the epoxy-based sizing intensifies the crystallization for both polymers byacting as heterogeneous nucleating agent. Thus, epoxy-sized carbon fibers representa foreign interface allowing B3S and C2000 to crystallize heterogeneously. This isconfirmed by the study of Sang and Wei [102]. They attributed the role of hetero-geneous nucleating agents on PA6 for carbon fibers that were surface-treated withsilane coupling agents and toughened elastomers. As well as the epoxy-sized fiberswithin this work, the surface treatment by Sang and Wei [102] resulted in a foreigninterface acting as nucleation site for PA6.The polyamide-based sizing is assumed to lower the effect of a heterogeneous nu-cleating agent by being chemically similar to B3S and to the aliphatic regions ofC2000.

Visual inspection

The considerable enhancement of ΔHc for CF-EPY/B3S was believed to be causedby possibly enhanced transcrystallization around epoxy-sized fibers. Polished andetched specimens of CF-EPY/B3S were visually inspected under the SEM in com-parison to CF-TP/B3S (see Figure 3-11). The oxygen plasma treatment provedto be sufficiently long and intense to visualize the crystal structures. Locations of

Possiblenuclei

a) CF-EPY/B3S b) CF-TP/B3S

2 μm 2 μm

Figure 3-11 Close-up view of B3S specimens with a) epoxy-sized fibers and b) polyamide-sizedfibers that were etched by oxygen plasma to visualize crystalline structures

possible nuclei are marked in Figure 3-11 where radial growth of crystals can bedetected. Comparing the structured surface of both CF-EPY/B3S and CF-TP/B3Sno significant differences can be detected qualitatively. The amount of crystal struc-tures around carbon fibers with epoxy- and polyamide-based sizing appears to beof comparable magnitude.

Crystallization of polyamides 53

3.3.4 Relation between mechanical properties and crystallinityratio

To eliminate the plasticizing effect of moisture on crystallinity and more impor-tant on mechanical properties, all test panels were dried at 80 °C for more than60 h under vacuum. This temperature lies above Tg of B3S and below Tg of C2000.The temperature for drying was chosen to ensure that moisture is removed mostefficiently and the mechanical properties remain comparable to Chapter 2. Thediffusion rate of water increases with increasing temperature due to increased mo-bility of molecular chains [103]. Considering the opposite of moisture removal, thewater absorption rate increases from 1.05 x 10−12m2/s at 40 °C to 14.19 x 10−12m2/sat 80 °C [104] of reinforced PA66. PA66 and PA6 resemble each other in their waterabsorption due to the same ratio of CH2/CONH groups. Drying at 40 °C, belowTg of B3S, would therefore require a long time. However, secondary crystallizationcannot be excluded by drying at temperatures above Tg.Pressure-induced changes in crystallinity during test panel manufacture can beexcluded since significant crystal modifications are usually obtained at pressuresabove 500 MPa [105]. All test panels were manufactured with a pressure of 1 MPa.Table 3-6 gives an overview of the produced test panels along with matrix massfraction determined by acid digestion according to DIN EN 2564 method B [106]at the central laboratory of SGL Carbon GmbH. In case of test panels, the ther-

Table 3-6 Overview of produced test panels along with matrix mass fraction determined by aciddigestion.

Cooling rate[ °C/min] Designation Manufacturing

processMatrix massfraction [wt%]

2 CF-TP/B3S_02 Static press 36.50±0.6920 CF-TP/B3S_20 Static press 35.30±0.62≈380* CF-TP/B3S_TF Thermoforming 37.20±0.202 CF-TP/C2000_02 Static press 27.20±2.1620 CF-TP/C2000_20 Static press 31.43±1.10≈320* CF-TP/C2000_TF Thermoforming 35.17±1.19*measured by thermocouple

mal history and the process-induced crystallinity is to be investigated. Thus, ΔHf

from the first heating (including ΔHcc) is divided by the maximum possible ΔHc

as described by CR in Equation 3-6. Maximum ΔHc was yielded by cooling thesamples with 5 °C/min for B3S and 20 °C/min for C2000. As investigated on neatpolymers, ΔHc reached maximum values for these cooling rates (see Figure 3-10).The relation of mean CR to the mean transverse flexural strength σf2 and the mean

54 Crystallization of polyamides

transverse flexural modulus Ef2 is presented in Figure 3-12 for both CF-TP/B3Sand CF-TP/C2000.

0 50 100 150 200 250 300 350 4000.0

0.2

0.4

0.6

0.8

1.0

CRCR undried

Cooling rate [°C/min]

CR

[-]

a

0

20

40

60

80

100

120

140

σf2σ f2

[MP

a]0 50 100 150 200 250 300 350 400

0.0

0.2

0.4

0.6

0.8

1.0

CR

Cooling rate [°C/min]C

R[-]

b

0

10

20

30

40

50

60

70

σf2

σ f2[M

Pa]

0 50 100 150 200 250 300 350 4000.0

0.2

0.4

0.6

0.8

1.0

CRCR undried

Cooling rate [°C/min]

CR

[-]

c

6

7

8

9

10

11

Ef2

Ef2

[GP

a]

0 50 100 150 200 250 300 350 4000.0

0.2

0.4

0.6

0.8

1.0

CR

Cooling rate [°C/min]

CR

[-]

d

0

2

4

6

8

10

12

Ef2

Ef2

[GP

a]

Figure 3-12 Influence of CR on σf2 and Ef2 of a), c) CF-TP/B3S and b), d) CF-TP/C2000.

In case of CF-TP/B3S, the CR exceeds 1 at a cooling rate of 2 °C/min. This isattributed to the fact that all samples were cooled at 5 °C/min. This cooling ratewas selected to reach maximum crystallinity as the results from neat polymers sug-gested. The values for ΔHc at a cooling rate of 2 °C/min and 5 °C/min howeverare very close to each other. Thus, the maximum crystallinity may be achieved at2 °C/min instead of 5 °C/min leading to an CR>1.As expected from non-isothermal measurements on neat B3S, CR decreases withincreasing cooling rates. In addition to DSC measurements on dried samples, theCR of undried specimen was determined to revise possible secondary crystalliza-tion induced by pre-drying. Undried samples show slightly lower CR than driedspecimen. However, the overlapping error bars indicate the deviation is not sig-nificant. The CR of CF-TP/C2000 develops also accordingly to measurements onneat polymer. At 20 °C, CR reaches its maximum. In general, fast cooling ratesthat occur during thermoforming have a larger impact on the crystallinity of C2000

Crystallization of polyamides 55

than on B3S as noticed for measurements on neat polymers.Usually, increasing crystallinity leads to improved mechanical properties such asstiffness and strength. However, the opposite behavior is observed regarding strengthσf2 for both polymers. Increases in CR involves decreased strength and vice versa.On the other hand, stiffness Ef2 declines with decreasing CR as expected.For CF-TP/B3S, drastic changes in mechanical properties effected by undercoolingare not anticipated since CR remains nearly constant for a wide range of differentcooling rates. C2000 reacts much more sensitive to different cooling rates. CR ofCF-TP/C2000 reduces from 0.98 at 20 °C/min to 0.25 at 320 °C/min. Despite themajor effect on CR, the mechanical properties do not change within the same mag-nitude. However, a significant increase in strength is observed for the lowest CR.The gain in strength is believed to arise rather from the lowest measured fiber vol-ume content than from a change in crystallinity. The transverse strength typicallyincreases with decreasing fiber volume content as less defects are present.

3.4 Conclusion and implications

The isothermal measurements on neat polymers by using DSC revealed hetero-geneous nucleation behavior for both B3S and C2000. For C2000, the nucleationis even independent of the temperature revealing that only foreign particles actas crystal nuclei. These foreign particles are assumed to be the monomers X thatare added to PA10T monomers by the manufacturer during synthesis (PA10T/X).The analysis of the isothermal crystallization kinetics according to Avrami yieldedfibrils for B3S and lamellae for C2000 as crystal forms. B3S revealed a crystalliza-tion rate slower than C2000 by two magnitudes. If it is desired to reach maximumcrystallinity after manufacturing, B3S is required to be cooled to 202 °C and heldconstant for approximately 10 min. C2000 may be cooled from melt to 228 °C andheld at this temperature for approximately 5 min to achieve maximum crystallinity.Although analysis of the non-isothermal crystallization kinetics on neat polymersfailed by using the method according to Ozawa [81] and Liu et al. [100], importantresults for processing were obtained. B3S revealed the maximum crystallizationat a cooling rate of 5 °C/min with decreasing crystallization ability for increas-ing cooling rates from 2 to 50 °C/min. In contrast, C2000 showed the maximumcrystallization at a cooling rate of 20 °C/min and a rapid decline of crystallinitytowards larger cooling rates up to 50 °C/min.The analysis of epoxy- and polyamide-sized carbon fibers on the crystallization ofB3S and C2000 showed a strongly nucleating behavior of the epoxy-sized fibers onboth polymers. There was no influence of the polyamide-sized carbon fibers on thecrystallization of C2000. B3S showed a slightly enhanced enthalpy of crystallization

56 Crystallization of polyamides

with increasing the cooling rate from 2 to 20 °C/min for B3S. Carbon fibers canact as nucleating agents. The low nucleating effect of the polyamide-sized fibers isattributed to the chemical similarity of the sizing to B3S and C2000.Test panels produced from CF-TP/B3S in a static press cooled at 2 and 20 °C/minas well as in a thermoforming unit cooled with approx. 380 °C/min revealed onlyslight differences in the CR and hence negligible differences in transverse flexu-ral strength. The transverse flexural modulus slightly decreased with decreasingCR. CR remained constant for CF-TP/C2000 when pressed and cooled at 2 and20 °C/min resulting in comparable strength but slightly decreased stiffness. A dras-tic drop in CR for CF-TP/C2000 thermoformed and cooled at 320 °C/min revealeda strong sensitivity of C2000 to large cooling rates. However, the transverse mod-ulus was not affected in the same manner and was found to be slightly decreased.The strength was even increased despite of the low CR and attributed to lowerfiber volume content of the thermoformed test panel compared to the staticallypressed panels.The introduced CR proved to be a suitable method to describe and compare theamount of crystallinity in fiber reinforced B3S and C2000.

4 Impregnation modelIn this chapter, a model is developed to observe the impregnation progress duringconsolidation of powder-coated tows to tapes and laminates. The governing equa-tions for the transport phenomena that occur during transverse flow are presented.Additionally, analytical models to represent the processing phenomena of the indi-vidual constituents, fiber and matrix, are summarized. Based on existing analyticalapproaches, a one-dimensional (1D) model is derived to compute the impregnationprogress through the thickness during processing of powder-coated tows and con-solidated tapes.The experimental work comprises the development of a method to determine thedegree of impregnation (DOI). This technique is further used to assess the resultsof an impregnation study which is set up by using design of experiments (DOE). Byvarying the most important process parameters that govern impregnation - time,temperature, and pressure - the study serves to verify the derived model experimen-tally for the material combinations CF-TP/B3S and CF-TP/C2000. In addition,the interlaminar shear strength (ILSS) is determined for both intermediates andrelated to the DOI.

4.1 Transverse resin flow

Under application of pressure and at elevated temperatures, intermediates areconsolidated that are initially pre-impregnated or pre-formed. Consolidation pro-cesses aim to complete impregnation, remove the entrapped air and surplus matrix.Double-belt press forming or processing of film-stacked prepregs, pre-impregnatedtows or powder-coated tows belong to the class of consolidation processes [107].This process is commonly described by resin flow or resin impregnation.Models to describe composites manufacturing processes generally combine funda-mental laws such as conservation of mass, energy and momentum with specificempirical models for e.g. viscosity and permeability [108]. The governing equationsof mass conservation consider a representative volume element (ΔV ) to averagethe material properties of all constituents forming a composite. Having two con-stituents in the composites, the equation for mass conservation is applied to bothfibers and matrix [109, 110]:

∂Vf

∂t+ ∇(Vf uf ) = 0 (4-1)

57

58 Impregnation model

∂(1 − Vf )∂t

+ ∇((1 − Vf )um) = 0 (4-2)

Vf denotes the fiber volume fraction, uf describes the average velocity of the fiberbed whereas um indicates the average velocity of the matrix.The conservation of momentum is described by Darcy’s Law [111] assuming satu-rated laminar flow and neglecting gravity:

(1 − Vf )(um − uf ) = −¯Kη

∇P (4-3)

¯K represents the permeability tensor of the fiber bed, η describes the dynamic ma-trix viscosity, and P the applied pressure. Darcy’s Law has been used by many re-searchers [25, 109, 110, 112–120] to describe thermoplastic matrix flow through fiberbeds. Originally, Henry Darcy formed the specialized momentum balance equationfor flow of water through a granular bed of sand in 1856. As a consequence, severalassumptions have to be made and fulfilled for thermoplastic matrix flow througha unidirectional carbon fiber bed, the focus of the present work. The boundaryconditions of Darcy’s Law and during processing of intermediates into parts madeof thermoplastic composites are compared in Table 4-1.

Table 4-1 Comparison of boundary conditions as present in Darcy’s Law and thermoplasticmatrix flow through carbon fiber bed [121].

Darcy’s Law CFRTP processing

Fluid Newtonian (water) non-Newtonian, visco-elastic

Porous mediaisotropic anisotropichomogeneous inhomogeneous

Flowcharacteristics

1D1D (unidirectional)quasi-1D, multidimensional(woven fabrics)

Darcy’s Law is based on ideally rigid, porous medium with spherical elementssimilar in size (sand). In case of spread unidirectional fiber tows, the distributionis assumed to be homogeneously as well in contrast to e.g. woven fabrics. Thefibers are also more or less spherical elements with regard to their cross-sectionalarea but with extension in length. Thermoplastics are generally characterized bynon-Newtonian behavior considering the dependance of the viscosity on the ap-plied shear rate. Under sufficiently low shear rates, thermoplastics behave quasi-Newtonian and fulfill therefore the requirements of Darcy’s law. In addition, thevalidity of Darcy’s Law is restricted to laminar flow with a Reynolds number < 1.

Impregnation model 59

Generally, the high melt viscosity of most thermoplastics causes low Reynolds num-bers and laminar flow [122].In principle, spontaneous impregnation can occur when matrix and carbon fibersizing increase the capillary forces. For thermoplastics known for high melt viscosi-ties, spontaneous impregnation is less likely to be achieved in a reasonable time.Hence, external pressure is generally applied [118].The pressure is assumed to be carried by fibers and matrix during the impregnationand consolidation process. In comparison to the applied pressure with a range of5 to 40 bar, the capillary pressure is low and hence disregarded [118]. Neglectinginertia-related forces, the stress equilibrium yields [109]

∇P + ∇¯σ = 0, (4-4)

where ¯σ denotes the average effective stress tensor that acts on the fiber bed. ∇P

describes the applied pressure [123].

4.2 Processing phenomena of individual constituents

The heterogeneous structure of composites involves also specific phenomena of theindividual constituents that occur during transverse flow. The packing geometryof the fiber bed influences the permeability and the compaction behavior. The vis-cosity of the matrix is characterized as function of temperature and shear rate. Allprocessing phenomena appear simultaneously during transverse flow and must beconsidered to develop a model reproducing reality comprehensively. In the follow-ing, several approaches to model processing phenomena of the individual compositeconstituents are presented.

4.2.1 Fiber bed properties

The reinforcing structure of a composite may be characterized by the packing ge-ometry that further influences the fiber volume content and hence the permeabilityas well as the compaction behavior of the fiber bed.

Packing geometry

Focusing on unidirectionally aligned 50k carbon fiber tows (see Table 2-1), thefibers can take different positions relatively to each other. The most commonlyconsidered idealized fiber packing are the quadratic and hexagonal arrangement,as presented in Figure 4-1 along with the maximum possible fiber volume content.

60 Impregnation model

The hexagonal arrangement leads to a dense packing with less interstices betweenthe fibers than the quadratic arrangement.

Quadratic fiber packing

Hexagonal fiber packing

Figure 4-1 Idealized arrangements of fiber packings in a composite yielding maximum fibervolume contents Vf,max of 78.5 % (quadratic) and 90.7 % (hexagonal).

Both arrangements underlie the assumption of a perfect alignment which may notreflect reality that often reveals mixed packing geometries.

Transverse permeability

Many different methods and modeling approaches have been developed over thelast decades to determine the permeability of different fiber architectures. Experi-mentally determined values for permeability are difficult to compare due to a lackof standardization. This leads to a large scatter of permeability values of more thanone magnitude when different experimental test setups are used. This is especiallythe case for the measurement of dual-scale permeability as found in woven fab-rics [124].The permeability of fiber tows however may be approximated by analytical expres-sions. Already in the early 20th century, Kozeny [125] and Carman [126] furtherdeveloped the empirical constant in Darcy’s Law on the basis of the capillary model.The Kozeny-Carman equation may be expressed as:

Kzz =r2

f

4kzz

(1 − Vf )3

V 2f

, (4-5)

where rf denotes the fiber radius, and kzz describes the so-called “Kozeny con-stant”. The Kozeny-Carman equation predicts isotropic permeability which is validfor isotropic porous media with flow resistance independent of the direction. In caseof unidirectional fibers, anisotropic behavior has to be considered: the permeabilityin fiber direction is much higher than transverse to the fibers.There have been several attempts to extend the Kozeny-Carman equation from anisotropic sphere bed to an anisotropic fiber bed. Gutowski et al. [127] proposed anextension of the Kozeny-Carman approach by adapting the Kozeny constant ac-

Impregnation model 61

cording to the respective fiber direction. Focusing on transverse flow, the expressionfor permeability of a fiber bundle is:

Kzz =r2

f

4k′zz

(√

V ′a

Vf

− 1)3

(V ′a

Vf

+ 1), (4-6)

where 4k′zz describes the Kozeny constant modified by Gutowski et al. [127]. Vf

is the current fiber volume content and V ′a expresses the maximum available fiber

volume content.Besides models that are based on the capillary approach by Kozeny and Carman,other analytical approaches replace the isotropic sphere bed by an array of alignedcylinders and compute the drag resistance to obtain the permeability. The modelsdeveloped by Gebart [128] as well as Bruschke and Advani [129] belong to thisclass and are widely used to predict permeability as a function of fiber packing,Vf , rf and fiber direction. The transverse permeability according to Gebart maybe expressed as

Kzz =C1(

√Vf,max

Vf

− 1)5

2r2f

, (4-7)

with C1 = 16/9π√

6 for hexagonal fiber packing and C1 = 16/9π√

2 for quadraticfiber arrangement.Bruscke and Advani [129] based their analytical approach on the lubrication theoryand a relationship between porosity and permeability in terms of a unit cell insteadof conventionally used idealized fiber packing:

Kzz =r2

f

3(1 − l2)2

l3 (3lB + l2

2 + 1)−1, (4-8)

where l2 is defined asl2 = 4

πVf (4-9)

and B is described by

B =arctan(

√1+l1−l

)√1 − l2

. (4-10)

Compaction

During the considered manufacturing processes in a double-belt or a static press,the fibrous reinforcement is subjected to considerable pressure in thickness direc-tion. This leads to a compression of the fibers involving an increase in fiber volumecontent Vf . This change in Vf has a significant impact on permeability as it is calcu-

62 Impregnation model

lated as a function of Vf in all previously introduced analytical models. The effectof transverse compression on Vf was investigated by Gutowski et al. [112, 127] forimpregnated, unidirectional fiber tows. Considering the compacted fibers as elas-tic springs embedded in a visco-elastic matrix, the fibers are modeled as bendingbeams between contact points that establish due to fiber misalignment or slightfiber waviness. More and more such contact points are assumed to evolve as soonas the composite suffers transverse compaction. The effective stress on the fibers iscalculated according to the following equation [127]:

σ = As

√Vf

V0− 1

(√

Va

Vf

− 1)4. (4-11)

As represents a constant, V0 expresses the initial fiber volume content, and Va

describes the available fiber volume content and represents the maximum possiblefiber volume content. As, V0, and Va are commonly determined by fitting the modelto experimental data.By neglecting inertia effects, Gutowski et al. [112] assume the applied pressure, theresin pressure, and the effective stress that is carried by the fibers to balance:

P = Pf · Vf + Pr. (4-12)

This forms a non-linear spring-damper set in parallel arrangement. This also allowsto take into account the elastic recovery of the fibrous reinforcement as soon as thepressure is reduced while the matrix is still molten [29].

4.2.2 Matrix

Flow behavior dominates the melt processing characteristics of thermoplastics andis commonly expressed in terms of the viscosity. Viscosity represents the resistanceof a fluid (here: polymer melt) against applied shear stress. Considering flow ofpolymer melts to take place between two plates where the upper one moves, New-tonian’s law of viscosity can be derived as ratio of shear stress τ to shear rate γ:

η = Shear stressShear rate = τ

γ. (4-13)

The viscosity of thermoplastics is not only dependent on their chemical composi-tion but also on temperature, pressure, shear rate and time. Pressure and hencethe production rate can be derived from viscosity data when it is determined

Impregnation model 63

as a function of the shear rate. A simple expression to describe non-Newtonian,shear-thinning behavior of thermoplastics represents the power-law model. Otherfrequently used models to account for shear-thinning of polymer melts were de-veloped by Cross [130] and Carreau-Yasuda [131]. In contrast to the power-lawmodel, the latter models capture not only Newtonian behavior at low shear ratesbut also shear-thinning behavior at high shear rates [131]. The Carreau-Yasudamodel offers an additional adjustable parameter than the Cross model so that thetransfer from Newtonian to the shear-thinning behavior can be described moreprecisely [131, 132]:

η = η0(1 + (λγ)a)n−1a . (4-14)

η0 denotes the zero-shear viscosity, λ, a, and n describe parameters that are fittedto experimental data. Given common production processes for thermoplastics, theoccuring shear rates can be classified as presented in Figure 4-2.

0.1 1 10 100 1000 100000.1

1

10

100

1000

10000

Vis

cosi

ty [P

a s]

Shear rate [1/s]

ExtrusionInjectionmolding

Static press/calender/

thermoforming

Figure 4-2 Comparison of production processes for thermoplastics with regard to shear rates;redrawn from [133].

Processing of thermoplastics in a static press, double-belt press (calendering pro-cesses) or in a thermoforming unit therefore involve shear rates lower than 100 1/s.The occurring shear rates allow for the quasi-Newtonian approximation of the vis-cosity as a function of the processing temperature when time-independent behavioris assumed. The dependance on temperature is generally expressed in terms of anArrhenius relationship [131]:

η0 = η0(T0)exp(

Ea

R( 1T

− 1T0

))

, (4-15)

64 Impregnation model

where η0 denotes the reference zero-shear viscosity, Ea describes the flow activationenergy, R represents the universal gas constant.

4.3 1D through thickness model

Several models have been developed to simulate impregnation and consolidationbehavior of powder-coated tows [117, 120, 134]. Detailed geometric models thataccount for bridges formed by powder droplets during application of heat are con-sidered in these models [134].The main application of the model developed in the following is the predictionof process parameters to adjust the DOI of intermediates as desired and to sim-ulate effects of changed viscosities arising from degradation and modification. Inaddition, the model is also applied to partially impregnated tapes produced in adouble-belt press in Chapter 7. Hence, a simplified and thus more general model isdeveloped to enable predictions on the DOI for powder-coated tows as well as forconsolidated tapes.

4.3.1 Assumptions

Matrix flow is assumed to occur transverse to the fibers in thickness direction onlyas flow in width direction is prevented by either a double-belt press or molds used inthe static press process [110]. In addition, it is assumed that the fiber bed is initiallycompacted to a certain fiber volume content that does not vary significantly whenthe matrix is transferred to the molten state.

4.3.2 Model derivation

Accounting for the load sharing between fibers and matrix the governing equationsare presented in the following:

∂Vf

∂t+ ∂

∂x(Vfuf ) = 0 (4-16)

− ∂Vf

∂t+ ∂

∂x

((1 − Vf )um

)= 0 (4-17)

(1 − Vf )(uf − um) = −K

η

∂P

∂x(4-18)

The fiber volume content Vf is assumed to be constant: as soon as pressure is ap-plied the fiber bed is compressed to a certain Vf that remains constant throughout

Impregnation model 65

impregnation. In addition, the velocity of the fiber bed is assumed to be 0. Withthese assumptions, the continuity equation for the matrix (Equation 4-17) reducesto

(1 − Vf ) · ∂

∂xum = 0 (4-19)

and results inum(x) = const (4-20)

after integration. When Equation 4-20 is inserted into Darcy’s law (Eq. 4-18) thefollowing expression is obtained:

(1 − Vf ) η

Kc = ∂P

∂x. (4-21)

Further integration of Eq. 4-21 results in an expression for the matrix velocityum(x):

um(x) = K

η

ΔP

(1 − Vf )(xi+1 − xi). (4-22)

The flow front at a certain position x is expressed as a function of the initial degreeof impregnation (DOIi):

xi = xt

2 (1 − DOIi). (4-23)

The progression of the flow front is defined as:

xi+1 = xi + um · (ti+1 − ti). (4-24)

A schematic of the flow front progression according to the developed 1D throughthickness model is presented in Figure 4-3.

Matrix

pFiber

Figure 4-3 Schematic of the flow front progression according to the derived 1D through thicknessmodel.

The derived model to describe transverse flow through a unidirectional fiber bed isimplemented in MATLAB® 2014b.

66 Impregnation model

4.4 Experimental work

4.4.1 Rheology

Rheological properties can be determined by several measurement techniques suchas rotational, capillary rheometers, falling sphere viscosimeter and extensionalrheometer [135]. Viscoelastic properties are commonly determined by rotationalrheometers with a parallel-plate or cone-plate setup as shown in Figure 4-4 [136].

a) parallel-plate

SampleR

ω

R

ω

b) cone-plate

α

Figure 4-4 a) Parallel-plate and b) cone-plate fixtures for use in rotational rheometers.

Flow characteristics are determined by presetting shear stress (stress-controlledmode) or deformation rate (strain-rate controlled mode). Tests can be conducted ina step-strain mode from low to high shear rates or in an oscillatory mode. Focusingon oscillatory measurements, the acting shear force ±F , the deflection ±s andangle of twist ξ are shown between two parallel plates in Figure 4-5. Performing

ξ

s

ξ

s

+F-F

h

Shear strain Shear stress

δ

time

a) b)

Figure 4-5 a) Oscillatory measurement and b) time-delayed shift of stress response comparedto applied strain rate; redrawn from [137].

oscillatory shear tests, a sinusoidal shear strain is set as a function of time t andangular frequency ω (depending on oscillation frequency f = 2π/ω):

γ(t) = γ0sin(ωt), (4-25)

Impregnation model 67

where γ0 represents the maximum deformation amplitude. In case of visco-elasticmaterials the stress response to shear strain is time-delayed by phase lag γ (seeFigure 4-5):

τ(t) = τ0sin(ωt + δ). (4-26)

The damping behavior of visco-elastic material is expressed as

tanδ = G′′

G′ . (4-27)

G′ represents the storage modulus indicating the proportion of stored elastic energywhereas G′′ describes the loss modulus stating the proportion of stored viscousenergy. Storage and loss moduli are defined as

G′ = τ0

γ0cosδ, (4-28)

G′′ = τ0

γ0sinδ. (4-29)

The results from an oscillatory shear experiment can be expressed in term of vis-cosity:

η′ = τ0

γ0sinδ = G′′

ω, (4-30)

η′′ = τ0

γ0cosδ = G′

ω. (4-31)

The complex viscosity is then defined as

η∗ = η′(ω) − iη′′(ω). (4-32)

According to the Cox-Merz rule, an empirical relationship, the dependency of thesteady state viscosity η is equal to the frequency dependance of the complex vis-cosity η∗ [131]:

η(γ) = |η∗(ω)|. (4-33)

Rheological measurements were conducted by using a rotational rheometer of thetype TA Instruments AR 2000 ex. The samples were prepared by pressing pow-dery samples to circular-shaped pellets with 1.2 mm in thickness and a diameterof 25 mm. The pellets were heated above Tm for 90 s using a laboratory pressCollin P 200 E. Prior to rheological measurement, all pressed pellets were dried at80 °C under vacuum for at least 4 h. Disposable plate-plate sample holders madefrom aluminum were used. The measuring cell of the rheometer was heated to Tm

+ 20K. After calibration of the zero-gap position to 1 mm, the pressed pellets wereinserted, cooled to Tm and subjected to a rotational speed of 0.1 rad/s for 10 s in

68 Impregnation model

order to remove entrapped air. A deformation of 0.01 % and an angular frequencyof 10 rad/s were gauged to ensure that oscillatory measurements take place in thelinear-viscoelastic region. During cooling, a deformation of 0.00125 % and an an-gular frequency of 0.1 rad/s were adjusted.

4.4.2 Experimental determination of degree of impregnation

The duration of impregnation is governed by the transverse impregnation of singlefiber bundles [138]. This means impregnation is not completed before every fiberbundle is completely impregnated. Hence, the DOI is evaluated for a representativeamount of single fiber bundles and averaged. The DOI may be defined as the ratioof non-impregnated to impregnated fibers within a single fiber bundle, as presentedin Figure 4-6. By subtracting the non-impregnated area (Aporosity within fiber bundle)from the entire fiber bundle area (Afiber bundle) divided by the entire fiber bundlearea [12, 139] the DOI is calculated:

Degree of Impregnation (DOI) = Afiber bundle − Aporosity within fiber bundle

Afiber bundle

(4-34)

The DOIi describes the impregnation state of intermediates before processing whilethe final degree of impregnation (DOIf ) characterizes the DOI after processing ofintermediates.

Impregnated area

Porosity

Matrix

Fiber

Fiber bundle

Figure 4-6 Schematic of the degree of impregnation of thermoplastic intermediates.

Currently, there are no standardized methods to determine the DOI of intermediatematerials. Besides a sufficient level of reproducibility, a simple specimen prepara-tion and an applicability to a variety of intermediates are required for an exper-imental method to determine the DOI. Imaging methods such as microscopy orcomputer-tomography facilitate the comparison and verification of the results. In-direct methods such as the water pick-up test or mercury porosimetry allow todetermine the DOI by measuring the weight or volume that is absorbed by thenon-impregnated areas of the specimen. All methods were investigated with regardto applicability to the examined polymers and intermediates within the Bachelor’s

Impregnation model 69

thesis of Heiko Baumann [140]. Since polyamides are hygroscopic and compressiblethe water pick-up test and mercury porosimetry were found to be inapplicable.Although computer-tomography gives a three-dimensional image of the specimen,the resolution of the specimen was found to be too low due to the insufficient con-trast between fibers and matrix. In addition, computer-tomography involves longrecording times and requires high computational effort during data processing.The preparation of micrographs with subsequent post-processing was found tobe a reproducible method to compare the DOI of different specimen types. Thepost-processing is conducted by using a MATLAB® routine, developed by Bau-mann [140]. Here, the micrographs are analyzed with regard to the entire fiberbundle area and the non-impregnated area within the fiber bundle first, as shownin Figure 4-7 a,b. Subsequently, binary pictures of fibers and matrix are computed.By stretch and compress operations the impregnated and non-impregnated areasare connected as presented in Figure 4-7 c,d. Finally, the DOI is calculated accord-ing to Equation 4-34.

a)

c)

b)

d)

Figure 4-7 Micrographs of cross-sections of tapes with a) highlighted entire fiber bundle area,b) highlighted non-impregnated area and binary pictures of c) the entire fiber bundlearea and d) the non-impregnated area within the fiber bundles.

4.4.3 Design of experiments

To calibrate the model derived in section 4.3 an impregnation study is conducted forboth B3S and C2000 reinforced with polyamide-sized carbon fibers (CF-TP). Theimpregnation study is designed and evaluated by design of experiments (DOE).DOE has established as method to plan and analyze experiments in an efficientmanner. Linear models fail as soon as the relationship between factor (input vari-able) and response becomes non-linear. Non-linear relationships may be describedby second-order designs. In 1960, Box and Behnken [141] introduced a novel classof incomplete three-level factorial designs that are known as Box-Behnken Designs.Considering the experiment space as a cube, the experiments of a Box-Behnken

70 Impregnation model

Design are placed in the middle of each cube edge, as shown in Figure 4-8 [142]. Inparticular with regard to non-linear relationships, the Box-Behnken Design allowsfor the determination of the optimum in between two extreme levels of a factor.In case of a full factorial design 27 experiments would be necessary for the presentexperimental design. By using the Box-Behnken Design, the amount of requiredexperiments reduces to 13 as the basic points of the cube are excluded. In addi-tion, the center point has to be repeated twice to guarantee orthogonality of theexperimental design. Thus, the Box-Behnken Design results in 15 experiments.

Factor A

Factor B

Factor C

-1

0

1

0 1-1

0

1

-1

Figure 4-8 Graphical representation of a Box-Behnken Design with three factors having twoextreme factor levels (−1 (minimum) and 1 (maximum)) and a center point (0);redrawn from [142].

The chosen factors to conduct the impregnation study represent the most dominantprocess parameters for impregnation: time, temperature and pressure. Two extremelevels were chosen for each factor and the center points were replicated twice. TheBox-Behnken Designs for both CF-TP/B3S and CF-TP/C2000 are presented inTable 4-2.

Table 4-2 Experimental design with three factors - time, temperature and pressure - for CF-TP/B3S and CF-TP/C2000.

Factor Factor levelsCF-TP/B3S CF-TP/C2000

Time [min] 1 5.5 10 1 3.5 6Temperature [°C] 260 270 280 280 290 300Pressure [bar] 5 17.5 30 5 17.5 30

For each material combination, the Box-Behnken Design with 15 experiments wascarried out separately in a randomized order and is stated in the appendix (section

Impregnation model 71

A.1). 12 plies of powder-coated tows made of CF-TP/B3S and CF-TP/B3S werestacked and pressed into unidirectional laminates with 100 mm in length (alongfiber direction) and 50 mm in width. Using a laboratory press Collin P 200 E, thestacks were inserted into the press unit, heated to the desired temperature. Afterthe ply stack has reached the temperature, pressure was applied. In addition, sometest panels produced according to the previously described procedure were immedi-ately cooled to room temperature after the applied pressure has built up. These testpanels (referred to as “0 min”-tests) serve to determine the impregnation progressof the test panels during the heating period until the desired process temperaturehas been reached. Micrographs are prepared from each laminate to evaluate theeffect of the corresponding process parameter set by the achieved DOI.The factor levels for time were selected to enable visual inspection of the impreg-nation progress. Narrow intervals were chosen for the factor temperature as theupper and lower boundary has been evaluated by rheology in earlier experiments.The extreme factor levels for pressure, 5 and 30 bar, are chosen to provide sufficientcontact between the single plies and to avoid fiber damage.

4.4.4 Interlaminar shear test

The impregnation study also aims to find the process parameter with the mostsignificant impact on mechanical properties. Besides micrographs, test specimensto determine the interlaminar shear strength (ILSS) are extracted from the testpanels produced during the impregnation study (Table 4-2). The test specimenscut to 10 mm in width, 20 mm in length and revealed a thickness of 2 mm.The interlaminar shear strength ILSS was determined according to DIN EN ISO14130 [143] as follows:

ILSS = 34

P

wt, (4-35)

where P represents the maximum force, w and t denote the specimen’s width andthickness, respectively.The used test setup represented a three-point bending design comprising of aloading nose with a radius of 5±0.2 mm and supporting noses with a radius of2±0.2 mm. The support span L was adjusted to L = 5t.

4.5 Results

The viscosity curves for both aliphatic polyamides and a co-polyamide were de-termined to derive the zero-shear viscosity that is used as input data for the de-rived model. Representative micrographs are depicted to visualize the impregnation

72 Impregnation model

progress. After the evaluation of the DOI from all test panels, the trends of themain effects and the interactions were analyzed by using the statistical softwaretool MINITAB 17. The experimental results are compared to values computed byusing the derived model. In addition, the impact of different process parameters onDOI and ILSS is correlated.

4.5.1 Viscosity data

Oscillation measurements yield the complex viscosity η∗ and the corresponding vis-cosity η is related to the shear rate by applying the Cox-Merz rule. The measure-ments were taken at three different temperatures: 260, 270 and 280 °C for aliphaticpolyamides B3S, B3L and B40; 290, 300 and 310 °C for the co-polyamide C2000.B40 and B3L were added for means of comparison. The viscosity curves for B3S,B3L, B40 and C2000 are presented in Figure 4-9.

10 100 1000

100

1000

10000

10 100 1000

100

1000

10000

Vis

cosi

tyη

[Pas

]

Shear rate [1/s]

C2000_290°CC2000_300°CC2000_310°C

a

B40_260°CB40_270°CB40_280°C

B3L_260°CB3L_270°CB3L_280°C

Vis

cosi

ty η

[Pas

]

Shear rate [1/s]

B3S_260°CB3S_270°CB3S_280°C

b

Figure 4-9 Viscosity curves at different temperatures for a) B3S, B3L, B40 and b) C2000.

The high-molecular weight PA6 grade, B40, showed the highest zero viscosity η0 incomparison to B3S, B3L and C2000. The viscosity curves for B3S and B3L resembleeach other although B3L is a PA6 type with elastomeric particles for toughening.C2000 shows higher viscosity values than the aliphatic B3S or B3L. For shear ratesbeginning from 100 1/s, all investigated polymers show pronounced shear-thinningbehavior. The viscosity curves of all polymers reveal a strong dependency on tem-perature. To describe the temperature-dependent behavior according to Arrhenius,the activation energy for flow, Ea, is determined first, by plotting the temperatureshift factor aT against the inverse of the product of universal gas constant and

Impregnation model 73

temperature 1/RT . By obtaining η0 from the viscosity curves in Figure 4-9, thetemperature shift factor aT is defined as [137]

aT = η0(T )η0(Tref ) . (4-36)

The slope of the straight plots of aT against 1/RT yields the flow activation energyEa as shown in Figure 4-10.

0.216 0.220 0.224

0.5

1

1.5

2

0.39 0.40 0.41 0.42

0.5

1

1.5

2

B3SB3LB40Linear fitLinear fitLinear fit

a T=

η 0(T)/

η0(

T 0)[-]

1/RT [mol K-1 kJ-1]

Ea = 39.61 kJ/mol (B3L)Ea = 36.67 kJ/mol (B3S)Ea = 20.74 kJ/mol (B40)

a

C2000Linear fit

a T=

η 0(T)/

η0(

T 0)[-]

1/RT [mol K-1 kJ-1]

Ea = 10.58 kJ/molb

Figure 4-10 Plots of temperature shift factor aT versus 1/RT for a) B3S, B3L, B40 and b)C2000; the slope of the lines represents the activation energy Ea.

By using the determined values for Ea, the viscosity can be calculated accordingto Eq. 4-15 in dependance of any temperature profile as presented in Figure 4-11.

200 400 600 800 1000

50

100

150

200

250

300

Temperature profile during thermoformingCalculated viscosity (Arrhenius) for B3S

Time [s]

Tem

pera

ture

[°C

]

100

1000

10000

η(T)

[Pa

s]

Figure 4-11 Viscosity curve of B3S calculated according to Arrhenius for a typical temperatureprofile during thermoforming.

74 Impregnation model

The used zero-viscosity data determined by rheology experiments are shown inTable 4-3 and Table 4-4.

Table 4-3 Zero-shear viscosity data for B3S used as input parameter.

Polymer η0 [Pa s] at260 °C 270 °C 280 °C

B3S 201.93±4.30 170.49±8.16 101.19±5.48

Table 4-4 Zero-shear viscosity data for C2000 used as input parameter.

Polymer η0 [Pa s] at280 °C 290 °C 300 °C

C2000 474.69±18.54 310.11±5.60 235.95±34.17

4.5.2 Influence of processing on impregnation progress

To visualize the influence of processing, selected micrographs (see subsection 2.3.4)are compared with regard to the impregnation state for extreme levels of one factorat a time. To begin with, the effect of the factor time on the DOI is compared forCF-TP/B3S in Figure 4-12. Extending the press time from 1 minute to 10 minutes,the DOI increases from 76 % to 100 %.

1 min/17.5 bar/260°C 10 min/17.5 bar/260°C

a) b)

~70 μm~100 μm

Figure 4-12 Micrographs of CF-TP/B3S test panels after a press time of a) 1 minute revealingnon-impregnated fiber bundles and b) 10 minutes showing complete impregnation.

The temperature appears to have a less pronounced influence on impregnation aspresented in Figure 4-13. When the temperature is raised from 260 to 280 °C whilethe pressure is held constant at 5 bar for 5.5 minutes, the DOI is increased from84 to 87 %. However, the selected extreme levels for the temperature are within

Impregnation model 75

a narrow range while the extreme levels for the pressure or the time yield muchlarger differences and hence more pronounced effects.

5.5 min/5 bar/260°C 5.5 min/5 bar/280°C

a) b)

~50 μm~70 μm

Figure 4-13 Micrographs of CF-TP/B3S test panels pressed with a) 260 °C showing non-impregnated fiber bundles and b) 280 °C revealing an advanced impregnationprogress.

Increasing the pressure from 5 to 30 bar has also a less pronounced effect on theimpregnation progress. The DOI is raised from 76 to 86 % when CF-TP/B3S wereprocessed at 270 °C for 1 minute as shown in Figure 4-14.

1 min/5 bar/270°C 1 min/30 bar/270°C

a) b)

~60 μm ~60 μm

Figure 4-14 Micrographs of CF-TP/B3S test panels pressed with a) 5 bar revealing non-impregnated fiber bundles and b) 30 bar with increased DOI.

Micrographs of CF-TP/C2000 reveal similar dependencies on process parametersas CF-TP/B3S.A more abstract representation of the relationship between process parametersand achieved DOI is enabled by considering the main effects and their interactions,analyzed by using the statistical software MINITAB 17. Figure 4-15 presents themain effects of time, temperature, and pressure on the development of DOI forCF-TP/B3S. The main effects plot reveal non-linear relationships between processparameters and DOI. For the factor time a DOI of more than 100 % is yielded.This is attributed to the insufficient least-square fit conducted for Box-BehnkenDesigns by using MINITAB 17. All DOI obtained from micrographs yielded valuesequal or below 100 %. Considering the maximum for each factor, the optimumprocess setting for CF-TP/B3S is found at a dwell time of 8 minutes at 275 °C

76 Impregnation model

Mea

nof

DO

I [%

]Pressure [bar] Temperature [°C] Time [min]

10 20 30 260 280 300 5 10

105.0

100.0

95.0

90.0

85.0

Figure 4-15 Main effect plot for DOI of CF-TP/B3S.

and 25 bar. The interaction plot presented in Figure 4-16 illustrates how a factorcan influence another one. With the three main factors - time, temperature andpressure - six possible interactions occur. The strongest interaction for CF-TP/B3Scan be identified for time on pressure and temperature as presented in Figure 4-16.

260 270 280

100

90

80

100

90

80

Mea

nof

DO

I [%

]

Pressure x Temp.

Pressure x Time Temperature x Time

Temp. x Pressure Time x Pressure

Time x Temperature

Pressure [bar] Temperature [°C] Time [min]10 20 30 3 6 9

Pressure [bar]

Temperature [°C]

Time [min]

517.5

30

15.510

260270280

100

90

80

Figure 4-16 Interaction plot for DOI of CF-TP/B3S.

The influence of both temperature and pressure on time as well as the effect ofpressure on time is found to be similar and less pronounced. Temperature appearsto have a small effect on pressure as can be derived from the almost horizontal

Impregnation model 77

interaction plot.In contrast to CF-TP/B3S, the main effect plot for CF-TP/C2000 in Figure 4-17shows a progressive course for the temperature without an optimum. The main

Mea

nof

DO

I [%

]

100.0

97.5

95.0

92.5

90.0

87.5

85.0

Pressure [bar] Temperature [°C] Time [min]

10 20 30 280 290 300 2 4 6

Figure 4-17 Main effect plot for DOI of CF-TP/C2000.

effect plot for pressure and time however reveals a maximum. This leads to anoptimum process setting for CF-TP/C2000 that is found at a dwell time of at least6 minutes at 300 °C and 25 bar. Compared to CF-TP/B3S, the effect of pressure ontemperature and time is more pronounced than the effect of time on pressure andtemperature for CF-TP/C2000 (see Figure 4-18).

Pressure x Temp.

Pressure x Time Temperature x Time

Temp. x Pressure Time x Pressure

Time x Temperature

100

90

80

Mea

nof

DO

I [%

]

Pressure [bar] Temperature [°C] Time [min]

100

90

80

100

90

80

10 20 30 2 4 6

Pressure [bar]5

17.530

Temperature [°C]

Time [min]1

3.56

280290300

280 290 300

Figure 4-18 Interaction plot for DOI of CF-TP/C2000.

78 Impregnation model

4.5.3 Model calibration

The best agreement between model and experiments was found by using the com-paction model (Eq. 4-11) and the modified equation for permeability according toGutowski (Eq. 4-6). Plotting the experimentally determined Vf as a function of theapplied pressure yields the compaction curve according to Eq. 4-11 that is shownin Figure 4-19.

a) b)

Figure 4-19 Development of the experimentally determined Vf (◦) as a function of appliedpressure for a) CF-TP/B3S and b) CF-TP/C2000.

For CF-TP/B3S, the empirical constant As was approximated by a least-square-fit and is found to be close to the value obtained by Gutowski et al. [127] forpoorly aligned fibers. When fitting As to the experimental data for CF-TP/C2000,it differs from literature by one order of magnitude which is attributed to localdeviations of Vf .All input parameters used to simulate the DOI that developed during the impreg-nation study are summarized in Table 4-5.

Table 4-5 Input parameters for model calibration.

Parameter Impregnation studyFiber radius rf [m] 0.0000035Initial Vf [-] 0.50Initial fiber bundle height [m] 0.00026Time increment [-] 1Pressure increment [-] 0.5Modified Kozeny constant k′

zz [-] 0.20Maximum fiber volume content V ′

a [-] 0.82Empirical constant As [Pa] 772.77Maximum fiber volume content Va [-] 0.892

Impregnation model 79

The Kozeny constant kzz = 0.2 along with the maximum possible fiber volumecontent V ′

a was taken from literature [127] and indicates poor fiber alignment. Inreality, the maximum fiber volume content Va lies between quadratic and hexagonalfiber packing and is set to 0.892 in accordance with [127]. The initial fiber bundleheight was determined to 260μm from micrographs.The agreement between model and experiment is compared for CF-TP/B3S andCF-TP/C2000 at a pressure of 17.5 bar in Figure 4-20.

a)

b)

c)

d)

e)

f)

Figure 4-20 Comparison of experimentally determined (◦) and calculated DOI for various tem-peratures at 17.5 bar for a-c) CF-TP/B3S and d-f) CF-TP/C2000.

A good agreement between model and experiment is found for CF-TP/B3S whenpressed at 260 °C and 270 °C. A slight deviation is observed when pressed at 280 °C.The first data point at t=0 min is obtained from the separately pressed “0 min” testpanels. The error is assumed to be caused by an inherent variability of the fiberdistribution for the different test panels.For CF-TP/C2000, model and experimental results from the impregnation studyshow good agreement when As is set to 772.77 Pa, as determined for CF-TP/B3S.The experimentally obtained Vf by microscopy to fit As are assumed to be lessaccurate than for CF-TP/B3S and lead to the significantly lower As. In addition,

80 Impregnation model

slight deviations between model and experiments can be observed, in particular forthe DOI from test panels pressed for 1 minute. This may be caused by deviationsin the quality of the micrographs.Altogether, a good agreement between the derived model and the impregnationstudy is found for both material combinations. The overall error between modeland experiment is 4.34 % for CF-TP/B3S and 5.42 % for CF-TP/C2000.

4.5.4 Influence of degree of impregnation on interlaminar shearstrength

In addition to the calibration of the developed model, the impregnation studyserved to identify a potential correlation between the DOI and ILSS. All testedspecimens showed plastic deformation instead of interlaminar shear failure as oftenobserved when testing CFRTP. According to Section 9.7d in [143] the yielded me-chanical property cannot be indicated as interlaminar shear strength but is allowedto be compared for specimens made of the same materials with comparable failurebehavior. As all specimens showed a similar failure, comparison of the yielded me-chanical property was allowed. Figure 4-21 plots the ILSS and the DOI at variousdwell times, temperatures and pressures as used during the impregnation study.With exception of the results obtained at 260 °C, 30 bar and 5.5 min, a correlationbetween ILSS and DOI of CF-TP/B3S can be found (Figure 4-21 a, b and c).In case of CF-TP/C2000, the same tendency is found for the development of ILSSand DOI (Figure 4-21 f). Minor deviations are observed for CF-TP/C2000 pressedat 300 °C, 5 bar and 3.5 min (Figure 4-21 b) as well as at 290 °C, 30 bar and 1 min(Figure 4-21 d). In general, the ILSS correlates well with the obtained DOI. Thistesting method is therefore found to be suitable to reflect the state of impregnationof thermoplastic composites.

4.6 Conclusion and implications

A model was developed to simulate the impregnation progress of various interme-diates. In a first step, the transverse resin flow of powder-coated tows was inves-tigated. Several processing phenomena of the individual constituents - fiber andmatrix - were taken into account. The actual resin flow was modeled by usingDarcy’s law. The impregnation characteristics of the used matrix systems in termsof rheology were determined. Additionally, the temperature dependance of the vis-cosity was modeled by the Arrhenius relationship to account for viscosity changesas a function of the applied temperature.

Impregnation model 81

5 bar/260°C

5 bar/280°C

30 bar/260°C

30 bar/280°C

20

30

40

50

60

70

80

90

100

B3S; ILSS; t=5.5 min

ILS

S[M

Pa]

a

20

30

40

50

60

70

80

90

100

B3S; DOI; t=5.5 min

DO

I[%

]

5 bar/280°C

5 bar/300°C

30 bar/280°C

30 bar/300°C

40

50

60

70

80

90

100

C2000; ILSS; t=3.5 min

ILS

S[M

Pa]

b

40

50

60

70

80

90

100

C2000; DOI; t=3.5 min

DO

I[%

]

5 bar/1 min

5 bar/10 min

30 bar/1 min

30 bar/10 min

20

30

40

50

60

70

80

90

100

B3S; ILSS; T=270°C

ILS

S[M

Pa]

c

20

30

40

50

60

70

80

90

100

B3S; DOI; T=270°C

DO

I[%

]

5 bar/1min

5 bar/6 min

30 bar/1min

30 bar/6 min

40

50

60

70

80

90

100

C2000; ILSS; T=290°CIL

SS

[MP

a]

d

40

50

60

70

80

90

100

C2000; DOI; T=290°C

DO

I[%

]260°C/1 min

260°C/10 min

280°C/1 min

280°C/10 min

20

30

40

50

60

70

80

90

100

B3S; ILSS; p=17.5 bar

ILS

S[M

Pa]

e

20

30

40

50

60

70

80

90

100

B3S; DOI; p=17.5 bar

DO

I[%

]

280°C/1min

280°C/6min

300°C/1min

300°C/6min

40

50

60

70

80

90

100

C2000; ILSS; p=17.5bar

ILS

S[M

Pa]

f

40

50

60

70

80

90

100

C2000; DOI; p=17.5bar

DO

I[%

]

Figure 4-21 ILSS and the DOI for a) B3S at constant time, b) C2000 at constant time, c) B3Sat constant temperature, d) C2000 at constant temperature, e) B3S at constantpressure, f) C2000 at constant pressure.

To verify the derived model experimentally, an impregnation study was conductedby varying the three main process parameters that drive impregnation: time, tem-perature and pressure. The experimental design followed the Box-Behnken method,a specialized DOE design.By post-processing micrographs with a MATLAB® routine, an experimental pro-cedure was developed to determine the DOI. The DOI was obtained for all ex-periments of the impregnation study and served as response to analyze the DOE.A stronger sensitivity of the DOI on pressure was found for CF-TP/C2000 thanfor CF-TP/B3S. In addition, impregnation is completed earlier for CF-TP/C2000

82 Impregnation model

despite the higher melt viscosity of C2000.In a subsequent step, the results obtained by the model were compared to experi-mental results from the impregnation study. The model showed good agreement toexperimental data for both CF-TP/B3S and CF-TP/C2000. The overall error wasabout 5 %.In addition to micrographs, test specimens were extracted from the panels preparedfor the impregnation study to determine ILSS. The testing method was found suit-able to provide information about the impregnation progress of CF-TP/B3S aswell as CF-TP/C2000 since ILSS and DOI showed similar tendencies for variousprocess parameter settings.For the following investigations within this work, the developed model can beused to determine the required processing parameters to complete impregnationof powder-coated tows and tapes in dependance of the DOI. The model can incor-porate the initial DOI and predict process parameters to complete impregnation ofpowder-coated tows or consolidated tapes. In addition, the model allows to investi-gate the effects of viscosity changes arising from applied temperature, degradationor modification by additives on the impregnation time.

5 Degradation of polyamidesAs presented in Chapter 1, common production processes for thermoplastic com-posites are characterized by repeated heating cylces from powder-coating to tapeor laminate consolidation until thermoforming to a final component. Degradationreactions arising from repeated melting may influence the complex viscosity andaffect the impregnation progress. Two temperature profiles are derived from typ-ical CFRTP production processes to simulate the thermal stresses acting on B3Sand C2000 during processing. Both polymers are subjected to thermal cycling byusing differential scanning calorimetry (DSC), thermogravimetric analysis (TGA)and rheometry. Gel permeation chromatography (GPC) serves to detect changesin the molecular weight distribution induced by thermal cycling. Besides the es-tablishment of an understanding of degradation reactions that occur during theprocessing of thermoplastic composites, it is further evaluated how degradationreactions may influence gradual impregnation throughout production to define asuitable processing window. Parts of the following section have been previouslypublished in [144].

5.1 Literature review

5.1.1 Thermal degradation

Under exclusion of oxygen, primary scission of polycaprolactam types is assumedto occur by re-equilibration to monomeric or oligomeric cyclic products. Many re-searchers found the scission of the macromolecules to mainly take place at the pep-tide bond or at bonds nearby [145]. Besides primary scission of the macromoleculesat the peptide bond or nearby, acidolysis or aminolysis are potential reaction mech-anisms during thermal degradation that can occur intra- or intermolecular [146].Secondary reactions occur especially at temperatures above 300 °C and were foundto be primarily responsible for crosslinking of aliphatic polyamides [145, 147, 148].The main mechanism behind the formation of branched structures is hydrolysis ofpeptide bonds and further condensation of carboxyl and amine end groups [148].The principal scheme is presented in Figure 5-1.

5.1.2 Thermo-oxidative degradation

In presence of oxygen, thermo-oxidative degradation is promoted when polyamidesare processed at high temperatures for long time and under shear stress during

83

84 Degradation of polyamides

H2O

CH2 CH (CH2)3 CONHCONH (CH2)5 CONH2

CONH (CH2)5 C N CONH (CH2)5 N C O CH3 (CH2)4 CONH

(CH2)5 C NH CH2 (CH2)4 C NH

O O

C NH (CH2)4 C N

HO

CH2 (CH2)4 C NH

O

-H+H

Figure 5-1 Principle scheme of thermal decomposition of PA6 [145].

e.g. molding or extrusion [149]. Absolute exclusion of oxygen during processing isnot feasible due to entrapped air or oxygen dissolved in the amorphous regions ofthe polymer [149, 150]. Marechal et al. [151] found the oxygen concentration in thepolymer melt to be twenty times higher when compared to solid polymer due toincreased molecular mobility along with enhanced solubility in the molten state.First studies [152] on the oxidation of polyamides led to the differentiation into thefollowing three main reactions:

1. Formation of N-acylamides (imides):

R CO NH CH2 R’ R CO NH CO R’ (5-1)

2. Formation of N-formamides resulting from C1-C2 scission:

R CO NH CH2 R’ R CO NH CHO (5-2)

3. Oxidative dealkylation to yield carbonyl compounds:

R CO NH CH2 R’ R CO NH2 + R’ CHO (5-3)

The commonly accepted mechanism for chain scission of aliphatic polyamides dueto thermo-oxidative degradation, introduced by Levantovskaya et al. [152], is ini-tiated by the formation of radicals [151, 153–155] as shown in Figure 5-2. Theformation of radicals may occur during processing at high temperatures and under

Degradation of polyamides 85

C NH CH2

O

C NH CH

O

C NH C

O O O

O2

C NH C

O O

C NH CH

O OHO2

C NH CH2

O

C NH CH

O OHC NH CH2

O

C NH C

O O O

C NH CH

O O

OH

Figure 5-2 Basic mechanism for chain scission during oxidation of aliphatic polyamides [156].

shear stresses. The formed radicals further attack the N-vicinal methylene groupand form more radicals for another radical chain mechanism. Besides the primaryattack of the N-vicinal methylene group, Allen et al. [157] investigated that anyother methylene group can be oxidized, preferably β -positioned groups. In addi-tion to the common thermo-oxidative mechanism, alkoxy radicals may also formunstable hydroxyl compounds that decompose to aldehydes and primary aminesas suggested by Lánská et al. [153].During thermo-oxidative degradation of PA6 chain scission can be superposed bythe formation of crosslinks. Analogous to thermal degradation, there are secondaryreaction mechanisms leading to branching or crosslinking.

5.1.3 Post-condensation

In addition to thermal and thermo-oxidative degradation reactions, PA6 tends tore-equilibrate with increasing temperature. New macromolecules are formed bypost-condensation of amine and carboxylic chain ends that are produced by chainscission of the molecules. In presence of water concentrations of up to 1.5 % [158],the activity of the chain ends increases and promotes polycondensation reactions.The molecular weight equilibrium is determined by an equilibrium between thechain ends, the amide groups and the remaining water concentration as Eq. 5-4shows [151]:

COOH + NH2 NHCO + H2O (5-4)

86 Degradation of polyamides

5.1.4 Influence of degradation on processing

Earlier studies refer to the thermal and thermo-oxidative degradation of neataliphatic nylons at temperatures below melting temperature Tm [156, 159, 160], veryhigh temperatures close to decomposition [145] or above for pyrolysis [161, 162]. Inaddition, the thermal and thermo-oxidative degradation behavior is mainly studiedwith focus on the theoretical development of degradation mechanisms [163, 164]and on long-term effects on the polymer properties by accelerated aging pro-cesses [165, 166].Direct effects of degradation reactions that occur during processing of thermoplas-tics and influence further processing were investigated to a limited extent. La Man-tia et al. [167] found the Newtonian viscosity decreasing with increasing passagesin a single-screw extruder during recycling of wet and dry PA6. The results fromgel permeation chromatography (GPC) reported in [168] show that the molecularweight of PA6 continuously decreases when repeatedly processed under oxidativeatmosphere while the molecular weight distribution (MWD) broadens at the sametime.During preliminary studies, substantial and irreversible increases in viscosity ofboth neat B3S and C2000 were detected when exposed to long dwell times atelevated temperatures. Similar observations were made by Khanna et al. [169] de-tecting sharp increases of the zero shear melt viscosity of PA6 that is assumed toarise from re-equilibration towards higher molar masses in presence of moisture.Reductions in solution viscosity of PA6 at temperatures below Tm were also re-ported by Dong and Gijsman [156]. A decreasing melt fluidity index (MFI) withincreasing dwell times at elevated temperatures due to thermo-oxidative degrada-tion was observed by Pakharenko et al. [170]. In contrast to the previous studies,La Mantia et al. [167] detected decreasing Newtonian viscosity compared to virginPA6. Su et al. [168] observed reductions in melt viscosity with reduced molecularweights.The ambiguous results from literature led to a comprehensive study of the meltviscosity as a function of the temperature profiles used during the manufacture ofCFRTP components. Additional methods such as DSC, TGA and GPC were ap-plied to find potential correlations between degradation and viscosity changes. Asubstantial increase in viscosity has considerable effect on processing intermediatesto CFRTP parts by slowing down the impregnation progress. Insufficient impregna-tion involving entrapped air leads to stress concentrations due to discontinued loadtransfer from fiber to fiber and thus reduces the mechanical performance. For theconcept of gradual impregnation, the comprehension of changes in viscosity overtemperature and time is crucial to ensure processability. The aim of this chapteris therefore to identify a processing window being suitable for impregnation.

Degradation of polyamides 87

5.2 Temperature profiles

Both polyamide types were subjected to two different temperature profiles to de-tect changes arising from repeated heating cycles as present during production ofCFRTP. The profiles follow a production process for CFRTP starting with powder-coating carbon fibers on both sides, consolidation to tapes in a double-belt press,consolidation to multi-ply laminates and thermoforming of multi-ply laminates toa final component as shown in Figure 5-3.

Powder-coating Tape consolidation

Double-belt pressPowder reservoir

Spread carbon fiber tows

Laminate production

Infrared heat source

Infrared heat source Static

press

Thermoforming

Consolidated tape

Multi-ply laminate

Heat

3D Tool

Powder-coated tows

Final component

Figure 5-3 Considered CFRTP production process to derive temperature profiles.

Altogether, the production of components made from CFRTP involves four heatingprocesses above Tg or Tm of the polymer due to impregnation, consolidation andforming. During repeated heating processes (thermal cycling) thermoplastics aresubjected to thermal and thermo-oxidative degradation.Temperature profile P1 in Figure 5-4 incorporates the processing temperatures re-quired for CFRTP production with a constant dwell time of 5 min. This profileis used to identify potential effects on the degradation by repeated heating only.Temperature profile P2 in Figure 5-4 includes prolonged dwell times as commonlyused in CFRTP production. A processing time of 10 min is considered for tapeconsolidation and thermoforming. This is an averaged time to account for consoli-dation and thermoforming of UD spread tows as well as for woven fabrics or NCF.Extreme processing times of 60 min for B3S and 30 min for C2000 were selectedfor the production of laminates from several tape plies in a press. This accountsfor prolonged impregnation times of woven fabrics where large flow paths have tobe overcome. In addition, the selected processing times allow for the observationof the viscosity development as a function of time under temperature exposure.

88 Degradation of polyamides

0 40 80 120 160 200 2400

70

140

210

280

350

420

0 40 80 120 160 200 2400

70

140

210

280

350

420

Tem

pera

ture

[°C

]

Time [min]

Temperature profile P1 (B3S)Temperature profile P2 (B3S)

Powder-Coating

TapeConsolidation

LaminateProduction

Thermoforminga

Tem

pera

ture

[°C

]

Time [min]

Temperature profile P1 (C2000)Temperature profile P2 (C2000)

Powder-Coating

TapeConsolidation

LaminateProduction

Thermoformingb

Figure 5-4 Temperature profiles P1 and P2 for a) B3S and ) C2000 derived from a CFRTPproduction process.

5.3 Experimental methods

The effects of thermal cycling on B3S and C2000 were generated and analyzed bysubjecting samples to both temperature profiles directly in a differential scanningcalorimeter, thermogravimetric analyzer and rheometer. Changes in the chemicalcomposition were analyzed by GPC after exposure to both temperature profiles ina differential scanning calorimeter.Parts of the experimental work have been conducted within the framework of theMaster’s Thesis from Christian Heckel [171]. Access to the central laboratories ofSGL Carbon GmbH to use DSC, TGA and rheometry within this Master’s Thesisis gratefully acknowledged.

5.3.1 Sample preparation

As-received B3S and C2000 pellets were cryogenically milled to powder by means ofliquid nitrogen. This was carried out by SGL Carbon GmbH in collaboration withA. Schulman. Due to centrifugal forces, the polymer pellets hit a circular sieve witha size of hole of 400 μm and size-reduce the polymer pellets. Powder with a particlesize distribution of 200 μm by majority (more than 80 %) was produced. Samplesin powder form account for the initial state of the polymers at the beginning toproduce CFRTP components from powder-coated tows (see Figure 5-3).

5.3.2 Differential scanning calorimetry

Using a TA Instruments Q100, B3S and C2000 in powder form were inserted in thecalorimeter and subjected to temperature profile P1 and P2 under air as well as

Degradation of polyamides 89

nitrogen atmosphere. A heating rate of 10 °C/min and a cooling rate of 20 °C/minwas applied for each process step. A cooling rate of 20 °C/min is commonly ap-plied during processing in a static press. Large cooling rates as present duringpowder-coating and thermoforming are not considered to suppress the influence oncrystallization. Thus, the effects of thermal cycling can be attributed to degrada-tion reactions only.By using DSC the development of Tm was analyzed arising from changes in themolecular composition and induced by thermal cycling. The evaluation was carriedout by using the software Universal V4.5A TA Instruments. Bimodal peaks wereobserved for C2000 that were evaluated by laying tangents alongside both peaksas shown in Figure 5-5.

210 220 230 240 250 260 270 280 290 300

-14

-12

-10

-8

-6

-4

Hea

tFlo

w[m

W]

Temperature [°C]

Tangents

Extrapolated Tm

Figure 5-5 Determination of Tm for bimodal melt peaks as present for C2000.

5.3.3 Thermogravimetric analysis

By using a thermogravimetric analyzer equipped with a thermobalance changes inmass of less than 10 μg can be detected as a function of time (static) and temper-ature (dynamic). TGA is a common method to measure the thermal stability andthe amount of single constituents in blends or compounds [76, 172].Measurements were performed on samples in powder form by using a TA Instru-ments Q500. All samples were subjected to both temperature profiles and atmo-spheres with a heating rate of 10 °C/min and cooling rate of 20 °C/min directly inthe thermogravimetric analyzer. However, the thermogravimetric analyzer does notprovide active cooling leading to a decrease of the actual cooling rate to 10 °C/minduring the measurements.The total mass loss compared to the original sample weight was determined sepa-

90 Degradation of polyamides

rately from thermograms for each process step of the temperature profiles by usingthe software Universal V4.5A TA Instruments. For each of the process steps frompowder-coating until thermoforming a new specimen was used resulting in eightsamples per polymer for temperature profile P1 and P2. The initial mass loss upto 140 °C was subtracted as it is attributed to water evaporation. Considering thethermogram in Figure 5-6, the overall mass loss of PA6 for the thermoforming stepafter exposure to temperature profile P2 yields 2.22 % neglecting the initial massloss of 2.10 %.

0 50 100 150 20095

96

97

98

99

100

0.14 %(0.03 mg)

1.35 %(0.26 mg)

0.28 %(0.05 mg)

0.45 %(0.09 mg)

B3S_P2_AirTemperature

Time [min]

Wei

ght[

%]

2.10 %(0.41 mg)

50

100

150

200

250

Tem

pera

ture

[°C

]

Figure 5-6 Procedure to determine the mass loss that has occurred during every process step ofP1 and P2 for B3S as an example.

5.3.4 Rheometry

Rheological measurements were conducted by using a TA Instruments AR 2000 exand TA Instruments DHR-3. The same experimental procedure including specimenpreparation was followed as described in subsection 4.4.1. During the rheologicalmeasurements the samples were cooled to 180 °C instead of 40 °C as depicted inFigure 5-4 to ensure operability of the viscosity experiments while reducing thethermal stresses by operating below Tm.

5.3.5 Gel permeation chromatography

Gel permeation chromatography/size exclusion chromatography (GPC/SEC) is aliquid column chromatographic technique that enables molecule separation in so-lution. By sorting molecules according to their size the number-average molecularweight (Mn), the weight-average molecular weight (Mw), and the molecular weight

Degradation of polyamides 91

distribution (MWD) can be determined. The material to be investigated is solvedand introduced to columns with porous packings made of semirigid organic gelsor rigid solids representing the stationary phase. The molecules are transportedthrough the column by the solvent which is the (liquid) mobile phase. The surfaceand the pores of the column packing temporarily restrain some dissolved moleculesfrom further migration through the column. Thus, the different velocities of thedissolved molecules to migrate through the column result in a time-delayed elutionfrom the column [173]. The basic principle of GPC measurements is sketched inFigure 5-7.

Retention time

PorousPacking

(A) (B) (C) (D)

Chromatogram

(A) Injection (C) Large solutes eluted

(B) Size separation

(D) Small solutes eluted

Concentrationdetector

Figure 5-7 Size separation and detection of dissolved molecules by GPC; redrawn from [173].

Modern GPC/SEC devices are equipped with physical and chemical detectors.Concentration detectors such as a differential refractometer can be combined withphysical detectors such as light-scattering photometer and/or viscosimeter. Theobtained elution curves are transformed into MWD curves by calibrating it withstandard polymers of known weight. By comparison of the investigated solute poly-mer with a calibration standard, Mn and Mw are generated [173].In cooperation with the WACKER-Chair of Macromolecular Chemistry from theTechnical University of Munich, GPC/SEC measurements were conducted to de-termine Mn, Mw, and MWD of samples subjected to thermal cycling. All testedsamples in powder form were dried and dissolved in hexafluoroisopropanol (HFIP)with a sample concentration of 1 mg/ml. The used Agilent 1200 Series with aHFIP gel column was calibrated with polymethylmethacrylate (PMMA) stan-dards for narrow MWD. The mobile phase ran at a flow rate of 0.5 ml/min ata temperature of 40 °C. All GPC measurements were evaluated with the softwarePSS WinGPC software.

92 Degradation of polyamides

5.4 Results

5.4.1 Influence on melting temperature

Since the most significant changes in Tm occurred during the exposure of B3S andC2000 to temperature profile P2 in air, the DSC thermograms for this profile areshown in Figure 5-8.

160 180 200 220 240 90 120 150 180 210 240 270

Thermoforming

Laminate Production

Tape Consolidation

Powder-Coating

Hea

tFlo

w[m

W]

Temperature [°C]

a

Hea

tFlo

w[m

W]

Temperature [°C]

Thermoforming

Laminate Production

Tape Consolidation

Powder-Coating

Cold crystallization peak

b

Figure 5-8 DSC thermograms for a) B3S and b) C2000 subjected to temperature profile P2 inair.

The development of Tm of B3S and C2000 during exposure to both profiles ispresented in Figure 5-9.

218

220

222

224

226

240

260

280

300

T m[°

C]

B3S_P1_Air2

B3S_P2_Air22

B3S_P1_N22

B3S_P2_N22

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

a

Powder-Coating

T m[°

C]

C2000_P1_Air2

C2000_P2_Air2

C2000_P1_N22

C2000_P2_N22

TapeConsolidation

LaminateProduction

Thermo-forming

b

Figure 5-9 Development of Tm of a) B3S and b) C2000 under air and nitrogen atmosphere whensubjected to temperature profiles P1 and P2.

Referring to B3S, Tm decreases with increasing heating cycles throughout bothtemperature profiles and atmospheres. At the beginning (powder-coating), Tm isfound to be increased in comparison to manufacturer’s data (220 °C). This maybe attributed to the thermal history during milling the pellets to powder undercryogenic conditions. However, preliminary studies on as-received pellets showed

Degradation of polyamides 93

also higher values for Tm than given by the manufacturer. Comparing measurementsin air and inert atmosphere, higher Tm are found in an inert atmosphere. This isassigned to early oxidative reactions during thermal cycling. With exception of thefirst heating, Tm remains constant for short dwell times and reveals slight decreasesfor long dwell times under an inert atmosphere. In an oxidative atmosphere, Tm

decreases to a greater extent for both short and long dwell times. Chain scissionsduring thermo-oxidative degradation in air are assumed to form more low-molecularcomponents lowering the Tm.The development of the extrapolated Tm of C2000 is reported in Figure 5-9b. Thebimodal melt peak of C2000 arises from two different crystal forms. As observed forB3S, Tm decreases with increasing heating cycles. Excluding the measurement forP1 in air and inert atmosphere, Tm rises from powder-coating to tape consolidation.Powder-coating is conducted at 260 °C which is below Tm and enables fusing of thelower melting crystal form only. At a moderate cooling rate of 20 °C/min in theDSC, the unstable γ - crystal phase is transformed into the more stable α - crystalphase leading to an increase of Tm as known from various nylons [174]. This γ → α -transition is promoted by long dwell times as present during temperature profileP2.Apart from the measurement following P1 in air, higher Tm and less reductionsare found in an inert atmosphere than in air as observed for B3S. Conductingmeasurements in air, Tm declines by 28 °C for short dwell times (P1) and by 14 °Cfor long dwell times (P2), referring to the maximum Tm.

5.4.2 Effect on mass loss

In accordance to the procedure presented in subsection 5.3.3, the mass loss of B3Sand C2000 throughout temperature profile P1 and P2 was determined as presentedin Figure 5-10. In general, the mass loss recorded for B3S rises with increasingheating cycles in air and becomes larger for long dwell times. Lower mass lossesare detected for short dwell times or when measured in an inert atmosphere.As C2000 is not entirely fused during the powder-coating step, a low mass loss isfound. Beginning with tape consolidation, the mass loss decreases with increasingheating cycles for short dwell times in both air and an inert atmosphere. For longdwell times, the mass loss increases until laminate production in air as found forB3S. Instead of further increase in mass loss for the thermoforming step overalllower mass losses were recorded although the temperature was raised by 20 K inthis step.

94 Degradation of polyamides

0

1

2

3

4

5

Mas

slo

ss[%

]

B3S_P1_Air2B3S_P2_Air2B3S_P1_N22B3S_P2_N22

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

a

0

1

2

3

4

5

Powder-Coating

Mas

slo

ss[%

]

C2000_P1_Air2C2000_P2_Air2C2000_P1_N22C2000_P2_N22

TapeConsolidation

LaminateProduction

Thermo-forming

b

Figure 5-10 Mass loss of a) B3S and b) C2000 samples subjected to temperature profile P1 andP2 under air and nitrogen gas atmosphere in TGA.

5.4.3 Impact on complex viscosity

The development of the complex viscosity η∗ of B3S and C2000 is presented inFigure 5-11 as a function of time during the exposure to temperature profile P2 inair and nitrogen atmosphere.

0 2000 4000 6000 8000102

103

104

105

106

B3S_P2_Air2

B3S_P2_N2

Time [s]

η*[P

as]

a

180

200

220

240

260

280Temperature profile P2

T[°

C]

0 2000 4000 6000102

103

104

105

106

C2000_P2_Air2

C2000_P2_N2

Time [s]

η*[P

as]

b

175

200

225

250

275

300

325Temperature profile P2

T[°

C]

Figure 5-11 Development of the complex viscosity η∗ of a) B3S and b) C2000 during subjectionto temperature profile P2 in air and nitrogen gas atmosphere.

For both polymers, a higher viscosity is found during powder-coating comparedto tape consolidation and laminate production because the temperature appliedduring powder-coating is 20 K lower than for the two subsequent process steps.During laminate production with extended dwell time in an oxidative atmosphere,both polymers show significant and irreversible increases in viscosity indicatingthermo-oxidative degradation. Although the temperature is further increased by20 K during thermoforming, the viscosity remains at the same level as found at theend of the laminate production.

Degradation of polyamides 95

In an inert atmosphere, B3S appears to be unaffected by long dwell times reveal-ing a constant viscosity during laminate production. As expected, the viscosity isfurther reduced during thermoforming as the temperature was raised by another20 K. C2000 however shows also increases in viscosity when temperature profile P2is followed in an inert atmosphere.Based on three measurements for each configuration, Figure 5-12 summarizes themean minimum of the complex viscosity η∗ of B3S as well as C2000 when subjectedto P1 and P2 in air and an inert atmosphere.

0

200

400

600

800

1000

1200

Min

imum

η∗ [Pa

s]

B3S_P1_Air2

B3S_P2_Air2

B3S_P1_N2

B3S_P2_N2

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

a

0

300

600

900

1200

1500

1800

Min

imum

η∗ [Pa

s]

C2000_P1_Air2

C2000_P2_Air2

C2000_P1_N2

C2000_P2_N2

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

b

Figure 5-12 Minimum complex viscosity η∗ of a) B3S and b) C2000 subjected to temperatureprofiles P1 and P2 in air and nitrogen atmosphere.

For short dwell times in both air and an inert atmosphere, the viscosity of B3S(Figure 5-11a) remains unaffected and decreases with increasing temperature usedfor the single process steps. After considerable increases in viscosity when exposedto long dwell times in air, the highest viscosity of B3S is found for the thermoform-ing step in Figure 5-12a. Compared to the beginning of the laminate production,the viscosity increases by 412 % until thermoforming.As Figure 5-11b shows, an early initiation of thermo-oxidative degradation reac-tions is found for C2000 at the beginning of the laminate production step. Anincreased viscosity was measured despite the application of constant temperatures(300 °C) for tape consolidation and laminate production. This phenomenon is ob-served for all tested C2000 specimens. Compared to the beginning of the laminateproduction step, the viscosity has increased by 608% until thermoforming.The significant viscosity increases of both B3S and C2000 observed during expo-sure to temperature profile P2 in air are attributed to diffusion of both dissolvedoxygen and moisture over time and temperature. In addition, reactive end groupsare formed by chain scission during thermo-oxidative degradation. Once a sufficientconcentration of reactive media is reached, post-condensation and crosslinking reac-tions are initiated that cause increases in viscosity. These secondary reactions from

96 Degradation of polyamides

thermo-oxidative degradation were found to be initiated in C2000 earlier than inB3S. In addition, the relative increase of the complex viscosity of C2000 duringlaminate production is 48 % higher than for B3S. Despite of the aromatic tereph-thalic acid in C2000 that enhances the heat resistance compared with aliphaticnylons, the investigated C2000 appears to be more sensitive to thermal as well asto thermo-oxidative degradation than the studied B3S. As a consequence of largermolecular distances of C2000 involving shorter diffusion paths, the critical concen-tration of reactive media may be reached more rapidly than in B3S. In addition,the higher sensitivity to degradation can be also attributed to a moderate heatstabilization of B3S whereas C2000 is not thermally stabilized according to manu-facturer’s data.Considering CFRTP production processes, external pressure is applied that acts interms of shear stress on the matrix polymer in addition to thermal stress. By usingrheometry, the effect of both shear and thermal stress is investigated when poly-mers are subjected to repeated heating cycles. Hence, rheometry can realisticallyrepresent the effects on further processing induced by thermal cycling on polymerproperties.

5.4.4 Effect on molecular composition

Figure 5-13 shows the molecular weight distribution (MWD) of B3S and C2000 afterexposure to temperature profile P2; air and nitrogen were employed as purging gas.In addition, the MWD of as-received, non-processed B3S and C2000 pellets act as

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

W(lo

gM

)[-]

Molar mass [g/mol]

As-received2

B3S_P2_Air2

B3S_P2_N2

a

W(lo

gM

)[-]

Molar mass [g/mol]

As-received2

C2000_P2_Air2

C2000_P2_N2

b

Figure 5-13 MWD of a) B3S and b) C2000 as-received and after exposure to temperature profileP2 under air and nitrogen gas atmosphere.

a reference before thermal cycling. Referring to B3S subjected to long dwell timesunder an oxidative atmosphere, the Mw is slightly increased and show comparableMn to as-received samples. Samples exposed to temperature profile P2 in an inert

Degradation of polyamides 97

atmosphere reveal higher Mw and slightly increased Mn in comparison to samplesexposed to an oxidative atmosphere and as-received samples. An increased Mn isattributed to post-condensation reactions enabled during processing. In addition,the MWD of B3S samples broadens after exposure to temperature profile P2 underoxidative and inert atmosphere. Altogether, thermal cycling has a low effect on themolecular composition of B3S.However, thermal cycling was found to have a larger impact on C2000 (Figure 5-13b). Lower Mw and broadened MWDs are found for C2000 processed accordingto temperature profile P2 in air and an inert atmosphere compared to as-receivedsamples. Subjecting C2000 samples to oxidative or nitrogen gas atmosphere forlong dwell times results in similar MWDs. However, the samples exposed to airshow higher Mw shifted to lower Mn.

5.4.5 Processing window

Considering the laminate production step of temperature profile P2, the effect ofthe significant increase of the melt viscosity on the impregnation time is calculatedby using the model developed in Chapter 4. The initial fiber bundle height was setto 500 μm to account for woven fabrics or NCF. The initial DOI was set to 0 andis representative of processing powder-coated tows in a static press for 60 min aspresented in Chapter 2.Figure 5-14 shows the development of the DOI of CF-TP/B3S over time for a con-stant viscosity of 185 Pa s, which is the minimum viscosity measured for B3S atthe beginning of the laminate production step, compared to the viscosity increaseas a function of time.

a) Constant η = 185 Pa s (B3S) b) Viscosity increase over time (B3S)

Figure 5-14 Development of the DOI of CF-TP/B3S as a function of a) constant viscosity and b)viscosity development as measured for temperature profile P2 during the laminateproduction step.

The increase in viscosity directly affects the impregnation time. For a constantviscosity, 100 % impregnation is yielded after 1526 s. The increase in viscosity as

98 Degradation of polyamides

measured for the laminate production step of temperature profile on neat B3S re-sults in an increase of the impregnation time by 13 %.In Figure 5-15a, the DOI of CF-TP/B3S is presented as a function of time for aconstant viscosity of 195 Pa s that was measured for C2000 at the beginning of thelaminate production step compared to the viscosity increase as a function of time(Figure 5-15b).

a) Constant η = 195 Pa s (C2000) b) Viscosity increase over time (C2000)

Figure 5-15 Development of the DOI of CF-TP/C2000 as a function of a) constant viscosityand b) viscosity development as measured for temperature profile P2 during thelaminate production step.

The effect of the viscosity increase on the impregnation time is more pronouncedfor C2000. Applying the constant viscosity of non-degraded C2000, complete im-pregnation is achieved after 1609 s. When the viscosity profile measured during thelaminate production step is used, the DOI reaches 42 % after 1700 s.To ensure processing at a constant viscosity and a time-efficient production ofthermoplastics with B3S and C2000 it is recommended to restrict each processstep to a dwell time of less than 5 min. This can be achieved by consolidationof powder-coated tows to tapes in a double-belt press prior to laminate produc-tion. The processing window applies in particular to woven fabrics or NCF withlarge flow paths. In addition, this processing window describes the time until theviscosity starts to increase and hence is a conservative value still allowing someimpregnation progress.

5.4.6 Conclusion and implications

To enable gradual impregnation during the production of thermoplastic compos-ites, B3S and C2000 were investigated with regard to thermal cycling as presentduring processing. Two temperature profiles were derived from the process stepspowder-coating, tape consolidation, laminate production and thermoforming re-quired for CFRTP production. The profiles revealed two different dwell times andwere applied to both polymer types.

Degradation of polyamides 99

The methods DSC, TGA, and rheometry served to detect changes in thermal prop-erties and complex viscosity as a result from thermal cycling. The results obtainedfrom all methods led to a more comprehensive understanding of degradation reac-tions that occur during thermal cycling. In the DSC with oxidative atmosphere, de-creasing melting temperatures Tm indicated thermo-oxidative degradation of bothB3S and C2000. The results from TGA that show increasing mass loss of B3S andC2000 with extending dwell times in the presence of oxygen support the indica-tions from DSC. In rheological studies, significant increases in complex viscosityη∗ for both B3S and C2000 were detected when subjected to long dwell times inpresence of oxygen. The results obtained from all methods confirm the thermo-oxidative degradation of both polyamide types during thermal cycling throughoutthe production of CFRTP. No correlation was found between the development ofthe complex viscosity η∗ and changes in number-average molecular weight Mn andweight-average molecular weight Mw as measured by GPC.In inert atmosphere, the complex viscosity η∗ of B3S remained unaffected, indicat-ing a low susceptibility to thermal degradation when additionally exposed to shear.In contrast, C2000 showed considerably enhanced complex viscosity also under ex-clusion of oxygen.Altogether, the results from this study set the basis for instructions during process-ing but also during recycling of both investigated polyamide types. The viscosityof both B3S and C2000 was increased by secondary reactions of thermo-oxidativedegradation. Using the model developed in Chapter 4, the effect of increasing vis-cosity during processing was simulated for both B3S and C2000. The increase inviscosity is more moderate for B3S than for C2000 but results in a rise in im-pregnation time by 13 % compared to a viscosity that remains at a constant level.The effect of increasing viscosity during processing was much more pronounced forC2000. While a constant viscosity would yield complete impregnation after 1700 s,increasing viscosity as measured on C2000 during the laminate production stephinders complete impregnation. The significant increase in viscosity led to insuffi-cient impregnation indicated by a DOI of 42 %.The number of thermal cycles was found to have a negligible effect on thermo-oxidative degradation of B3S and C2000 when dwell times are limited to 5 min.Considering the concept of gradual impregnation during the production of ther-moplastic composites, complete impregnation is hindered by additionally increasedviscosities. Therefore, the impregnation process is recommended to be completedbefore thermo-oxidative degradation is initiated which leads to a processing win-dow of less than 5 min per process step for the investigated polyamide types. Asthis processing window describes the time until the viscosity starts to increase, itis a conservative value that still allows some impregnation progress.

6 Thermal stabilization and flowpromotion of polyamides

The effect of thermal cycling on neat polymers was studied in Chapter 5 and re-vealed significant increases in viscosity when B3S and C2000 were subjected todwell times of more than 5 min in an oxidative atmosphere. Considering gradualimpregnation of partially impregnated intermediates throughout CFRTP produc-tion, the substantial viscosity increase can slow down the impregnation progressand even hinder complete impregnation in case of C2000.By adding various antioxidants to both B3S and C2000, the polymers are thermallystabilized to prevent drastic increases in viscosity. In addition, several lubricants areselected and compounded to both polymers to further decrease the viscosity facili-tating gradual impregnation. After investigations on single-modified polymers withonly one additive, the most effective antioxidant and lubricant is further combinedand potential interactions are studied on multi-functionalized B3S and C2000. Bycoating CF-TP fibers with the multi-functionalized powder, consisting of antiox-idant, lubricant and polymer, intermediates are produced. These are stacked andpressed to multi-ply test panels. Flexural properties are determined from these testpanels to investigate the effect of these additives on the mechanical behavior onthe composite level. Parts of the following section have previously been publishedin [144].

6.1 Thermal stabilization

The addition of stabilizers cannot prevent the initiation of thermo-oxidative degra-dation [104] but has proven to limit it during reprocessing or recycling of polyamides[165, 168]. Antioxidants are commonly classified into chain breaking (CB) and pre-ventive antioxidants due to their functional principle [175].

6.1.1 Chain breaking antioxidants

CB antioxidants terminate the oxidation process by removing the propagating rad-icals according to the donor or the acceptor mechanism. Chain breaking donorantioxidants (CB-D) donate electrons or hydrogen atoms and react with a rad-ical to form a stable, reduced radical. During the chain breaking acceptor (CB-A) mechanism the antioxidants oxidize alkyl radicals under formation of a stable

101

102 Thermal stabilization and flow promotion of polyamides

molecule [175]. A stabilization reaction of CB-D antioxidants can proceed as fol-lows [175]:

AH + ROO ROOH + A (6-1)

where A denotes the antioxidant with hydrogen atom H to be donated to a hy-droperoxyl radical ROO . For an effective stabilization, the antioxidant radical Amust form stable products during propagation reactions [175]:

A +RH AH + R ROO (6-2)

A+O2/RH

AOOH + R ROO (6-3)

The stabilization reaction of CB-A antioxidants by oxidizing alkyl radicals com-petes with the propagation reactions and are most effective in absence of oxy-gen [175]:

A + R O2 deficient Non-reactive products (6-4)

Typical examples for CB-D antioxidants are sterically hindered phenols and aro-matic amines [104]. Phenolic antioxidants react with oxygen radicals by donatinga hydrogen atom to form hydroperoxide ROOH as shown in Figure 6-1.

R3

R2OH

R1

+ ROO

R3

R2O

R1

+ ROOH

Figure 6-1 Stabilization reaction using sterically hindered phenols [176]; R1,R2, and R3 denotemoiety.

Sterically hindered phenols are often used in aliphatic polyamides [86, 177, 178]but have a lower antioxidative effect in polyamides than copper salts [179, 180] andare more effective in polyolefins [159]. Besides copper-based stabilizers, aromaticamines represent the most effective antioxidants for aliphatic polyamides [181], es-pecially in PA6 types that were hydrolytically and anionically prepared [155]. Theprinciple stabilization reaction of aromatic amines is shown in Figure 6-2.

NH

R + ROO N R + ROOH

Figure 6-2 Stabilization reaction of aromatic amines [176].

Due to discoloration of the polymer the use of aromatic amines is restricted topigmented polymers or shielded products [104].

Thermal stabilization and flow promotion of polyamides 103

6.1.2 Radical scavengers

Sterically hindered amines (HAS) are designed to stabilize polymers at low tem-peratures and are also known as hindered amine light stabilizers (HALS) dueto their excellent stabilization against light [176]. The stabilization mechanismof HAS/HALS is based on the ability to scavenge alkyl radicals by formation ofnitroxyl radicals [104]. These radicals further react with alkyl radicals to alkyl hy-droxylamines by the CB-A mechanism. In general, the reaction rate of oxygen withalkyl radicals is extremely high in comparison to reaction mechanisms through CB-A/CB-D stabilization or radical scavenging [176]. Hence, many radicals are rapidlyformed as they cannot be neutralized by stabilizers at the same rate.

6.1.3 Preventive antioxidants

Preventive antioxidants can retard thermo-oxidative degradation by preventing thegeneration of new radicals. The most important preventive antioxidants are repre-sented by hydroperoxide decomposers [175].Hydroperoxide decomposers react with hydroperoxides into stable products withoutformation of further radical products [104, 176]. Hence, the generation of radicalsis already prohibited by using hydroperoxide decomposers. They must exhibit ahigher thermal stability than the hydroperoxide itself for efficient decomposition.Then, chain branching initiated by alkoxy and hydroxyl radicals can be prevented.While the hydroperoxide decomposer oxidizes, it initiates the reduction of the hy-droperoxide group to an alcohol [176].Common representatives of preventive antioxidants are organic compounds basedon phosphorous (phosphites, phosphonites) or sulfur (sulfides), metal salts, anddithiophosphates [176]. Phosphites and phosphonites, often used in combinationwith sterically hindered phenols, are known for stabilization during processing [172,182]. The general reaction scheme of phosphites during stabilization proceeds asfollows [183]:

P(OR)3 + ROOH O P(RO)3 + ROH (6-5)

P(OR)3 + ROO O P(RO)3 + RO (6-6)

P(OR)3 + RO O P(RO)3 + R (6-7)

where R can be an alkyl or aryl group. Using phosphite compounds, the recy-cling of any polycondensation polymer including PA6 is enabled with maintained

104 Thermal stabilization and flow promotion of polyamides

mechanical properties as water is removed by this antioxidant [184]. In addition,sulfur-based thiosynergists are used to decompose hydroperoxides suitable for sta-bilization during the utilization phase [172, 182].Copper salts or in combination with halogen ions are traditionally used as sta-bilizers in polyamides [185]. Among various examined metal salts such as CoCl2,CuCl, CuCl2, CuI, NiCl2 or ZnCl2, copper salts were found to be most effective todecelerate thermo-oxidative degradation [180, 186].When stabilization during the processing and utilization phase is desired, preven-tive antioxidants are often combined with CB-D antioxidants [176]. As this studymainly focuses on the stabilization of polymers during processing to obtain CFRTPcomponents, preventive antioxidants are preferred.

6.1.4 Investigated antioxidants

In this chapter, B3S and C2000 are modified with several antioxidants to reducethe extent of thermo-oxidative degradation reactions. To ensure efficiency and pro-cessability, the antioxidants must fulfill the following requirements:

• Thermal stability

• Hydrolytic stability

• Low volatility

• Chemical compatibility

Based on these requirements, commercially available CB and preventive antioxi-dants were selected that are described in more detail in the following.

Hostanox O310

Hostanox O310 is supplied by Clariant and composes of sterically hindered phenols.The low Tm of 90 °C allows for processing via compounding. Amounts of 0.1-0.5 wt%are recommended. The additive is characterized by hydrolitic stability leading toan effective stabilization [187].

Brüggolen H320

Brüggolen H320 supplied by Brüggemann Chemical is based on a copper-halogencompound consisting of CuI and KI. According to manufacturer’s instructions, theaddition of 100 ppm Cu yields sufficient stabilization, leading to a weight per-centage of 0.2-0.4 wt%. Brüggolen H320 is designed for long-term stabilization atservice [188].

Thermal stabilization and flow promotion of polyamides 105

Hostanox P-EPQ

Hostanox P-EPQ composes of phosphorous compounds with diphosphonite as maincomponent. Being a processing stabilizer against thermo-oxidative degradation,the melt viscosity can be held constant even at high shear stresses and elevatedtemperatures. In addition, the secondary antioxidant is suitable for the use inpolyamides and is compatible to a wide range of other polymers [189].

Brüggolen H10

Brüggolen H10 received from Brüggemann Chemical is suitable for the stabiliza-tion of polyamides during processing and helps to reduce yellowing during meltprocessing. It is composed of inorganic phosphonate powder and is compoundedto the host polymer with mass fractions between 0.1 and 0.3 wt%. The stabilizerpossesses no volatility and an excellent resistance against high thermo-mechanicalstress as present during compounding and processing [188].

Nylostab S-EED

Nylostab S-EED supplied by Clariant is a multifunctional additive with a HALSstabilizer originally designed to protect pigmented polyamides against light-induceddecomposition. The aromatic amide structure of Nylostab S-EED is responsible forthe high thermal stability and is highly compatible to polyamides due to a resem-bling molecular composition. The hindered piperdine part leads to stabilization ofthe host polymer against light [190]. Nylostab S-EED shows also characteristics ofa processing aid improving melt stability [191].

Table 6-1 summarizes the selected antioxidants used to limit thermo-oxidativedegradation of B3S and C2000 during processing.

Table 6-1 Selected additives for thermal stabilization of B3S and C2000.

Designation Type Chemical Basis Addition[wt%]

Hostanox O310 CB-D Sterically hindered phenols 0.35Brüggolen H320 Preventive AO Copper salt, CuI, KI 0.30

Hostanox P-EPQ Hydroperoxidedecomposer Phosphor compounds 0.20

Brüggolen H10 Hydroperoxidedecomposer Inorganic phosphonate 0.30

Nylostab S-EED Radical scavenger,processing aid HALS 0.50

106 Thermal stabilization and flow promotion of polyamides

6.2 Flow promotion

In principle, reductions in viscosity of thermoplastics are achieved by increases inprocessing temperature as shown in subsection 4.5.1. However, an excessive rise intemperature leads to decomposition, release of low-molecular products or degrada-tion [192]. The modification of thermoplastics with lubricants offers another pathto reduce the viscosity, ideally without changing the mechanical properties. Lu-bricants are not only applied to increase flowability but also to lower materialdamage during processing. Less stress acts on the polymer when friction is reducedas shear forces decrease. In addition, dissipation is lowered that prevents local over-heating [192].In general, lubricants may be divided into internal and external lubricants. Theaddition of lubricants aims to decrease internal and external friction. As a con-sequence, shear (and therefore temperature as well as degradation reactions) islowered, equipment wear is reduced, production rates can be increased, and lessenergy is consumed [175].

6.2.1 Internal and external lubricants

External lubricants are typically insoluble and are applied externally to polymersto e.g. reduce the adhesion to the tool surface. They are also referred to as releaseagents or slip agents.Internal lubricants represent processing or flow aids and are further classified intoinner and outer lubricants. With 50%, lubricants are mainly applied to polyvinylchloride (PVC) followed by engineering thermoplastics and other polymers [193].The main purpose of inner (internal) lubricants is flow promotion. Inner lubricat-ing agents are usually soluble in polymers and polar. They facilitate processingby reducing the inner friction between molecular chains and hence, the viscos-ity [192, 194]. Outer lubricants offer low solubility resembling external lubricants.They diffuse to the surface of the liquid or molten polymer and form a separativelayer between component and tool upon cooling to reduce the friction between meltand tool [175, 176, 195]. This effect is known as the slip-effect and can be inducedby amides of several monounsaturated fatty acids [176].Nowadays, lubricants typically consist of polar and nonpolar groups leading toflow promoters with internal and external characteristics [175]. The extent of in-ternal and external characteristics is dependent on the lubricant itself, the poly-mer and the metering as Figure 6-3 shows. General lubricating agents are fattyalcohols/long-chain alcohols, fatty acids and salts, fatty acid amides, fatty acid es-ters, montan waxes, polar and non-polar polyolefin waxes and paraffin waxes [176,

Thermal stabilization and flow promotion of polyamides 107

192, 195]. Waxes represent the largest group of lubricants for polymers followedby fat derivatives [176]. Fatty alcohols and fatty acids are highly effective in PVCbut show a high volatility at the same time. For engineering plastics such as PA6,often stearic acid esters are employed owing to their high thermal stability and lowvolatility [192].

Solubility

Pro

cess

abili

ty

Figure 6-3 Effect of lubricants as a function of solubility in the host polymer; redrawn from [192].

If amides are used as lubricants it is commonly referred to as bis-stearyl ethylene-diamine/ethylenbisstearamide (EBS) or amide wax [175]. In contrast to other rep-resentatives such as oleic acid amide or erucamide, EBS is not characterized by apronounced slip-effect [192].Polyolefine waxes, paraffin waxes or microcrystalline waxes belong to the classof hydrocarbons. Some esters that are used as lubricants are simple esters, glyc-erol esters, polyglycerol esters or montan esters. With exception of montan esters,these lubricants are fat derivatives. Montan waxes resemble fatty acids in theirchemistry but are longer than fatty acids which changes properties significantly.Montan waxes and its derivatives in form of esters and soaps are widely used inmany polymers. They are characterized by a well-balanced mix of external andinternal effects. In addition, montan waxes and its derivatives show high thermalstability, low tendency for migration and low volatility [176, 192].Lubricants that are effective in particular in polyamides are montan wax, amidewax, copolymer wax, wax esters and polyol wax by improving the flowability andrelease effect. The best dispersion is achieved by montan waxes, polyol esters andwax esters [176].Due to the aim of this study to lower the viscosity during processing, only lubri-cants with high thermal stability came into consideration. High thermal stabilityis ensured by using amide waxes, montan esters or fatty acid esters [192].

108 Thermal stabilization and flow promotion of polyamides

6.2.2 Investigated lubricants

Commercially available, internal lubricants were selected (Table 6-2) to investigatethe effect on the flowability of B3S and C2000. The selection was based on thefollowing requirements:

• Thermal stability

• Compatibility to host polymer and other additives

• Migration resistance

• Low volatility

Brüggolen P14

The processing agent Brüggolen P14 supplied by Brüggemann Chemical combinesinternal and external characteristics with low volatility and high thermal stability.It is applied to polyamides to broaden the processing window [188].

Licowax C

Licowax C received from Clariant used in engineering thermoplastics is an amidewax of the N,N’-bisstearoylethylenediamine (EBS) type that possesses high ther-mal stability. In addition, the vegetable-based Licowax C is characterized by lowvolatility [196].

Nylostab S-EED

As previously mentioned, Nylostab S-EED is a multifunctional additive with an-tioxidative and lubricating character. The composition imitates the structure ofpolyamides and leads to significantly improved melt processing. An additional meltstabilizing effect was observed for Nylostab S-EED which has not been detectedfor other HALS structures. Transamidation reactions that occur during melt pro-cessing may be the reason for improved melt stability [190].The introduced lubricants are summarized in Table 6-2 along with the used amounts.

Table 6-2 Selected additives to increase flowability.

Designation Type Chemical Basis Addition [wt%]Brüggolen P14 Lubricant Polyethylene wax 0.35Licowax C Lubricant Amide wax 0.30Nylostab S-EED HALS, processing aid Amide basis 0.50

Thermal stabilization and flow promotion of polyamides 109

6.3 Experimental methods

To begin with, the efficiency of the selected antioxidants and lubricants was stud-ied separately (single-modified polymers). The additives found most effective werefurther combined to produce multi-functionalized polymers. For the investigations,the temperature profiles derived from a typical CFRTP production process as de-scribed in section 5.2 were used. In addition, the influence of the combined use ofadditives on the mechanical behavior is investigated on composite level.

6.3.1 Polymer samples

Sample preparation

In cooperation with the Institute of Medical and Polymer Engineering of the Tech-nical University of Munich, the single and combined additives were compounded toneat B3S and C2000. For compounding, a twin-screw co-rotating extruder (Cope-rion ZSK 18ML) with volumetric dosing (Brabender DSR 28) was used. A strandpelletizer (Pell-Tec SP30) served to produce pellets after passing a cooling sec-tion (Pell-Tec CT120). The extruder possesses seven, independently controllableheating zones. The used temperature profiles of the heating zones or compoundingadditives to B3S and C2000 are presented in Figure 6-4 for B3S and C2000.

Volumetric dosing

Twin-screwextruder

B3S

C2000

270°C 270°C 265°C 260°C 255°C260°C 250°C

330°C 330°C 325°C 320°C 315°C320°C 310°C

Figure 6-4 Temperature profiles used for compounding additives to B3S and C2000 in a twin-screw extruder.

Using Nylostab S-EED, the high Tm of 272 °C required the production of a mas-terbatch with 10 wt% additive by increasing the depicted temperature profiles by20 K. In another compounding step, the desired mass fraction of 0.5 wt% was ad-justed. The rotational speed was 150 rpm for B3S and 250 rpm for C2000. Thehaul-off speed of the strand pelletizer was 30 m/min.

110 Thermal stabilization and flow promotion of polyamides

Differential scanning calorimetry

The oxidation induction time (OIT) of single-modified and multi-functionalized(MF) polymers was determined to study the effectiveness of the selected antiox-idants and combined additives by using a calorimeter of the type TA Instru-ments Q100. After determination of the required holding temperature using thedynamic method based on DIN EN 728 [197], the static method was applied. Un-der nitrogen atmosphere, the samples were heated with 20 °C/min to the requiredholding temperature. After holding the temperature constant for 3 min the purg-ing gas was switched to air. The temperature was maintained until the exothermicdegradation reaction occurred. The oxidation induction time was determined fromthe point at which the gas was switched until the degradation reaction becameapparent. The OIT was determined again for polymers with the most efficientantioxidant and lubricant.

Thermogravimetric analysis

The mass loss of single-modified and multi-functionalized B3S and C2000 was mea-sured by using a TA Instruments Q500 during exposure to temperature profile P1under an oxidative atmosphere. For each of the process steps from powder-coatinguntil thermoforming a new specimen was used resulting in four samples per poly-mer. Here, the thermal stability and the effectiveness of antioxidants is studiedwhen the samples are heated above Tm for several times. The procedure and eval-uation procedure developed in subsection 5.3.3 was maintained for single-modifiedand multi-functionalized B3S and C2000. Results for neat polymers were trans-ferred from Chapter 5.

Rheometry

Rheometry served to determine the most efficient lubricant for B3S and C2000.Polymer samples modified with the selected lubricants were subjected to tempera-ture profile P2 under an oxidative atmosphere. The testing procedure is describedin subsection 5.3.4. Using temperature profile P2 under oxidative and inert con-ditions, rheological properties were determined for multi-functionalized B3S (MF-B3S) and C2000 (MF-C2000) to identify possible interactions in comparison tosingle-modified polymers. Results for neat polymers were transferred from Chap-ter 5.

6.3.2 Composite samples

In addition to studies on single-modified and multi-functionalized polymers, theeffect of the combined use of additives was investigated on composite level as de-

Thermal stabilization and flow promotion of polyamides 111

scribed in the following. The main purpose is to identify potential changes in themechanical properties by the modification with additives. The multi-functionalizedpolymers (MF) produced via compounding were cryogenically ground to powderby means of liquid nitrogen. The laboratory mill ZM200 by Retsch was adjustedto a rotational speed of 8000 rpm. Due to centrifugal forces, the polymer pelletshit a circular sieve with a size of hole of 500 μm and crush the polymer pellets intopowder. Powder with a particle size of 250 μm by majority (more than 80 %) wasproduced.For reference, as-received polymers were compounded and milled in the same wayas described above for the multi-functionalized polymers. In addition, spread car-bon fibers from the same mother spool were used for all test panels. Thus, fullcomparability of test panels with multi-functionalized and non-modified polymersis ensured.

Four-point bend test

The milled powder was further used to manufacture powder-coated tows reinforcedwith polyamide-sized carbon fibers (CF-TP) as described in section 2.3.1, stackedand pressed to obtain test panels in accordance to section 2.3.2. Two different dwelltimes were chosen: half and full dwell time. Test panels made from CF-TP/MF-C2000 were pressed for 15 (half) and 30 min (full), respectively. The press time forlaminates produced from CF-TP/MF-B3S was set to 30 (half) and 60 min (full).The full dwell times represent the laminate production step of temperature profileP2.According to DIN EN ISO 14125 B [62], seven specimens were tested per configu-ration to determine the longitudinal and transverse flexural strengths (σf1 and σf2).Longitudinal tests were conducted with a support span of 81 mm and a load spanof 27 mm at a crosshead speed of 5.00 mm/min. Transverse tests were carried outwith a support span of 45 mm and a load span of 15 mm at a crosshead speed of2.00 mm/min. The used test equipment is described in section 2.2.1. The outer fiberstrain was derived from the deflection in accordance with DIN EN ISO 14125 B -section 10.2.3. The fiber volume content (FVC) of each test panel was determinedat three different positions by acid digestion according to DIN EN 2564 B [106] atthe central laboratory of SGL Carbon GmbH.

Gel permeation chromatography

GPC measurements were conducted to detect changes in Mn, Mw and MWD arisingfrom modification with most efficient antioxidant and the combined used of addi-tives. In addition, samples extracted from multi-functionalized test panels were ana-lyzed in comparison to non-modified laminates. GPC measurements were conducted

112 Thermal stabilization and flow promotion of polyamides

in cooperation with the WACKER-Chair of Macromolecular Chemistry from theTechnical University of Munich by using the chromatograph Agilent 1200 Series.The same procedure was followed as previously described in subsection 5.3.5.

6.4 Results

The effectiveness of the selected antioxidants is compared by measuring the OITand the mass loss obtained by TGA. Rheometry was used to determine the effecton flowability induced by lubricants. Based on these methods, the most effectiveantioxidant as well as lubricant is selected. Possible interactions of both additivesare monitored by measuring the OIT, mass loss and rheological properties of multi-functionalized polymers. By producing intermediates via coating carbon fibers withthe multi-functionalized polymers and manufacture of test panels, the effect of theadditives on mechanical properties is tested on composite level.

6.4.1 Effectiveness of additives

Effectiveness of antioxidant

The effectiveness of the selected antioxidants (see Table 6-1) was evaluated by mea-suring the OIT using DSC. The oxidation induction temperature of neat polymersamples was determined to 320 °C for B3S and to 340 °C for C2000. The results forthe OIT of neat and modified B3S and C2000 are reported in Figure 6-5.

Neat

H10

H320

Nylostab

O310

P-EPQ

0 10 20 30 40 50Oxidation induction time (OIT) [min]

OIT of C2000 at 340°COIT of B3S at 320°C

Figure 6-5 The OIT of B3S samples at 320 °C and C2000 samples at 340 °C, both neat andmodified with antioxidants.

In comparison to neat B3S and C2000, specimens modified with inorganic phospho-nate (H10), copper salts (H320) and phosphor compounds (P-EPQ) can improve

Thermal stabilization and flow promotion of polyamides 113

their OIT. The modification with sterically hindered phenols (O310) even reducedthe oxidative stability of B3S and showed no improvement for C2000. The low Tm

of approx. 90 °C of O310 may lead to an early decomposition of the antioxidant atthe high oxidation induction temperatures for B3S and C2000.The multifunctional additive Nylostab S-EED is found to be very effective in B3Ssamples but shows only a marginal improvement of the thermal stability of C2000.The lower effectiveness of Nylostab S-EED in C2000 may be attributed to the cyclicstructures present in additive as well as host polymer. These structures can preventa sufficient approximation of both molecules to develop an effective stabilization.In general, secondary antioxidants on the basis of phosphorous such as H10 andP-EPQ show the largest improvements of the OIT. This is attributed to the highthermal stability of these structures. The higher effectiveness of P-EPQ is explainedby the higher oxidation state (+5) than present in H10 (+3): oxidation of phos-phonites (P-EPQ) yields phosphonates (H10) first. Hence, P-EPQ reveals a higherpersistence due to an additional possibility to react with radicals. The improvedefficiency of the primary antioxidant H320 compared to phenols can be explainedby the higher thermal stability of copper-based structures. However, the stabiliz-ing effect is lower than found for phosphorous-based antioxidants due to the loweroxidation state of copper.The maximum enhancement of the thermo-oxidative stability for both B3S (by28 % to 46.36 min) and C2000 (by 117 % to 41.55 min) was achieved by using thepreventive antioxidant P-EPQ that is based on phosphonites.The efficiency of antioxidants was further investigated by means of the TGA. Byusing temperature profile P1 in an oxidative atmosphere, the effect of repeatedheating above Tm is studied. The mass loss measured across various process stepsfrom powder-coating until thermoforming is compared for neat and single-modifiedpolymers as shown in Figure 6-6.

0.00.20.40.60.81.01.21.41.61.82.0

Mas

slo

ss[%

]

Neat B3S_P1B3S+H10_P1B3S_H320_P1B3S+Nylostab_P1B3S+O310_P1B3S+P-EPQ_P1

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

a

0

1

2

3

4

5

6

7

Mas

slo

ss[%

]

Neat C2000_P1C2000+H10_P1C2000_H320_P1C2000+Nylostab_P1C2000+O310_P1C2000+P-EPQ_P1

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

b

Figure 6-6 Mass loss of neat and single-modified a) B3S samples and b) C2000 samples duringtemperature profile P1.

114 Thermal stabilization and flow promotion of polyamides

During powder-coating, only B3S modified with Nylostab S-EED showed less massloss than the neat polymer. The last step (thermoforming) represents the most in-teresting result as it shows the effect of thermo-oxidative reactions upon heating thesamples for four times. Here, P-EPQ yields the largest improvement of the thermalstability related to mass loss. The phosphonate H10 and the phenol-based O310appear to lose their stabilizing effect as soon as the sample is heated to 280 °C.Considering neat and modified C2000 as presented in Figure 6-6b, all antioxidantshave a detrimental effect on the recorded mass loss during thermoforming, withexception of P-EPQ. The substantial increases in mass loss indicate a decliningstabilizing effect. However, P-EPQ presents the lowest stabilization during powder-coating and laminate production.Altogether, the results from OIT clearly indicate that P-EPQ yields the largestincrease in thermal stabilization. The results from TGA referring to the antioxida-tive effect under process-related conditions are more ambiguous. With focus on thelast process step after four heating cycles, P-EPQ still reveals a stabilizing effecton both B3S and C2000 indicated by decreased mass loss. Hence, this secondaryantioxidant was selected for the following investigations.

Effectiveness of lubricant

With neat B3S and C2000 as a reference, the complex viscosity of samples modi-fied with the multifunctional additive Nylostab S-EED (Nylostab), the amide waxLicowax C (Licowax), and the polyolefine wax (P14) was determined. Figure 6-7ashows the results for neat and modified B3S when subjected to temperature profileP2 in air and inert atmosphere.

0

200

400

600

800

1000

1200

η*[P

as]

Neat B3S_P2_Air2

B3S+Nylostab_P2_Air2

B3S+P14_P2_Air2

B3S+Licowax_P2_Air2

Neat B3S_P2_N2

B3S+Nylostab_P2_N2

B3S+P14_P2_N2

B3S+Licowax_P2_N2

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

a

0

500

1000

1500

2000

2500

3000

3500

η*[P

as]

Neat C2000_P2_Air2

C2000+Nylostab_P2_Air2

C2000+P14_P2_Air2

C2000+Licowax_P2_Air2

Neat C2000_P2_N2

C2000+Nylostab_P2_N2

C2000+P14_P2_N2

C2000+Licowax_P2_N2

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

b

Figure 6-7 Complex viscosity of neat and modified a) B3S and b) C2000 subjected to temper-ature profile P2 under oxidative and inert atmosphere.

From the powder-coating step until laminate production, a reduction in complexviscosity can be achieved by every added lubricant in comparison to as-received

Thermal stabilization and flow promotion of polyamides 115

B3S. In an oxidative atmosphere, the largest viscosity reduction of 32 % on aver-age is yielded by adding Licowax C to B3S. In completely inert atmosphere, theaddition of P14 gains the largest decrease in mean viscosity by 25 %, followed byLicowax C with 19 % and Nylostab S-EED with 12 %. The considerable increase inviscosity of neat B3S observed during the thermoforming step results from subject-ing the samples to a dwell time of 60 min at 260 °C. The long dwell time well aboveTm was assumed to cause thermo-oxidative degradation reactions, as described inmore detail in Chapter 5. The addition of lubricants led to further increases incomplex viscosity up to 24 % (Licowax C) under an oxidative atmosphere and upto 44 % (Nylostab S-EED) under exclusion of oxygen.Focusing on C2000 when subjected to temperature profile P2 as shown in Figure 6-7b, the addition of Licowax C resulted in the largest decrease in viscosity by 25 %on average from powder-coating until laminate production in an oxidative atmo-sphere. The lubricating effect of Licowax C was even more pronounced in an inertatmosphere, revealing a viscosity reduction by 32 % on average compared to neatC2000. The viscosity could also be reduced by 17 % when adding P14 to C2000in an oxidative surrounding. However, the viscosity remains almost unaffected bymodification with P14 under an inert atmosphere. The same phenomenon can beobserved for C2000 modified with Nylostab S-EED.As observed for B3S, C2000 also showed a considerable enhancement of the com-plex viscosity during thermoforming after dwelling at 300 °C for 30 min and in anoxidative environment (laminate production). Thermo-oxidative degradation reac-tions are assumed to cause the considerable increase in viscosity and are discussedin more detail in Chapter 5. Adding lubricants pronounced this effect even more. Afurther increase of up to 162 % is caused by adding P14 to C2000. The modificationwith Nylostab S-EED and Licowax C led to an additional increase of 94 % and92 % respectively. In contrast to the observations for B3S, the viscosity of C2000modified with lubricants further decreased in an inert atmosphere. The addition ofLicowax C enabled another viscosity reduction by 54 %.Table 6-3 summarizes the effects of the selected lubricants on the flowability of B3Sand C2000. The additional increase in viscosity caused by lubricants is attributedto their chemical composition. The additives are composed of carbonyl and aminegroups that can form further radicals and increase the extent of thermo-oxidativereactions. In addition, the thermal stability of the lubricants may be sufficient forcompounding and short processing times. However, the lubricants may start to de-compose during long dwell times at elevated temperatures and form further radicalsthat initiate thermo-oxidative degradation at more sites than in neat polymers. Asa consequence, additional thermal stabilization by antioxidants is required whenan increase in flowability by lubricants is desired.Altogether, Licowax C was found to be most effective in both B3S and C2000 when

116 Thermal stabilization and flow promotion of polyamides

Table 6-3 Effectiveness of different lubricants on B3S and C2000, compared across the firstthree process steps (powder-coating until laminate production) and the final step(thermoforming) of temperature profile P2 in an oxidative and inert atmosphere.

ModificationChange in η∗ [%] Change in η∗ [%]

(Step 1-3) [%] (Thermoforming) [%]Nylostab P14 Licowax Nylostab P14 Licowax

B3S_P2_Air -25 -19 -32 +14 +21 +24B3S_P2_N2 -12 -25 -20 +44 +19 +16C2000_P2_Air -21 -17 -25 +94 +162 +92C2000_P2_N2 -2 +1 -32 -21 -28 -54

the first processing steps from powder-coating until laminate production are con-sidered. The large increase in viscosity after laminate production may be reducedas soon as thermo-oxidative reactions can be decreased. Hence, Licowax C is chosenas most effective lubricant combined with the most effective antioxidant P-EPQ toproduce multi-functionalized polymers. Potential interactions and arising effects ofcombining both additives are investigated in the following section.

6.4.2 Effects of combined use of additives

Oxidation induction time

The OIT of B3S and C2000 modified with both P-EPQ as well as Licowax C iscompared with neat and thermally stabilized polymers in Figure 6-8. The use of

Neat

P-EPQ

P-EPQ +Licowax C

0 10 20 30 40 50 60

Oxidation induction time (OIT) [min]

OIT of C2000 at 340°COIT of B3S at 320°C

Figure 6-8 Complex viscosity of neat and modified C2000 subjected to temperature profile P2under oxidative and inert atmosphere.

both additives yields a further increase of the thermo-oxidative stability. The OIT

Thermal stabilization and flow promotion of polyamides 117

of multi-functionalized B3S is enhanced by 23 % to 57 min in comparison to B3Smodified with P-EPQ only. For C2000, an additional increase of the OIT by 19 %is yielded when Licowax C is added to the single-modified C2000. The synergisticeffect of P-EPQ and Licowax C is attributed to increased mobility of the molecularchains due to less friction induced by the lubricant. This facilitates diffusion ofantioxidants as well as radicals and enables a more effective stabilization.

Mass loss

Table 6-4 summarizes the mass loss recorded for neat and multi-functionalized B3Sand C2000 when exposed to temperature profile P1 in an oxidative surrounding.

Table 6-4 Mass loss of neat and multi-functionalized (MF) B3S and C2000 recorded duringsubjection to temperature profile P1 under an oxidative atmosphere.

Process Step B3S_P1 MF-B3S_P1 C2000_P1 MF-C2000_P1

Powder-coating 0.49 0.80 0.74 1.54Tape consolidation 0.88 1.49 2.99 4.60Laminate production 1.08 2.05 1.71 6.23Thermoforming 1.79 2.70 1.35 7.90

The results from TGA reveal a different trend than the measured OIT. The massloss detected for multi-functionalized polymers are ever larger than for non-modifiedB3S and C2000. Substantial mass losses are recorded for the multi-functionalizedpolymers during the laminate production and thermoforming. These mass losses areattributed to the added lubricant as less decomposition reactions were present inpolymers modified with P-EPQ only. The melting range of Licowax C lies between139 and 144 °C but the temperatures used for temperature profile P1 are muchhigher and can lead to early decomposition of Licowax C. Parts of the detectedmass loss can be attributed to decomposed lubricant. Additionally, degradation ofthe lubricant may result in an enhanced amount of radicals. The reduced frictionbetween molecular chains may promote the diffusibility of the formed radicals andaccelerate degradation reactions. At the same time, the efficiency of antioxidantsis improved due to increased diffusibility but decreases with the amount of newlyformed radicals.

Complex viscosity

As presented for single-modified polymers, the complex viscosity of neat and multi-functionalized B3S and C2000 is plotted against the single process steps of temper-ature profile P2 under an oxidative and inert atmosphere in Figure 6-9. In contrast

118 Thermal stabilization and flow promotion of polyamides

0

200

400

600

800

1000

1200

1400

η*[P

as]

B3S_P2_Air2

MF-B3S_P2_Air2

B3S_P2_N2

MF-B3S_P2_N2

a

0

200

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800

1000

1200

1400

η*[P

as]

C2000_P2_Air2

MF-C2000_P2_Air2

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MF-C2000_P2_N2

b

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

Powder-Coating

TapeConsolidation

LaminateProduction

Thermo-forming

Figure 6-9 Complex viscosity of neat and multi-functionalized a) B3S and b) C2000 subjectedto temperature profile P2 under oxidative and inert atmosphere.

to the thermogravimetric results, a synergistic effect on melt viscosity by combin-ing antioxidant with lubricant is found for both B3S and C2000. Compared to neatpolymers, the combination of P-EPQ and Licowax C yielded an overall decreasein complex viscosity of 13 % for B3S and 50 % for C2000 in air atmosphere. Underexclusion of oxygen, the viscosity was increased by 2 % for MF-B3S and reducedby 68 % for MF-C2000.

Four-point bend test

After the production of intermediates with neat and multi-functionalized polymersreinforced by CF-TP fibers, test panels were pressed with half and full dwell times.This serves to identify potential effects on the viscosity and mechanical properties.It is expected that adding P-EPQ and Licowax C positively influences the impreg-nation progress by reducing the viscosity. The results obtained by four-point bendtesting longitudinal (0°) and transverse (90°) to the fiber direction are presented inFigure 6-10. The results for longitudinal flexural strength σf1 were normalized tothe measured fiber volume content that is stated in the appendix in section A.2.Considering the effect of the dwell time on σf1 overlapping error bars indicate nosignificant influence. For all dwell times, the multi-functionalized laminates yieldedslightly lower σf1. Laminates produced from non-modified and multi-functionalizedC2000 show increased values for σf1. A better alignment of the fibers along 0°-direction may be the reason for increased σf1.The effect of the dwell time on σf2 is more pronounced for B3S and reveals in-creased values for both non-modified and multi-functionalized test panels. The op-posite behavior is observed for test panels produced from non-modified and multi-functionalized C2000 showing decreased σf2 with increased dwell time. The overalllower σf2 for C2000-based test panels can be explained by the lower tensile strength

Thermal stabilization and flow promotion of polyamides 119

CF-TP/B3S

CF-TP/MF-B3S

CF-TP/C2000

CF-TP/MF-C2000

0

200

400

600

800

1000

1200

1400

σ f1[M

Pa]

Half dwell timeLong dwell time

a

CF-TP/B3S

CF-TP/MF-B3S

CF-TP/C2000

CF-TP/MF-C2000

0

20

40

60

80

100

120

σ f2[M

Pa]

Half dwell timeLong dwell time

b

Figure 6-10 Results from four-point bend testing a) in fiber direction and b) perpendicular tofiber direction of test panels produced from non-modified and multi-functionalizedpolymers at different dwell times in a press.

of 62 MPa compared to 90 MPa of B3S. In general, a lower σf2 is found for testpanels made from multi-functionalized polymers. As presented, the suppressionof thermo-oxidative degradation reactions with the help of P-EPQ can avoid thesubstantial increase in viscosity when the polymers are subjected to high tempera-tures for long dwell times. The increase in viscosity is explained by crosslinking andpost-condensation reactions (see Chapter 5 for a more detailed discussion). If thesereactions are suppressed by adding antioxidants a decrease in mechanical perfor-mance is expected. In addition, the modification with P-EPQ and Licowax C mayalter the compatibility to the polyamide-based sizing of CF-TP fibers. The molec-ular composition of P-EPQ and Licowax C is different from the host matrices andfrom the polyamide sizing. The intercalation of the additives in the host matricescan act as disturbing factors in the interface region and reduce the establishmentof covalent bonds.

Effect on molecular composition

The effect of adding P-EPQ and Licowax C to B3S and C2000 on the molecularcomposition was tested on extracted specimens from laminates that were testedin the four-point bend setup. For means of comparison, the MWD of as-received(non-processed), single-modified and multi-functionalized polymer samples is pre-sented in Figure 6-11.Referring to Figure 6-11 a, the MWD of B3S modified with P-EPQ after subjectionto temperature profile P2 in air results in a similar, slightly broadened MWD ofnon-processed B3S. However, Mw is reduced significantly and shifted to lower Mn

when multi-functionalized B3S was investigated. The long subjection to high tem-peratures (P2) is expected to result in a shift to lower masses as observed for multi-functionalized polymers due to stabilization reactions. Potential post-condensation

120 Thermal stabilization and flow promotion of polyamides

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

1x102 1x103 1x104 1x105 1x1060.0

0.2

0.4

0.6

0.8

1.0

W(lo

gM

)[-]

Molar mass [g/mol]

B3S_as-receivedB3S+P-EPQ_P2_AirMF-B3S_P2_Air

a

W(lo

gM

)[-]

Molar mass [g/mol]

C2000_as-receivedC2000+P-EPQ_P2_AirMF-C2000_P2_Air

b

W(lo

gM

)[-]

Molar mass [g/mol]

B3S_4PBMF-B3S_4PBc

W(lo

gM

)[-]

Molar mass [g/mol]

C2000_4PBMF-C2000_4PBd

Figure 6-11 MWD of as-received without processing, single-modified and multi-functionalizedpolymers after subjection to temperature profile P2 in air of a) B3S and b) C2000;MWD of non-modified and multi-functionalized samples extracted from four-pointbend test panels of c) B3S and d) C2000.

reactions promoted by P-EPQ but suppressed by Licowax C are assumed to causea similar MWD as found for neat B3S. A different behavior is observed for samplesbased on C2000 (Figure 6-11 b): modification with P-EPQ only and with both ad-ditives yield an almost superposable MWD. In comparison to as-received C2000,the MWD reveals lower Mw and Mn. This confirms the assumptions of stabilizingreactions without crosslinking reactions.Comparing the MWD of samples extracted from the composite samples (Figure 6-11 c, d), Mw is reduced and shifted to lower Mn when adding P-EPQ and Li-cowax C to both B3S and C2000 in comparison to non-modified polymers. Mw

of non-modified samples is higher than for multi-functionalized B3S and C2000whereas more pronounced for B3S. The larger Mw with higher Mn of non-modifiedsamples is attributed to the stabilization reaction from P-EPQ. The antioxidantsuppresses crosslinking and post-condensation reactions by reacting with emergingradicals. The given MWDs also support the outcome from four-point bend testing.

Thermal stabilization and flow promotion of polyamides 121

By suppressing crosslinking reactions, smaller Mn and lower Mw lead to a decreasein strength.

6.4.3 Influence of matrix modification on impregnation

Analogous to the effect of degradation on the impregnation progress (see subsec-tion 5.4.5), the impact by modifying B3S and C2000 with P-EPQ and Licowax Cis studied on the impregnation time by using the impregnation model developed inChapter 4. Since the rheological measurements record the development of viscosityover time, the viscosity curves can be used as input parameter for the model. Allother input parameters remained the same as used in subsection 5.4.5.In Figure 6-12, the effect on the DOI due to the increasing viscosity during thelaminate production step of non-modified B3S and multi-functionalized B3S iscompared.

a) Constant η (B3S) b) Rise in η duringprocessing (B3S)

c) Rise in η duringprocessing (MF-B3S)

Figure 6-12 Development of the DOI of CF-TP/B3S as a function of a) constant viscosity andviscosity development as measured for temperature profile P2 during the laminateproduction step of b) non-modified B3S and c) multi-functionalized B3S.

The impregnation time can be reduced by 18 % when multi-functionalized B3S (Fig-ure 6-12 c) is used during the laminate production step compared to non-modifiedB3S (Figure 6-12 b). The modification with antioxidant and lubricant yields alsoa reduction in impregnation time of 7 % when compared to a constant viscosityof 185 Pa s (Figure 6-12 a). This arises from an initial reduction in viscosity to149 Pa s of multi-functionalized B3S due to modification. Figure 6-13 compares theeffect of an increasing viscosity during the laminate production step on the DOIfor non-modified C2000 and multi-functionalized C2000.The modification with lubricant and antioxidant (Figure 6-13 c) reduces the initialviscosity to 60 Pa s at the beginning of the press cycle during the laminate produc-tion and increases more moderately than non-modified C2000 (Figure 6-13 b). Asa consequence, modification enables complete impregnation whereas non-modifiedC2000 results in a DOI of 42 % due to the rise in viscosity induced by thermo-

122 Thermal stabilization and flow promotion of polyamides

a) Constant η (C2000) b) Rise in η duringprocessing (C2000)

c) Rise in η duringprocessing (MF-C2000)

Figure 6-13 Development of the DOI of CF-TP/C2000 as a function of a) constant viscosity andviscosity development as measured for temperature profile P2 during the laminateproduction step of b) non-modified C2000 and c) multi-functionalized C2000.

oxidative degradation. Even if no degradation reactions become apparent (constantviscosity, see Figure 6-13 a) the impregnation time is reduced by 15 % when C2000is multi-functionalized.

6.5 Conclusion and implications

Substantial increases in melt viscosity due to thermo-oxidative degradation reac-tions led to the investigations of antioxidants and lubricants to increase the thermalstability and flowability of B3S and C2000.Based on requirements such as thermal stability, low volatility and chemical com-patibility, several antioxidants and lubricants were selected and compounded to B3Sand C2000 separately (single-modified). The modified polymers were subjected totemperature profiles derived from a CFRTP production process and analyzed byDSC, TGA and rheometry. The preventive phosphorous-based antioxidant P-EPQwas found to be most efficient showing an enhanced OIT measured by DSC anddecreasing mass loss measured by TGA compared to neat B3S and C2000. How-ever, results from TGA show a decreasing effectiveness of P-EPQ for long dwelltimes. The longer the dwell time of the polymers at high temperatures, the morewater is produced by condensation reactions and the more new radicals are formedthat cannot be entirely removed by P-EPQ. The amide wax Licowax C showed thelargest lubricating effect on both B3S and C2000 indicated by significant reductionsin viscosity as long as thermo-oxidative degradation has not started yet.Combining P-EPQ and Licowax C (multi-functionalized) showed synergistic ef-fects by increasing the OIT and viscosity reduction even after subjection to hightemperatures for a long time. For multi-functionalized B3S an average viscosityreduction by 12 % in air and by 28 % in nitrogen atmosphere can be expected.

Thermal stabilization and flow promotion of polyamides 123

Multi-functionalized C2000 yielded a decrease in viscosity by 50 % in air and by54 % under exclusion of oxygen.The influence of modification on mechanical properties on composite level wasanalyzed by the four-point bend test. The flexural properties in fiber directionshow comparable values for strength when powder-coated tows are produced frommulti-functionalized polymers which were pressed with different dwell times. Thetransverse strength was found to be more affected by the dwell time than the lon-gitudinal strength. B3S-based laminates showed an increase in transverse strengthwith increasing dwell time, C2000-based laminates revealed the opposite behavior.The addition of P-EPQ and lubricant to B3S and C2000 reinforced by polyamide-sized fibers yielded lower transverse flexural strength than laminates produced fromnon-modified polymers. This is attributed to the stabilization reaction of P-EPQ.The suppression of thermo-oxidative degradation resulted in less low-molecularproducts that is assumed to hinder crosslinking and branching as shown by theresults from GPC.The influence of the modification on viscosity was compared to non-modified poly-mers for the laminate production step of temperature profile P2 by using the im-pregnation model developed in Chapter 4. Due to modification with the lubricantLicowax C, the impregnation time of B3S can be reduced by 18 % because of an ini-tial reduction in viscosity. The addition of P-EPQ led to a more moderate viscosityincrease as degradation reactions can be suppressed to a certain extent. The sig-nificant increase in viscosity due to degradation prevented complete impregnationwhen non-modified C2000 was subjected to long dwell times. The modification withLicowax C and P-EPQ led to the completion of the impregnation. Additionally, theimpregnation time was decreased by 15 % compared to non-modified C2000 whenthe initial viscosity at the beginning of the laminate production remains constant.Usually, mechanical properties in fiber direction are the dominating criteria indesigning composites. Since the flexural strength in fiber direction remained unaf-fected by using multi-functionalized polymers, the use of the selected antioxidantas well as the lubricant is recommended as the viscosity can be reduced. This re-sults in an improved impregnation progress leading to a decrease in manufacturingcosts for intermediates.

7 Gradual impregnation duringproduction

The manufacture of completely impregnated intermediates is time-consuming dueto the high viscosity of most thermoplastics. This leads to high manufacturingcosts. As investigated in Chapter 6, the impregnation times may be reduced byadding lubricants resulting in an increased operational throughput. By using par-tially impregnated tapes that are manufactured with increased production ratesin a double-belt press or pultrusion equipment the manufacturing costs can bereduced, too. These tapes are intended to completely impregnate during the sub-sequent heating and consolidation processes required for CFRTP production. Thisprinciple is also known for Out-of-Autoclave (OoA) prepregs [198, 199].In this chapter, the gradual impregnation of partially impregnated tapes duringthe considered CFRTP production process is investigated. To begin with, powder-coated tows are prepared and consolidated to tapes by varying process conditionsin a double-belt press to produce unidirectional tapes (UD tapes) with three dif-ferent initial degrees of impregnation (DOIi). The produced UD tapes are stackedand consolidated in a press by using different dwell times. Some laminates are addi-tionally thermoformed to account for a common manufacturing process of CFRTP.Flexural properties of the pressed laminates are determined to evaluate the effect oftapes with different DOIi on composite properties. Using the four-point bend test,the impact on the composite properties arising from partially impregnated tapescan be comprehensively evaluated due to the combined loading in tension andcompression. In addition, the effect of the DOIi is assessed on fiber- and matrix-dominated properties by testing in and transverse to the fiber orientation. Asidefrom mechanical testing, the impregnation state before and after processing is ana-lyzed on the basis of micrographs and related to the obtained mechanical properties.In a subsequent analysis of the process costs, the manufacturing costs required toproduce differently impregnated tapes are calculated and correlated to the yieldedmechanical properties for CF-TP/B3S.Parts of the following work have been previously published in [200].

7.1 Manufacture of differently impregnated tapes

Carbon fibers with polyamide sizing (CF-TP) were coated on both sides with thealiphatic polyamide 6 (B3S) to produce powder-coated tows as described in sub-section 2.3.1. The powder-coated tows were then fed into an isobaric double-belt

125

126 Gradual impregnation during production

press of the type Heather 5 from Held Technologie GmbH to produce the differentlyimpregnated tapes. With this double-belt press, an isobaric surface pressure of upto 80 bar and processing temperatures of up to 400 °C can be realized. The typeHeather 5 possesses a work width of 1.5 m, a length of 3.4 m and seven indepen-dently controllable heated sections. The double-belt press used for the productionof differently impregnated tapes is presented in Figure 7-1 along with the appliedtemperature profile.

Material feed

1 2 3 4 5

6 7

1> 220°C

2> 300°C

5≈ 300°C

3≈ 300°C

4≈ 300°C

7< 80°C

6< 200°C Process

direction

Figure 7-1 Double-belt press with seven heated sections used to produce CF-TP/B3S tapes withdifferent DOIi; modified from [201].

By increasing the operating speed and reducing the pressure of the double-beltpress, less time and force is left for the polymer to further impregnate the carbonfibers. The temperature profile of the double-belt press was maintained while oper-ating speed and pressure were varied to yield tapes with three different DOIi. Whenthe operating speed was changed it was waited until the temperature stabilized inthe different heated sections before the actual tape production started. However,the increasing operating speeds of 4 m/min and 8 m/min caused a retardation ofthe set temperatures. While the temperature profile given in Figure 7-1 could befollowed when the speed was set to 2 m/min, the increase in speed to 8 m/min ledto a decrease of the temperature in the heating sections 2 to 5 to 280 °C. However,this is still well above the melting temperature and sufficient for impregnation.Micrographs were prepared (see subsection 2.3.4) from the produced tapes and

Gradual impregnation during production 127

analyzed by post-processing with a MATLAB® routine (see subsection 4.4.2) todetermine the DOIi. Since the DOI is governed by the transverse impregnation offiber bundles [138], at least five fiber bundles were analyzed to determine the DOIi

of the produced tapes. Representative examples of micrographs used for evaluationof DOIi are provided in Figure 7-2.

100% DOIi

80% DOIi

a)

b)

c)

90% DOIi

~70 μm

~70 μm

~70 μm

Figure 7-2 Differently impregnated CF-TP/B3S tapes with highlighted non-impregnated areasproduced in a double-belt press with a) 2 m/min at 40 bar b) 4 m/min at 5 bar andc) 8 m/min at 5 bar.

The highlighted non-impregnated areas within the fiber bundles become largerwith increasing operating speed in the double-belt press. Table 7-1 summarizes theused process settings of the double-belt press along with experimentally determinedDOIi.

Table 7-1 Double-belt press settings and yielded different DOIi of UD tapes determined frommicrographs.

Tape designationDBP Settings DOIi [%] analyzed

from micrographsOperating speed[m/min]

Pressure[bar]

Tape_100 % 2 40 99.98±0.00Tape_90 % 4 5 86.43±0.10Tape_80 % 8 81.94±0.20

128 Gradual impregnation during production

7.2 Manufacture of thermoplastic composites

7.2.1 Prediction of dwell times

The dwell time required to complete impregnation of tapes with a DOIi of 80 % and90 % in a static press was predicted by using the impregnation model developedin Chapter 4. The input parameters used for the prediction are summarized inTable 7-2. The viscosity was set to 202 Pa s, as measured in a rheometer at 260 °C(see subsection 4.5.1). In addition, Vf was set to 60 % and the pressure to 10 bar.The computed dwell times for processing in a static press were 90 s and 300 s. Theexperimentally determined DOIi of the differently impregnated tapes were used asa starting DOI to complete the impregnation during the applied dwell times.

Table 7-2 Input parameters for the prediction of dwell times to completely impregnate partiallyimpregnated UD tapes during press forming.

Input parameter ValueFiber radius rf [m] 0.0000035Initial Vf [-] 0.60Modified Kozeny constant k′

zz [-] 0.20Maximum fiber volume content V ′

a [-] 0.82Empirical constant As [Pa] 772.77Maximum fiber volume content Va [-] 0.892Initial fiber bundle height [m] 0.00026Time increment [-] 1Pressure [105 Pa] 10Viscosity [Pa s] 202

Figure 7-3 presents the impregnation progress computed by the model for partiallyimpregnated tapes made from CF-TP/B3S that are pressed for 90 s as well as300 s. Obviously, a press time of 90 s is insufficient to complete impregnation ofneither 90 % or 80 % impregnated tapes. Impregnation is found to be completedafter 122 s for 90 % impregnated tapes and after 157 s for 80 % impregnated tapes.Thus, a dwell time of 300 s is assumed to be suitable for complete and gradualimpregnation of the partially impregnated tapes during processing in a static press.Since the model captures transverse flow within fiber bundles of a single tapeply only, extended press times of 600 s and 1200 s were chosen in addition to thedwell times 90 s and 300 s. Thus, extra time is available that may be required forconsolidation phenomena such as the establishment of intimate contact, autohesionand reptation to take place between the tape plies.

Gradual impregnation during production 129

a) DOIi=80%, 90s

c) DOIi=80%, 300s d) DOIi=90%, 300s

b) DOIi=90%, 90s

Figure 7-3 Predicted final degree of impregnation (DOIf ) after press forming tapes with a DOIi

of a),c) 80 % and b),d) 90 % at a dwell time of 90 s and 300 s.

7.2.2 Test panel production

The tapes with different DOIi were stacked according to the lay-up of [011] and [012]to obtain laminates with a nominal thickness of 2 mm. The production followedthe procedure as previously described in subsection 2.3.2.Four different dwell times (90 s, 300 s, 600 s and 1200 s) were applied at a temper-ature of 260 °C while holding the pressure constant at 10 bar. Then, the laminateswere cooled to 80 °C at a cooling rate of 20 °C/min. Some of the laminates pressedfor 90 s were additionally thermoformed as described in subsection 3.2.3. Thus, thepotential impregnation progress during a typical CFRTP production is consideredwhile maintaining the flat shape of the test panels.

7.2.3 Four-point bend test

According to DIN EN ISO 14125 B [62], four-point bend tests were conducted todetermine the effect of the DOIi of UD tapes on flexural properties of the composite.Seven specimens were tested per configuration to determine the flexural modulus(Ef1 and Ef2) and flexural strength (σf1 and σf2) in and transverse to the fiberdirection. Prior to testing, the specimens were dried at 80 °C for at least 60 h

130 Gradual impregnation during production

under vacuum. The used test setup is described in subsection 6.3.2. For qualitycontrol, the fiber volume content (FVC) of each test panel was determined at threedifferent positions by acid digestion according to DIN EN 2564 B [106] at thecentral laboratory of SGL Carbon GmbH.

7.3 Influence of initial degree of impregnation onmechanical properties

7.3.1 Final degree of impregnation after production

After processing the differently impregnated tapes with various dwell times, micro-graphs served to determine the final degree of impregnation (DOIf ) of the producedtest panels. The micrographs were prepared by following the same procedure as de-scribed in subsection 2.3.4. Representative micrographs of the manufactured testpanels are compared in Figure 7-4.

100% DOIi

80% DOIi

90% DOIi

90s 300s 600s 1200s90s + TFLaminate Production - DOIfInitial state - DOIi

Figure 7-4 Micrographs of laminates produced with varying dwell times based on differentlyimpregnated tapes with highlighted non-impregnated areas if applicable.

The micrographs and the analyzed DOIf reveal complete impregnation after pro-cessing tapes with different DOIi after a dwell time of 90 s with additional thermo-forming or when the press time exceeds 600 s.Tapes with a DOIi of 80 % show non-impregnated fiber bundles after a dwell timeof 300 s while complete impregnation is observed when these tapes were thermo-formed. In general, the dwell time of the polymer above Tm and under pressure

Gradual impregnation during production 131

is longer for laminates pressed with 300 s compared to laminates pressed with 90 sthat were additionally thermoformed. However, additional impregnation progressduring thermoforming is assumed to occur in the heating phase by infra-red sources.Besides the laminate, the surrounding aluminum mold has to be heated. Comparedto the use of aluminum foils as tooling for press forming, the heating phase is longerduring thermoforming due to the use of a tool. The effect of the low pressure (ap-proximately 0.004 bar) applied to the laminate by the upper mold half duringthermoforming on the impregnation progress is negligible.With exception for the dwell time of 90 s, the processing of tapes with a DOIi of90 % with prolonged press times or thermoforming resulted in completely impreg-nated fiber bundles. Table 7-3 summarizes the different dwell times used for theproduction of test panels in a static press and via thermoforming from differentlyimpregnated tapes along with the resulting DOIf .

Table 7-3 Dwell times used to process differently impregnated tapes in a static press and viathermoforming along with the final DOIf of all test panels.

Designation oftest panels

Dwell time DOIf analyzed frommicrographs [%]static press [s] thermoforming [s]

80 %_90 s 90 - 94.51±0.0190 %_90 s 90 - 97.29±0.01100 %_90 s 90 - 100.00±0.0080 %_90 s+TF 90 8 100.00±0.0090 %_90 s+TF 90 9 99.40±0.01100 %_90 s+TF 90 8 99.81±0.0080 %_300 s 300 - 95.66±0.0190 %_300 s 300 - 100.00±0.00100 %_300 s 300 - 100.00±0.0080 %_600 s 600 - 100.00±0.0090 %_600 s 600 - 100.00±0.00100 %_600 s 600 - 100.00±0.0080 %_1200 s 120 - 100.00±0.0090 %_1200 s 1200 - 100.00±0.00100 %_1200 s 1200 - 100.00±0.00

The results from FVC and thickness measurements of all test panels can be foundin section A.3 of the appendix. Averaging the measured FVC over all test panels,a mean of 45.23±1.08 % is obtained. The thickness of test panels pressed for 600 sand 1200 s is increased as another ply ([012]) was added compared to the remaininglaminates. The consolidated tape thickness was found to decrease with increasingpress time in earlier studies. By adding another ply, an almost constant FVC isachieved for all test panels enabling comparability.

132 Gradual impregnation during production

7.3.2 Flexural properties

Longitudinal flexural properties

The effects of the DOIi of tapes on the longitudinal flexural strength σf1 and mod-ulus Ef1 of test panels produced with different dwell times in a static press andthermoforming are shown in Figure 7-5.

80%_1200s90%_1200s100%_1200s

80%_600s90%_600s100%_600s

80%_90s+TF90%_90s+TF100%_90s+TF

80%_90s90%_90s100%_90s

Dwell time [s]

σ f1[M

Pa]

80%_300s90%_300s100%_300s

90s90s+TF 300s 600s 1200s

0

200

400

600

800

1000

1200

1400

1600

Dwell time [s]

Ef1

[MP

a]

90s90s+TF 300s 600s 1200s

0

20

40

60

80

100

120

a

80%_1200s90%_1200s100%_1200s

80%_600s90%_600s100%_600s

80%_300s90%_300s100%_300s

80%_90s+TF90%_90s+TF100%_90s+TF

80%_90s90%_90s100%_90s

b

Figure 7-5 a) Longitudinal flexural strength σf1 and b) longitudinal flexural modulus Ef1 of testpanels made from CF-TP/B3S tapes with different DOIi that were processed withvarying dwell times in a static press or thermoformed (+TF).

As discussed in Chapter 3, the significantly different cooling rates during processingin a static press (up to 20 °C/min) and during thermoforming (up to 380 °C/min)were found to have negligible influence on the crystallization and hence on themechanical properties of CF-TP/B3S. Therefore, the mechanical test results fromlaminates produced in a static press can be compared to thermoformed test panelswithout limitation.Referring to short dwell times of 90 s, the longitudinal flexural strength σf1 slightlydecreases with decreasing DOIi. This goes along with incomplete impregnation

Gradual impregnation during production 133

(95-97 %) found after processing tapes with a DOIi of less than 100 % for 90 s.An overall low strength level is found for test panels pressed for 90 s comparedto laminates with extended press times. This indicates that a dwell time of 90 sis not only insufficient to achieve complete impregnation in the case of 80 %- and90 %-impregnated tapes but also deficient for consolidation between plies referringto 100 %-impregnated tapes. Extending the dwell time or with additional thermo-forming leads to similar σf1 for completely impregnated tapes and tapes with aDOIi of 90 %. For all test series, the tapes with a DOIi of 80 % yielded the lowestσf1. The longitudinal stiffness (see Figure 7-5b) is not affected by the DOIi of theused tapes, with exception of tapes with a DOIi of 80 % press formed for 90 s.

Transverse flexural properties

The transverse flexural properties for test panels produced from CF-TP/B3S tapeswith different DOIi and with different process settings are reported in Figure 7-6.

0

20

40

60

80

100

0123456789

10

80%_1200s90%_1200s100%_1200s

80%_600s90%_600s100%_600s

80%_300s90%_300s100%_300s

80%_90s+TF90%_90s+TF100%_90s+TF

σ f2[M

Pa]

Dwell time [s]

80%_90s90%_90s100%_90s

90s

90s+TF 300s

600s

1200s

a

80%_1200s90%_1200s100%_1200s

80%_600s90%_600s100%_600s

80%_300s90%_300s100%_300s

80%_90s+TF90%_90s+TF100%_90s+TF

E f2[G

Pa]

Dwell time [s]

80%_90s90%_90s100%_90s

90s

90s+TF 300s 600s 1200s

b

Figure 7-6 a) Transverse flexural strength σf2 and b) transverse flexural modulus Ef2 of lam-inates made of tapes with different DOIi that were processed with varying dwelltimes in a static press or thermoformed (+TF).

134 Gradual impregnation during production

A press time of 90 s yielded a comparable transverse flexural strength σf2 withinstandard deviation independent of the DOIi. As observed for the longitudinal flex-ural strength, σf2 of laminates pressed for 90 s remains at a lower level than testpanels that were pressed for prolonged press times or thermoformed. However, thelevel of σf2 obtained from laminates based on completely impregnated tapes is notreached when tapes with DOIi of 80 % and 90 % were pressed for 300 s or 90 s withadditional thermoforming.When pressed for 600 s, test panels made from partially impregnated tapes approx-imate the σf2 of laminates made from completely impregnated tapes. Tapes witha DOIi of 80 % and 100 % show even comparable values for σf2. When test panelswere pressed for 1200 s σf2 becomes independent of the DOIi.In contrast to the measured longitudinal properties, σf2 of test panels producedfrom tapes with a DOIi of 90 % is found to be slightly lower than of laminatesmade from tapes with a DOIi of 80 %. Slight increases in FVC for the laminatesmade from 90 % impregnated tapes may cause decreased values for σf2. Since themicrographs document complete impregnation for tapes with a DOIi of 90 % uponpressing for 300 s or thermoforming, the lower σf2 cannot be attributed to insuffi-cient impregnation. In addition, insufficient air evacuation and insufficient time forrelaxation of the fiber bed during manufacture is assumed to cause lower σf2.After pressing for 90 s, the transverse flexural modulus Ef2 (Figure 7-6b) of 80 %impregnated tapes yields significantly lower values than laminates produced fromtapes with a DOIi of 90 % or 100 %. As soon as the press time exceeds 300 s thetransverse modulus becomes independent of the DOIi of tapes.Obviously, the transverse flexural strength is more dependent on the DOIi than thelongitudinal flexural strength. In general, tapes with a DOIi of 90 % faster convergeor exceed reference values from completely impregnated tapes than tapes with aDOIi of 80 % upon processing.

7.4 Cost analysis

For process costing, the consumption of resources during processing is determinedand evaluated monetarily. So, the arising expenses, divided into cost types, are cap-tured for each process step. The cost types comprise of material costs, labor costs,overheads, direct expenses and other recurring costs [7]. Recently, process costinghas been favored over parametric, empirical-based and performance-oriented mod-els for cost accounting of components made of CFRP. This is mainly attributed tothe low amount of data required for the estimation of process-based costs and thehigh transparency of the used cost accounting type.

Gradual impregnation during production 135

7.4.1 Procedure

In the following, the theory is presented to calculate process-based costs by usingthe machine-hour-rate approach. The data collection was conducted by an expertconsultation to establish the basis for process costing. In dependance of varyingoperating speeds in a double-belt press, the machine hour rate and the throughputis calculated.The data collection and cost analysis has been conducted within the frame of thestudent project from Miriam Ernst and Patrick Consul [202].

Process costing

Within the frame of process costing, the production costs are calculated by addingmaterial costs and manufacturing costs. These were further divided into direct costsand overheads as presented in Equation 7-1. In addition, special direct costs thatcannot be attributed to material or manufacture are taken into account.

ΣProduction costs = Material costs + Manufacturing costs

= Direct material costs + Overheads

+ Direct manufacturing costs + Overheads

+ Special direct costs

(7-1)

With the presented cost analysis, the effect of partially impregnated tapes on pro-duction costs shall be evaluated. The material costs for carbon fibers, matrix as wellas the manufacturing costs for spreading the fibers, grinding the polymer pelletsand powder-coating prior to tape consolidation in a double-belt press are indepen-dent of the operating speed. Thus, the costs for powder-coated tows are excludedfrom the cost analysis. The focus lies on the direct costs and overheads for manu-facturing in a double-belt press.To begin with, the production volume was estimated to set the basis for the costanalysis. The assumptions made to calculate a realistic production volume in 2020are summarized in Table 7-4 and are based on the market report by the AVK Fed-eration of Reinforced Plastics and Carbon Composites e.V. in 2015 [3].

Table 7-4 Assumptions made for cost analysis based on forecast for 2020.

Forecast 2020 QuantitySales quantity of carbon fibers in 2020 100,000 t/aCarbon fibers used with thermoplastics (2014) 15 %Proportion of prepreg layup processes (2014) 45 %Target market share in 2020 1 %

136 Gradual impregnation during production

Based on the data presented in Table 7-4, the production volume of carbon fiberreinforced UD tapes in tons per year for 2020 was calculated as follows:

Production volume (2020) = Forecast (2020) · CFRTP share ·Prepreg layup share · Target market share

= 67.5 t/a

(7-2)

Subsequently, the manufacturing costs are determined by using the machine hourrate approach. The overheads for manufacturing is also known as factory overheadand comprise fixed maintenance costs, tooling costs, costs for auxiliary material,costs for power and occupancy. Since the required personnel to operate a double-belt press can be directly allocated to the equipment, the labor costs are alsoattributed to the direct manufacturing costs.The running time or occupation time of the double-belt press is determined bysubtracting the non-productive time from the productive time as Equation 7-3shows. The non-productive time consists of down time and test time whereas maintime, setup time, secondary time and additional time account for the productivetime.

Running time = Productive time − Non-productive time (7-3)

Related to the double-belt press, the non-productive time is composed of the setuptime, cleaning time, heating-up period and servicing period. In combination withthe shift length, the actual working time is calculated that corresponds to therunning time of the double-belt press. Based on the running time and the factoryoverhead, the machine hour rate can be determined:

Machine hour rate [e/h] = Factory overhead

Running time

= Labor costs [e/h] · Working time

Running time

+ Maintenance costs [e/h]+ Auxiliary costs [e/h]+ Equipment costs [e/h]+ Occupancy costs [e/h]+ Depreciation [e/h]+ Downtime [e/h]

(7-4)

The depreciation of the double-belt press can be calculated as follows:

Depreciation [e/a] = Initial costs [e]Recovery period [a] · Imputed interest

(7-5)

Gradual impregnation during production 137

The downtime of the equipment is evaluated by the failure probability of the util-ities, expressed as costs:

Downtime costs [e] = Failure probability [−] · Utilities [e] (7-6)

The total costs are further divided into fixed and variable costs. The fixed costscomprise of the imputed depreciation and interest, the occupancy costs and thefixed maintenance costs. These costs incur time-dependent, are independent of thecapacity utilization level as well as of the operating rate. Costs for tooling, variablemaintenance, auxiliary and equipment represent the variable costs.

7.4.2 Data collection

The feasibility of the selected operating speeds with regard to the production ofpartially impregnated tapes was proven by the previous study on the impregnationprogress during processing. Aside from this information, written questionnaireswere prepared to conduct interviews of experts to gather information about thefixed and variable costs of double-belt press equipment. The questionnaires weresent to the selected experts and complemented by telephone interviews.The experts were carefully selected from suitable companies and institutes. Forthe tape production process, experts from SGL Carbon GmbH and the researchinstitute Thüringisches Institut für Textil- und Kunststoffforschung (TITK) wereinterviewed.Experts from Sandvik TPS, Hymmen GmbH and Held Technologie GmbH wereinterviewed for data collection about double-belt press equipment. Sandvik TPSis the market leader for isochoric double-belt press, Held Technologie GmbH forisobaric double-belt press. Hymmen GmbH produces isobaric as well as isochoricdouble-belt press equipment [7]. Consulting experts from all three companies fordouble-belt press equipment enables the generation of a reliable information basis.Due to the use of questionnaires, the data collection was carried out in a systematicmanner that allows to compare the responses from the interviewed experts.The collected data from expert consultation and inquiry were allocated to thecategories labor costs and manufacturing costs as stated in Table 7-5. It is assumedthat the operation of a double-belt press requires operatives only and no engineersor office workers.The heating period is divided by the number of shifts as another heating at thebeginning of the second shift is not applicable. Considering edge trim and scrapduring the production, the actual produced volume becomes smaller. The required

138 Gradual impregnation during production

Table 7-5 Collected data used for the cost analysis.

Cost type QuantityLabor costs

Required staff [persons] 2Personnel costs [e/h] 45Number of shifts [-] 2Shift length [h] 8Working time [h/week] 35Working days per year (2016) [d/a] 253Non-productive time [h] 2.28Productive time [h] 5.72Additional charges (R&D, administration,sales) [%]

20

Manufacturing costsOccupancy costs [e/(m2month)] 2.20Power costs [e/kWh] 0.15Initial costs [e] 4,000,000Floor space required (equipment, periphery,logistics) [m2]

275

Work width [m] 1.5Amount of tapes fed into double-belt press [-] 3Tape spool length [m] 250Imputed interest [%] 5Recovery period [a] 10Additional costs (e.g. reconstruction) [e] 0Power input [kW] 500Power requirement [%] 30Operating performance (Power input *Power requirement) [kW]

150

Auxiliary materials (release agent) based on2 m/min [e/h]

25

Cleaning time [h] 0Setup time [min per spools/shift] 5Scrap [%] 2Edge trim [%] 5Heating-up period [h] 2Failure probability of equipment [%] 0Utilities (steel strip) [e] 60,000Service life (steel strip) [h] 10,833

production volume has to incorporate edge trim and scrap to produce the desiredproduction volume as Equation 7-7 presents:

Required production volume [t/a] = Production volume += Edge trim + Scrap

= 72.23 t/a

(7-7)

Gradual impregnation during production 139

Based on the required production volume the number of identical double-belt pressequipment is determined to manufacture the necessary amount of UD tapes.

Required equipment = Round up

(Required production volume

Throughput per year

)(7-8)

7.4.3 Monetary effect of partially impregnated tapes

Based on the collected data, the labor and manufacturing costs are determined forthe varying double-belt press speeds as used during the study on gradual impreg-nation. Thus, the machine hour rate and finally the manufacturing costs for theproduced partially impregnated tapes can be calculated as presented in Table 7-6.

Table 7-6 Calculation of the machine hour rate for varying operating speeds of a double-beltpress.

Cost type QuantityFixed costs

Number of required equipment 1Labor costs [e/h] 125.87Occupancy costs [e/h] 2.51Depreciation [e/h] 145.11Power costs [e/h] 77.39

Variable costsOperating speed [m/min] 2 4 8Throughput per shift [kg/shift] 196.03 392.06 784.12Throughput of kg UD Tape per hour [kg/h] 34.27 68.54 137.08Edge trim, scrap costs [e/h] 57.39 114.78 229.57Maintenance costs [e/h] 3.87 7.75 15.49Auxiliary materials and utilities [e/h] 34.97 69.93 139.86Downtime costs [e/h] 0

Machine hour rate [e/h] 447.11 543.34 735.80Manufacturing costs tape [e/kg] 13.04 7.93 5.36Reductions in manufacturing costs [%] Reference -39.18 -58.90

Obviously, the machine hour rate increases with increasing operating speed sincemore auxiliary materials are required and the incurring edge trim as well as scrapincrease. At the same time the throughput increases proportionately to the oper-ating speed. The manufacturing rate thereby decreases with increasing operatingspeed. Producing tapes with an operating speed of 4 m/min, the manufacturingcosts for UD tapes decrease by 39 % compared to tape production with 2 m/min.Increasing the operating speed further to 8 m/min the manufacturing costs reduceby 59 % compared to UD tapes that are produced with 2 m/min.

140 Gradual impregnation during production

7.5 Correlation of mechanical properties andmanufacturing costs

The influence of the different DOIi of tapes on flexural properties of test panels uponprocessing with different dwell times was previously determined during the studyon gradual impregnation. Thus, the effects on mechanical performance resultingfrom the use of partially impregnated tapes can be correlated to the incurringmanufacturing costs estimated by the cost analysis in Figure 7-7. The reference isrepresented by completely impregnated tapes that were pressed for 1200 s and isshown in the center of each of the diagrams below.

-60 -40 -20 0 20 40 60

-40

-30

-20

-10

0

10

20

30

-60 -40 -20 0 20 40 60

-50

-40

-30

-20

-10

0

10

20

30

40

-60 -40 -20 0 20 40 60

-40

-30

-20

-10

0

10

20

30

-60 -40 -20 0 20 40 60

-40

-30

-20

-10

0

10

20

30

90s90s+TF300s600s1200s

DOIi=80%

σf1 / σf1_1200s

Cos

t(DO

I i)/C

ost(D

OI i=

100%

) 90s90s+TF300s600s1200s

a

DOIi=80%

DOIi=100%

DOIi=90%

σf2 / σf2_1200sC

ost(D

OI i)

/Cos

t(DO

I i=10

0%) b

DOIi=100%

DOIi=90%

90s90s+TF300s600s1200s

Ef1 / Ef1_1200s

Cos

t(DO

I i)/C

ost(D

OI i=

100%

)

DOIi=100%DOIi=90%

DOIi=80%

c 90s90s+TF300s600s1200s

Ef2 / Ef2_1200s

Cos

t(DO

I i)/C

ost(D

OI i=

100%

)

DOIi=80%DOIi=90%

DOIi=100%

d

Figure 7-7 Correlation of costs to mechanical performance for a) σf1, b) σf2, c) Ef1 and d) Ef2compared to reference values obtained from completely impregnated tapes pressedfor 1200 s.

Considering the strength in fiber direction as potential design criteria, the manu-facturing costs of UD tapes can be decreased by 39 % when a loss of about 11 % isacceptable after thermoforming or a press time of 300 s. Increasing the dwell time

Gradual impregnation during production 141

to 600 s and 1200 s, only a minor loss of 5 % and 2 % is to be expected. Less than10 % reduction in σf1 associated with cost savings in tape manufacturing of 59 %are achieved when tapes with a DOIi of 80 % are pressed for 1200 s.As previously mentioned, the transverse flexural strength is affected to a greaterextent than the longitudinal flexural strength. A loss in mechanical performance ofabout 10 % may be acceptable when considerable savings in manufacturing costsare achieved at the same time. When pressing tapes with a DOIi of 80 % and 90 %for 1200 s, less than 10 % loss in σf2 can be expected while reducing tape costs by39 % to 56 %.If stiffness in fiber direction is meant to be the dominating design allowable re-ductions of less than 5 % have to be approved resulting from the use of partiallyimpregnated tapes. Transverse to the fiber direction, a moderate loss in stiffnesscan be expected when partially impregnated tapes are pressed for at least 600 s.

7.6 Conclusion and implications

In a comprehensive study, the influence of varying DOIi on flexural composite prop-erties was tested. UD tapes with a DOIi of 80 %, 90 % and 100 % were produced byvarying the operating speed and pressure in a double-belt press. Test panels madeof these tapes were produced in a static press with press times from 90 s up to1200 s. Some laminates were also thermoformed to account for a typical productionprocess for CFRTP components.The impregnation state before (DOIi) and after processing the partially impreg-nated tapes (DOIf ) was analyzed from micrographs. The micrographs as well asthe obtained DOIf document a further impregnation progress during CFRTP pro-duction.To enable the evaluation of effects of DOIi on composite properties, four-pointbend tests were conducted for all laminates. A dwell time of 300 s or 90 s withadditional thermoforming was sufficient to complete impregnation of tapes with aDOIi of 90 % confirmed by the results for the longitudinal flexural strength σf1.Tapes with a DOIi of 80 % appeared to be completely impregnated after press-ing for 90 s with additional thermoforming. Exceeding the press time of 600 s, thelongitudinal strength became independent of the DOIi. Stiffness in fiber directionwas found to be less affected by the DOIi of the used tapes and shows comparablevalues for all test panels upon a dwell time of 300 s or thermoforming.The flexural properties transverse to the fiber direction are affected by the DOIi

to a stronger extent indicating that dwell times need to be increased to enablefiber bed relaxation and sufficient consolidation of tape plies with a DOIi of lessthan 100 %. σf2 became independent of the DOIi upon pressing for 1200 s. How-

142 Gradual impregnation during production

ever, 80 % and 100 % impregnated tapes yielded comparable values for transversestrength when pressed for 600 s. Stiffness in transverse direction was found to fur-ther increase with extending dwell time while it became independent of DOIi assoon as the press time exceeds 300 s.The loss in mechanical performance was opposed to the monetary impact by in-creased operating speeds in the double-belt press on the manufacturing costs ofUD tapes. The monetary effect was evaluated by a cost analysis according to theprocess costing approach. Based on a comprehensive data collection by expert con-sultation, the machine hour rate and the throughput of a double-belt press wascalculated in dependance on various operating speeds.When tapes with a DOIi of 80 % are produced, the manufacturing costs are re-duced by 59 %. Producing tapes with a DOIi of 90 % cost savings of up to 39 %can be expected. Considering the flexural properties of the produced compositescost savings of 39 % are possible when a loss in mechanical performance between 2and 5 % are acceptable. Higher reductions in manufacturing costs of 59 % go alongwith mechanical properties that are found to be decreased by 5 to 10 %.

8 Summary and outlookPre-impregnated and pre-formed intermediates typically used for the productionof CFRTP components are characterized by high material costs. They arise fromthe time-consuming impregnation step due to the high melt viscosity of most ther-moplastics. In this work partially impregnated tapes were developed and manu-factured with increased production rates to reduce the costs of intermediates. Thepartially impregnated tapes were intended to completely impregnate during subse-quent heating and consolidation processes required to produce a final componentmade from CFRTP.

8.1 Summary and conclusion

Fiber-matrix compatibility

To obtain high-performance composites with ideal load transfer between the fibers,the fiber-matrix compatibility between various polyamide types and carbon fiberswith different sizings was studied. Due to the high level of comparability and flex-ibility, the powder-coating method was chosen for the production of the variousmaterial combinations made from two carbon fiber types with a polyamide-basedsizing (CF-TP) as well as an epoxy-compatible sizing (CF-EP) and four polyamidetypes. The high sensitivity of the transverse four-point bend test as well as thedouble-cantilever beam test allowed to compare the different material combinationswith regard to the fiber-matrix adhesion. In relation to epoxy-sized fibers, the useof carbon fibers with polyamide-based sizing yielded a considerable increase of thetransverse flexural strength by 152 % and a moderate gain in interlaminar fracturetoughness of 11 % when combined with an aliphatic PA6 (B3S). These observa-tions were also supported by fracture analysis using SEM. Nano-indentation testson CF-EPY/B3S and CF-TP/B3S revealed the establishment of an interphase withincreased hardness on micro-mechanical level.The use of a toughened PA6 grade (B3L) and a PA6 grade with high molecularweight (B40) led to additional improvements in fracture toughness by 72 % and55 % when combined with polyamide-sized carbon fibers. However, B3L led to adecrease in transverse flexural strength compared to B3S and the high viscosity ofB40 was found to be detrimental to time-efficient impregnation.Using the co-polyamide C2000 as matrix, a decrease by 23 % in transverse flexuralstrength and an increase by 44 % in interlaminar fracture toughness was obtainedwhen polyamide-compatible carbon fibers were used. In this case, the insufficientspreading behavior of CF-TP fibers was found to be responsible for the decreased

143

144 Summary and outlook

transverse flexural strength whereas the fracture analysis showed improved adhe-sion to polyamide-sized carbon fibers.As a result from mechanical testing and fracture analysis, the material combina-tions CF-TP/B3S and CF-TP/C2000 were selected for further investigations.

Crystallization behavior

Since processing in a static press and thermoforming were considered as produc-tion processes to investigate gradual impregnation during CFRTP production, thesignificantly different cooling rates have to be taken into account when mechanicalproperties of semi-crystalline CFRTP are compared. Therefore, the non-isothermalcrystallization behavior of B3S and C2000 was studied on polymer, intermediate,as well as on the composite level by using DSC.Neat B3S revealed the maximum crystallization at a cooling rate of 5 °C/min withdecreasing crystallization ability for increasing cooling rates from 2 to 50 °C/min.In contrast, neat C2000 showed the maximum crystallization at a cooling rate of20 °C/min and a rapid decline of crystallinity towards larger cooling rates up to50 °C/min.Analyzing the influence of different carbon fiber sizings on the crystallization be-havior of B3S and C2000 a strong nucleating behavior of carbon fibers with anepoxy-based sizing on both polymers was revealed. A new method to describe theamount of crystallinity in fiber reinforced thermoplastics, the Crystallinity RatioCR, proved to be a suitable method to describe and compare the crystallized frac-tion in carbon fiber reinforced B3S as well as C2000.To evaluate the effect of cooling rates on composite level, spread polyamide-sizedcarbon fibers were coated with B3S and C2000 (powder-coated tows), stackedand processed in a static press with three distinct cooling rates of 2, 20 and>300 °C/min. The produced test panels were subsequently evaluated by transversefour-point bend testing. For B3S-based composites, only slight differences in theamount of crystallinity were detected, leading to comparable mechanical proper-ties over a large range of cooling rates. In C2000-based composites, different coolingrates yielded large differences in the amount of crystallinity. However, the flexuralproperties were not affected within the same extent.In addition to non-isothermal crystallization, the isothermal crystallization kineticswas investigated on neat polymers to identify the temperatures where the crystal-lization proceeds most rapidly and comprehensively. These temperatures may beused upon processing during cooling to maximize crystallization or to reduce resid-ual stresses.

Summary and outlook 145

Impregnation model

Based on literature review, an impregnation model was developed to enable simula-tion of the transverse impregnation progress for various intermediates. The actualresin flow was modeled by using Darcy’s law and the temperature dependance ofthe viscosity by the Arrhenius relationship.An impregnation study was conducted to verify the developed model experimen-tally. Following the Box-Behnken method, DOE was used to design this experimen-tal study. The main process parameters that drive impregnation - time, temperatureand pressure - were varied during the production of test panels from powder-coatedtows.By post-processing micrographs from the produced test panels with a MATLAB®

routine, an experimental procedure was developed to determine the DOI. In a sub-sequent step, the results obtained by the model were compared to experimental datafrom the impregnation study. The model showed good agreement to experimentaldata for both CF-TP/B3S and CF-TP/C2000. The overall error was approximately5 %.In addition to micrographs, the test panels produced during the impregnation studywere used to conduct ILSS tests. This test method was found suitable to pro-vide information about the impregnation progress of CF-TP/B3S as well as CF-TP/C2000, since the results from ILSS tests and DOI showed comparable trendsfor various process parameter settings.In general, the developed model serves the determination of required processing pa-rameters to complete impregnation of powder-coated tows and tapes in dependanceof the DOI. The model can also incorporate the initial DOI of various intermedi-ates and viscosity data as a function of time to predict process parameters thatare required to complete impregnation of powder-coated tows or consolidated tapesduring subsequent processing. In addition, the model allows for the investigationof effects of viscosity changes, arising from applied temperature, degradation ormodification by additives on the impregnation time.

Degradation of polyamides

To enable gradual impregnation of carbon fibers during subsequent heating andconsolidation processes as present during the production of thermoplastic compos-ites, the effects induced by thermal cycling on the matrix were investigated. Thederived temperature profiles with different dwell times follow a typical CFRTP pro-duction process starting with powder-coating, consolidation to tapes in a double-belt press, consolidation to multi-ply laminates and thermoforming of multi-plylaminates to a final component. DSC, TGA and rheometry were used to gener-ate thermal cycling and analyze its effect on neat B3S as well as on C2000. The

146 Summary and outlook

results obtained from all three methods led to a more comprehensive understand-ing of degradation reactions that arise during thermal cycling. The results fromDSC and TGA indicated thermo-oxidative degradation of both B3S and C2000. Inrheological studies, significant increases in viscosity for both B3S and C2000 weredetected, when exposed to elevated temperatures in presence of oxygen as a re-sult from secondary reactions of thermo-oxidative degradation. No correlation wasfound between the increases in viscosity and changes in the number-average molec-ular weight Mn and weight-average molecular weight Mw as measured by GPC.The effect of increasing viscosity on the impregnation progress of both B3S andC2000 was simulated by using the model developed in Chapter 4. The increase inviscosity of B3S led to a rise in impregnation time by 13 % compared to the viscos-ity of a non-degraded B3S. The effect of thermo-oxidative degradation on furtherprocessing was found to be more pronounced for C2000. While the viscosity ofnon-degraded C2000 would yield complete impregnation after 1700 s, the increas-ing viscosity, as measured on C2000 during the laminate production step, hinderscomplete impregnation leading to a DOI of 42 %.The number of thermal cycles was found to have a negligible effect on thermo-oxidative degradation of B3S and C2000 when dwell times are limited to 5 min-utes. Considering the concept of gradual impregnation throughout the productionof thermoplastic composites, a processing window of 5 minutes per process stepwas therefore identified for both B3S and C2000. The identified processing windowdescribes the time until the viscosity starts to increase and hence is a conservativevalue that still allows some impregnation progress.

Thermal stabilization and flow promotion

The substantial increases in melt viscosity due to thermo-oxidative degradationreactions led to the investigations of antioxidants and lubricants to increase thethermal stability as well as the flowability of B3S and C2000.Based on requirements such as thermal stability, low volatility and chemical com-patibility, several antioxidants and lubricants were selected and compounded toboth polyamide types separately. Single-modified B3S and C2000 were subjected tothe temperature profiles previously derived from a CFRTP production. The analy-sis was conducted by using DSC, TGA and rheometry. The preventive phosphorous-based antioxidant P-EPQ was found to be most efficient, showing an enhanced OITmeasured by DSC and decreasing mass loss measured by TGA compared to neatB3S and C2000. However, results from TGA showed a decreasing effectiveness ofP-EPQ for long dwell times. The amide wax Licowax C showed the largest lubricat-ing effect on both B3S and C2000, indicated by significant reductions in viscositybefore thermo-oxidative degradation starts.Combining P-EPQ and Licowax C showed synergistic effects by increasing the

Summary and outlook 147

thermal stability and reducing the viscosity even after subjection to elevated tem-peratures for considerable time. For B3S modified with both additives, an averageviscosity reduction of 12 % in air and 28 % in nitrogen atmosphere can be expected.C2000 modified with P-EPQ and Licowax C yielded a decrease in viscosity by 50 %in air and 54 % under exclusion of oxygen.The influence of the modification on the impregnation behavior was compared tonon-modified polymers by using the model developed in Chapter 4. By adding thelubricant Licowax C, the impregnation time of B3S can be reduced by 18 % be-cause of an initial reduction in viscosity. In combination with P-EPQ, the increasein viscosity proceeded more moderate as degradation reactions can be suppressedto a certain extent. When non-modified C2000 was subjected to long dwell timesthe resulting increase in viscosity due to degradation prevented complete impreg-nation. By modifying C2000 with Licowax C and P-EPQ complete impregnationwas enabled.The four-point bend test was used to determine the influence of the polymer mod-ification on the mechanical properties on composite level. The transverse strengthwas found to be more affected by the dwell time than the longitudinal strength.The addition of P-EPQ and lubricant to B3S and C2000 reinforced by polyamide-sized fibers yielded lower transverse flexural strength than laminates produced fromnon-modified polymers. This effect is attributed to the stabilization reaction of P-EPQ as the suppressed thermo-oxidative degradation resulted in less low-molecularproducts. Thus, crosslinking and branching are prevented as shown by the resultsobtained from GPC.In the design of composites, the mechanical properties in fiber direction are usu-ally the dominating design criteria. As the longitudinal flexural strength remainedunaffected by the selected antioxidant as well as the selected lubricant their use isrecommended as the viscosity can be reduced. An improved impregnation progressdue to the use of modified polymers can be expected leading to a decrease inmanufacturing costs for intermediates.

Gradual impregnation

In a comprehensive study, the influence of varying DOIi of tapes on composite prop-erties after processing to CFRTP components was tested. UD tapes with a DOIi of80 %, 90 % and 100 % were produced by varying the operating speed and the pres-sure in a double-belt press. With completely impregnated tapes as a reference, thepartially impregnated tapes were processed by press forming and thermoformingwith different dwell times from 90 s up to 1200 s to simulate a typical production ofa CFRTP component. As discussed in Chapter 3 significantly different cooling ratesas present in a static press (20 °C/min) or during thermoforming (>300 °C/min)were found to have negligible influence on the crystallization and on flexural prop-

148 Summary and outlook

erties of CF-TP/B3S. Thus, complete comparability between test panels that werepressed in a static press or thermoformed is given.The experimental method to determine the DOI developed in Chapter 4 was usedto analyze the impregnation state before (DOIi) and after press forming or thermo-forming (DOIf ). The post-processed micrographs from the test panels documentgradual and complete impregnation of the partially impregnated tapes upon pressforming them for more than 600 s.The effects of tapes with different DOIi on the composite performance was eval-uated by longitudinal and transverse four-point bend tests. A dwell time of 300 sor 90 s with additional thermoforming was sufficient to complete impregnation oftapes with a DOIi of 90 %, confirmed by the results for the longitudinal flexuralstrength. Tapes with a DOIi of 80 % appeared to be completely impregnated af-ter processing in a static press for 90 s with additional thermoforming. After pressforming the laminates for more than 600 s, the longitudinal flexural strength be-came independent of the DOIi with completely impregnated tapes as a reference.The flexural properties transverse to the fiber direction were found to be affectedby the DOIi to a stronger extent, indicating that dwell times need to be increasedto enable fiber bed relaxation and sufficient consolidation between the plies of par-tially impregnated tapes. The transverse flexural strength became independent ofthe DOIi after processing for 1200 s compared to completely impregnated tapes.With a process-based cost analysis, the effect of partially impregnated tapes onproduction costs was evaluated. Based on a comprehensive data collection by ex-pert consultation, the machine hour rate as well as the throughput of a double-beltpress depending on various operating speeds was calculated. The analysis of themanufacturing costs yielded cost savings between 39 to 59 % as the tapes with aDOIi of 90 % can be produced with double and tapes with a DOIi of 80 % withquadruple production rates, compared to completely impregnated tapes.

8.2 Future work

Considering the application of partially impregnated tapes, the feasibility of pro-cessing in a static press or in a thermoforming unit was proven for polyamide 6.The gradual impregnation during the CFRTP production has to be verified alsofor other types of thermoplastics first before this concept can be established inindustry.In addition, the use of partially impregnated tapes in back-injection molding pro-cesses where they can be heated during the injection of long-fiber reinforced orneat thermoplastics may represent another application area. In this case, the cost-efficient production of partially impregnated tapes can meet the requirements for

Summary and outlook 149

automotive applications. The use of partially impregnated tapes in combinationwith AFP or filament winding can come into consideration for lay-up processeswith subsequent consolidation in a press or an autoclave. However, if in-situ consol-idation is desired partially impregnated tapes may not be completely impregnateddue to the short processing times and low consolidation forces.The use of partially impregnated tapes in the future requires the development ofsuitable criteria to evaluate the quality of the tapes. The non-impregnated areasmay be described by specifications with regard to width and height across the cross-section of the tapes. Close tolerances and exact processing guidelines can guaranteethat partially impregnated tapes are completely impregnated after processing themto a final component. Beyond, suitable non-destructive test methods need to betested and further developed to ensure the detection of the non-impregnated areas.For future process designs, thermoplastic matrix modified with antioxidant andlubricant may be used and processed with increased operating speeds in a double-belt press. Considering tapes with a DOIi of 90 % sufficient flow promotion of thematrix may enable complete impregnation after production in a double-belt presswhile the operational output as well as the economic attractiveness is increased.The reduction of the manufacturing costs for intermediate materials can representa key to an increased acceptance for thermoplastic composites.

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[194] H. Domininghaus, P. Elsner, P. Eyerer, and T. Hirth, eds., Kunststoffe: Eigen-schaften und Anwendungen. VDI-Buch, Dordrecht: Springer, 2011.

[195] M. Bonnet, Kunststofftechnik: Grundlagen, Verarbeitung, Werkstoffauswahlund Fallbeispiele. Wiesbaden: Springer Vieweg, 2014.

[196] Clariant International Ltd, “Waxes for engineering thermoplastics.”https://www.clariant.com/de/Solutions/Products/2014/03/18/16/33/-Licowax-C-powder, 2013.

[197] DIN EN 728, “Rohre und Formstücke aus Polyolefinen - Bestimmung derOxidations-Induktionszeit,” 03/1997.

[198] T. Centea and P. Hubert, “Modelling the effect of material properties andprocess parameters on tow impregnation in out-of-autoclave prepregs,” Com-posites Part A: Applied Science and Manufacturing, vol. 43, no. 9, pp. 1505–1513, 2012.

[199] T. A. Cender, P. Simacek, and S. G. Advani, “Resin film impregnation infabric prepregs with dual length scale permeability,” Composites Part A:Applied Science and Manufacturing, vol. 53, pp. 118–128, 2013.

[200] V. Radlmaier, H. Baumann, P.-V. Brudzinski, A. Erber, H. Koerber, andK. Drechsler, “Effects of differently impregnated carbon fiber reinforced tapeswith polyamide 6 on further processing: Proceedings of the 17th ECCM,Munich, Germany,” 26-30 June 2016, pp. 1–8.

[201] Held Technologie GmbH, “Double-belt press Heather 5.”http://www.umformtechnik.net/binary_data/3190474_114498_BL-15-

168 Bibliography

04_088xx0415bl_1.jpg, 2015.

[202] M. Ernst and P. Consul, Prozess- und Kostenanalyse in der Herstellung end-loscarbonfaserverstärkter Thermoplaste. Student project, Technische Univer-sität München, Munich, Germany, 2016.

A AppendixA.1 to Section 4.4.3

Table A-1 Three-factor Box-Behnken Design for CF-TP/B3S

Standard order Run order Pressure [bar] Temperature [°C] Time [s]15 1 17.5 270 5.53 2 5 280 5.58 3 30 270 1013 4 17.5 270 5.55 5 5 270 14 6 30 280 5.52 7 30 260 5.59 8 17.5 260 110 9 17.5 280 112 10 17.5 280 1014 11 17.5 270 5.511 12 17.5 260 107 13 5 270 101 14 5 260 5.56 15 30 270 1

169

170 Appendix

Table A-2 Three-factor Box-Behnken Design for CF-TP/C2000

Standard order Run order Pressure [bar] Temperature [°C] Time [s]12 1 17.5 300 610 2 17.5 300 15 3 5 290 18 4 30 290 69 5 17.5 280 111 6 17.5 280 64 7 30 300 3.514 8 17.5 290 3.52 9 30 280 3.56 10 30 290 17 11 5 290 615 12 17.5 290 3.53 13 5 300 3.513 14 17.5 290 3.51 15 5 280 3.5

Appendix 171

A.2 to Section 6.4

Table A-3 FVC averaged over three samples of test panels produced from non-modified andmulti-functionalized B3S and C2000 reinforced by CF-TP carbon fibers

Test panel FVC [%] Standard deviation [-]CF-TP/B3S - Half time 49.0 1.2CF-TP/B3S - Full time 48.6 0.6CF-TP/MF-B3S - Half time 53.7 1.9CF-TP/MF-B3S - Full time 53.5 0.9CF-TP/C2000 - Half time 48.7 2.0CF-TP/C2000 - Full time 48.7 1.0CF-TP/MF-C2000- Half time 46.9 0.6CF-TP/MF-C2000 - Full time 49.1 2.8

A.3 to Section 7.3.1

Table A-4 FVC averaged over three samples of test panels produced with various dwell timesfrom differently impregnated tapes

Test panel FVC [%]Standarddeviation[-]

Thickness[mm]

Standarddeviation[-]

80 %_90 s 45.2 1.4 2.26 0.0390 %_90 s 44.4 1.8 2.02 0.02100 %_90 s 45.0 0.7 2.02 0.0280 %_90 s + TF 45.4 1.8 2.08 0.0290 %_90 s + TF 46.1 2.8 2.04 0.04100 %_90 s + TF 44.8 1.8 2.01 0.0480 %_300 s 44.1 2.2 2.08 0.0290 %_300 s 45.6 3.2 2.04 0.05100 %_300 s 43.4 1.4 2.02 0.0380 %_600 s 44.8 0.7 2.24 0.0590 %_600 s 45.5 0.5 2.20 0.05100 %_600 s 46.0 1.7 2.20 0.0380 %_1200 s 48.9 0.7 2.21 0.0490 %_1200 s 47.5 3.4 2.15 0.06100 %_1200 s 46.1 0.8 2.16 0.02

B PublicationsJournal papers

[P1] Radlmaier, V., Heckel, C., Winnacker, M., Erber, A., Koerber, H. “Effectsof thermal cycling on polyamides during processing”, Thermochimica Acta,vol. 648, pp.44-51, 2017.

[P2] Radlmaier, V., Obermeier, G., Ehard, S., Kollmannsberger, A., Koerber, H.;Ladstätter, E. “Interlaminar fracture toughness of carbon fiber reinforcedthermoplastic in-situ joints”, AIP Conference Proceedings, vol. 1779, no.090003, 2016.

[P3] Radlmaier, V., Ehard, S., Ladstätter, E., Drechsler, K. “Thermoplastic auto-mated fiber placement - current fields of application and future prospects”,JEC Composites Magazine, 97, pp. 39-45, 2015.

Conferences

[C1] Radlmaier, V., Henne, F., Ladstätter, E., “Novel environmentally friendlymanufacturing approach for thermoplastic composite helicopter doors”, inGreener Aviation 2014 Conference, Brussels, 12-14 March 2014.

[C2] Radlmaier, V., Erber, A., Brudzinski, P.-V., Koerber, H., Drechsler, K., “In-fluence of different sizing on fracture toughness and flexural properties ofcarbon fiber reinforced polyphthalamide”, in Proceedings of the 20th ICCM,Copenhagen, Denmark, 19-24 July 2015.

[C3] Radlmaier, V., Baumann, H., Brudzinski, P.-V., Erber, A., Koerber, H.,Drechsler, K., “Influence of differently impregnated carbon fiber reinforcedtapes with polyamide 6 on further processing”, in Proceedings of the 17thECCM, Munich, 26-30 June 2016.

173

C Supervised student thesesDuring my employment at the Chair of Carbon Composites –Lehrstuhl für CarbonComposites – I supervised the following student theses:

[S1] J. Specks, Diploma thesis, restricted until 2015.

[S2] T. Leipnitz, “Mikroskopische Untersuchungen an faserverstärkten Kunststof-fen im wissenschaftlichen Betrieb”, Student thesis, 2014.

[S3] M. Juretko, “Rheologische Charakterisierung von Polyamid 6 und Polyph-thalamid”, Diploma thesis, 2014.

[S4] C. Heckel, Master’s thesis, 2015, restricted until 2018.

[S5] A. García López, “Characterization of carbon fiber reinforced PPS jointsproduced in situ by Automated Fiber Placement”, Student thesis, 2015.

[S6] H. Baumann, Bachelor’s thesis, 2015, restricted until 2018.

[S7] S. Ender, “Literaturrecherche: Herstellungsmethoden für unidirektional car-bonfaserverstärkte Tapes mit thermoplastischer Matrix und anschließenderProzessanalyse”, Student thesis, 2015.

[S8] P. Consul, Student thesis, 2016, restricted until 2019.

[S9] E. Lell, Bachelor’s thesis, 2016, restricted until 2019.

[S10] P. Consul and M. Ernst, “Prozess- und Kostenanalyse in der Herstel-lung endloscarbonfaserverstärkter Thermoplaste”, Study in cooperation withLehrstuhl für Technologie- und Innovationsmanagement, 2016.

Parts of the following theses contributed to the present doctoral thesis as indicatedin the text: [S4], [S6], [S8] and [S10].

175