An efficient approach to the synthesis of a calcium phosphate bone-cement and its reinforcement by hydroxyapatite crystals of various particle morphologies

UNIVERSIDADE DE SANTIAGO DE COMPOSTELA INSTITUTO DE CERÁMICA DE GALICIA AN EFFICIENT APPROACH TO THE SYNTHESIS OF A CALCIUM PHOSPHATE BONE-CEMENT AND ITS REINFORCEMENT BY HYDROXYAPATITE CRYSTALS OF VARIOUS PARTICLE MORPHOLOGIES Memoria presentada por Da. María Inés Sánchez Neira para optar al grado de Doctor por la Universidad de Santiago de Compostela. Junio 2008 3 ACKNOWLEGMENTS En primer lugar, quiero agradecer al Prof. Francisco Guitián Rivera la oportunidad de trabajar bajo su supervisión en el Instituto de Cerámica de Galicia. Gracias por los consejos, comentarios y apoyo que me han guiado a lo largo de estos años. I would like to thank Prof. Masahiro Yoshimura for his kindly acceptance of my stay in his research group at the Tokyo Institute of Technology (Materials and Structures department, Japan). His knowledge about hydrothermal technique and his will to share it made this work possible. Thank you to Prof. Peter Fratzl and Dr. Himadri S. Gupta for their gently agreement for my research stay at the Biomaterials department of the Max Planck Institute of Colloids and Interfaces (Potsdam, Germany). Their understanding of bone and the share of their knowledge improve much this research and open a new biomimetic world for me. Thank you to Prof. Gustav Van Tendeloo and Dr. Oleg I. Lebedev from the Electron Microscopy for Materials Science (University of Antwerp, Belgium) for their help with transmission electron microscopy studies. Gracias a Mar, Jose, Eva, Gabriel, María y Déborah por los buenos ratos compartidos dentro y fuera del Instituto. ¡Uno siempre necesita oxigenarse!. Gracias por vuestra amistad. Gracias también a Pablo por su ayuda con el porosímetro y a Alfredo y Jaime por sus enseñanzas sobre los aparatos del Instituto. Gracias a Mercedes por ser como es. En general, gracias a todos los compañeros, a los que todavía están en el Instituto y a los que ya no. Thank you also to the members of Yoshimura-sensei’s lab. Thanks to Gusan, Yury and Machan for the “Creative company”. For the share of bin-birru and long funny crazy nihon-jin nights. Machida!. Thank you to Watanabe-san, Suda, Naota, Oga-chan, Kudo, Wagata, Yuzaburo, Sakamoto-san, Ruwan-san, Li-san, Ishigaki (Sacho!), Nakagawa and Subramani to make me feel like at home so far from home. Thank you to Matsushita-sensei for his kindest and to Yoshioka-san for her amazing efficiency. You all made my stay in Japan an incredible experience that I will never forget. The lab life (dinners, parties, travels,…) was just perfect. Thanks for make from me “Ine-chan”. Arigato gozaimasu. Thanks also to Sugata, Bidisa, Neli, Plamen and Vitor for the home parties, the discussions about Nihon, Nihon-jins and “the meaning of life”. Spopa!.Ummm… Gracias a María, Enrique e Irene por el toque español en Japón. 4 Thanks too to the members of the Biomaterials department at Potsdam. Thank you to Christoph, Claudia and Anna for the share lunches and coffees. Thanks to Maxim for the discussions on the train and at the office. Thanks to Jan and Antje for their friendly kindness. Thanks to Annemarie for her amazing preparation of my samples and to Petra for show me all what I need to know about the nanoindentor. Gracias a Josefina por su argentino soplo de aire fresco. ¡Está tan bien poder olvidarse del inglés de vez en cuando!. ¡Y por las discusiones sobre mecánica!. Thanks to Kirill, Julia, Peter, Tom, Almudena, Rosa and Carine for make Berliner life more funny and party-full. Gracias a Iria, Carba, Raquel, Juan, Jorge y Paco por vuestra amistad continuada a lo largo de los años, ¡estemos donde estemos cada uno!. Gracias a mi familia por su apoyo constante y cariño. Спасибо Коленькам за то, что приняли меня в семью. Самое главное, большое спасибо моему любимому Юре. Спасибо за его помощь с диссертацией и за все хорошее. Спасибо за то, что связал свою жизнь со мной. Без тебя ничего этого бы не было. Я тебя люблю. This work was partially supported by the Spanish Ministry of Education and Science with a Formación de Personal Investigador fellowship (project MAT2002-03857) and the European Union with a Marie Curie Early Stage Training fellowship on Biomimetic Systems (MEST-CT-2004-504465). ありがとう, dankeschön, Спасибо. MUCHAS GRACIAS A TODOS TABLE OF CONTENTS ACRONYMES............................................................................................................................i RESUMEN................................................................................................................................iii I.- RESEARCH MOTIVATION AND BACKGROUND......................................................1 1.1.- MOTIVATION.............................................................................................................3 1.2.- BACKGROUND ..........................................................................................................5 II.- INTRODUCTION ..............................................................................................................7 2.1.- BONES .........................................................................................................................9 2.1.1. Bone constituents and structural hierarchy ...........................................10 2.1.1.1. Bone structure .......................................................................... 10 2.1.1.2. Bone composition: collagen and hydroxyapatite. Woven and lamellar bone. ..................................................................... 12 2.1.2. Microarchitecture .................................................................................. 14 2.1.2.1 Trabecular bone.........................................................................15 2.1.2.1 Cortical bone ............................................................................. 16 2.1.3. Bone shape and function ....................................................................... 17 2.1.3.1. Long bones...............................................................................18 2.1.3.2. Short bones...............................................................................18 2.1.3.3. Flat bones .................................................................................18 2.1.3.4. Other geometries ...................................................................... 19 2.1.4. Bone remodelling.................................................................................. 19 2.1.5. Bone microdamage. Microcracks..........................................................20 2.1.6. Bone fracture......................................................................................... 20 2.2.- TEETH........................................................................................................................ 21 2.2.1. Structure ................................................................................................ 21 2.2.1.1. Enamel......................................................................................22 2.2.1.2. Dentine ..................................................................................... 23 2.2.1.3. Cementum ................................................................................23 2.2.2. Mechanical properties ...........................................................................24 2.2.3. Hydroxyapatite in teeth......................................................................... 25 2.3.- BONE REPAIR AND SUBSTITUTION...................................................................26 2.3.1. Bone defects..........................................................................................26 2.3.2. Natural bone grafts. Autogeneous, allogeneous and xenogeneous implants................................................................................................. 26 2.3.3. Synthetic bone grafts: Biomaterials ......................................................28 2.3.3.1. What is a biomaterial?..............................................................28 2.3.3.2. Hystorical background .............................................................28 2.3.3.3. Requirements of biomaterials. Biocompatibility, bioactivity and bioresorbability..................................................................29 2.3.3.4. Biomaterials for bone repair.....................................................31 2.3.4. Metallic biomaterials.............................................................................32 2.3.5. Polymeric biomaterials..........................................................................34 2.3.6. Ceramic biomaterials.............................................................................36 ii NI – Nano-indentation; OCP – Octacalcium phosphate Ca8(HPO4)2(PO4)4·5H2O; PMMA – Polymethylmethacrylate; PSD – Particle size distributions; PTFE – Polytetrafluoroethylene; SBF – Acellular simulated body fluid; SEM – Scanning electron microscopy; S/L ratio – Solid to liquid ratio (g/mL); ST – Final setting time; TCP – Tricalcium phosphate Ca3(PO4)2; TEM – Transmission electron microscopy; TGA/DTA – Thermo-gravimetric and differential thermal analyses; TTCP – Tetracalcium phosphate Ca4(PO4)2O; XRD – Powder X-ray diffraction; Resumen iii RESUMEN Motivación Los huesos humanos fracturados tienen la capacidad de autorrepararse a través de un proceso natural. Sin embargo, existen muchas situaciones en las que las fracturas no cicatrizan de forma apropiada o espontánea. Es por ello que, hoy en día, la reparación y sustitución de huesos mediante el uso de implantes fabricados con diversos materiales artificiales son prácticas médicas corrientes. Estos materiales artificiales empleados para reemplazar una parte o una función del cuerpo humano de un modo seguro, fiable, económico y psicológicamente aceptable, reciben el nombre de biomateriales. En este contexto los cementos de fosfato cálcico (CPCs) presentan unas propiedades excelentes, tanto biológicas como fisiológicas, para ser empleados en la reparación o sustitución de elementos óseos. Los CPCs son altamente bioactivos, biocompatibles y osteoconductivos, es decir, son resaborbidos por el suero fisiológico del cuerpo humano, permitiendo así su progresiva sustitución por hueso nuevo, recién formado. De forma general, los CPCs se obtienen mediante la mezcla de uno o varios ortofosfatos de calcio, que actúan como precursores, en un medio acuoso. Dependiendo de la cantidad de fase líquida empleada, se obtiene una especie de masilla, fácilmente moldeable, o un líquido inyectable. Ambos productos son capaces de fraguar in vivo. Es por esto que los CPCs se presentan como unos biomateriales altamente prometedores en su uso para aplicaciones clínicas, como puede ser el relleno de cavidades óseas. Además, en la mayoría de los CPC se obtiene hidroxiapatito deficiente en calcio como producto de la reacción de fraguado. Dicho hidroxiapatito (HA) es química y estructuralmente muy similar al apatito biológico, el principal componente inorgánico de los huesos, lo que confiere a los CPC unas características excelentes desde un punto de vista biológico y fisiológico para su uso en implantes óseos. En general, la rápida conversión de los precursores de los CPCs al producto final HA no sólo conlleva un aumento de las propiedades fisiológicas adecuadas, sino también de las propiedades mecánicas. Sin embargo, hay que tener en cuenta que durante un proceso quirúrgico el CPC debe fraguar lo suficientemente lento como para permitir la implantación quirúrgica, y lo bastante rápido como para minimizar la duración de la operación. En consecuencia, desde un punto de vista fundamental y de aplicación, es importante diseñar CPCs que presenten tiempos de fraguado entre 10 y 30 minutos, así como una rápida conversión al producto final, preferiblemente, HA. Resumen iv Debido a que los CPCs presentan una resistencia a la fractura menor que los huesos (~0.6-1.5 MPa/m2 frente a 2-12 MPa/m2), sus aplicaciones clínicas están, en la actualidad, limitadas a áreas dónde los huesos no soportan carga, por ejemplo, reparaciones craneofaciales o periodontales. Uno de los métodos más empleados para mejorar las propiedades mecánicas de los CPCs es la adición de fibras que actúen como fase de refuerzo. Por ejemplo, las cerámicas de HA han sido reforzadas con éxito mediante el empleo de fibras de SiC, carbón, Si3N4, Al2O3 o ZrO2. Sin embargo, dichos materiales son, en el mejor de los casos, bioinertes, es decir, no interaccionan con el tejido óseo circundante, y su empleo conlleva una disminución de la bioactividad de los correspondientes composites cerámicos. En contraste con estas fibras bioinertes, si se emplean partículas de HA como fase de refuerzo, éstas no sólo actúan como tal, sino que también mejoran la bioactividad de los composites finales, siendo ésta la gran ventaja de las fibras de HA empleadas como refuerzo de biomateriales cerámicos o de CPCs. Cuando se quiere reforzar un material, es de vital importancia tener en cuenta la morfología de las partículas empleadas como fase de refuerzo. Así, en composites isotrópicos, las inclusiones con forma de lámina son las más efectivas, seguidas de las fibras, siendo la morfología menos efectiva la esférica. Por tanto, es importante desarrollar una ruta sintética efectiva que nos permita obtener cristales puros de HA sin defectos y con una morfología definida y controlable. Así, con unos materiales de partida adecuados, es posible obtener nuevos biocomposites (composites empleados como biomateriales) para su empleo en sustitución y reparación ósea. Estos composites están compuestos por un CPC con una rápida conversión a HA y un tiempo de fraguado adecuado como matriz, y cristales de HA con diferentes morfologías como fase de refuerzo. Dichos materiales tienen la ventaja de que todos sus componentes son biocompatibles, bioactivos y osteoconductivos. Es decir, son aceptados por los tejidos circundantes y son capaces de establecer uniones con ellos, proporcionando un soporte en el que el nuevo hueso puede formarse. Además de la morfología, el tamaño y la fracción volumétrica ocupada por el refuerzo de HA son factores clave en las propiedades mecánicas finales de estos biocomposites, y por lo tanto, necesitan ser optimizadas. Resumen v Objetivos Tal y como se ha mencionado, los cementos de fosfato cálcico son uno de los materiales con mayor potencial para su empleo en reparación y/o sustitución ósea. Uno de los principales objetivos de este trabajo es preparar un cemento de fosfato cálcico bioactivo, que fragüe entre 10 y 30 minutos y con una rápida conversión al producto final, HA. El segundo objetivo es el desarrollo de una ruta hidrotermal fácil y económica que nos permita obtener partículas de HA para su empleo como fase de refuerzo. Variando las condiciones de síntesis, es posible sintetizar monocristales de HA sin defectos y con distintas morfologías. El CPC preparado, así como los cristales de HA sintetizados, deben de ser adecuadamente caracterizados con el objetivo de determinar su estructura, composición y propiedades físico-químicas. Una vez logrados los componentes adecuados, se estudiará el refuerzo del CPC empleando los cristales de HA como inclusiones, obteniendo así una nueva familia de biocomposites con unas propiedades mecánicas mejoradas con respecto al CPC puro. Para ello, se estudiará la influencia de la morfología, tamaño y fracción volumétrica de los cristales de HA empleados como refuerzo en las propiedades mecánicas del biocomposite final. Por último, se optimizarán estos factores con el fin de obtener nuevos y funcionales biomateriales con unas propiedades mecánicas, biológicas y fisiológicas mejoradas con respecto al CPC puro. Estrategias –. Cemento de fosfato cálcico: Se empleó fosfato dicálcio dihidratado (DCPD) y fosfato tetracálcico (TTCP) como precursores en la obtención del CPC, ya que dichos compuestos reaccionan para dar hidroxiapatito como producto final. La estrategia empleada para acelerar la reacción de fraguado del cemento y su conversión a HA fue la disminución del tamaño de partícula de los precursores mediante molienda. El empleo de una fase líquida saturada con iones fosfatos es otro de los métodos empleados para activar la reacción de los precursores de los CPCs. Sin embargo, esta técnica presenta varias limitaciones, por ejemplo, la incorporación de iones extraños en la estructura cristalina del HA final. Así mismo, existe un exceso de iones reactivos que no participan en la reacción en el medio, y los bajos o altos valores iniciales de pH que ello conlleva suelen facilitar la formación de especies intermedias. Estos efectos pueden crear problemas de Resumen vi biocompatibilidad o disminuir la velocidad de reacción del CPC. Por el contrario, el empleo de DCPD y TTCP finamente molidos asegura una rápida disolución de los precursores, lo que implica una mayor velocidad de la reacción de fraguado del CPC. –. Cristales de hidroxiapatito (HA): El método hidrotermal fue la ruta seleccionada para sintetizar cristales de HA (Ca10(PO4)6(OH)2) debido a su exitosa aplicación para la síntesis de la fase deseada bajo condiciones de reacción suaves y mediante un solo paso. Así mismo, la técnica hidrotermal permite la obtención de monocristales sin defectos, altamente cristalinos, con un tamaño de partícula relativamente controlable y con diversas características morfológicas. Así, se sintetizaron monoscristales de HA con varias morfologías mediante un tratamiento hidrotermal, empleando nitrato de calcio tetrahidratado e hidrogenofosfato de diamonio como reactivos. Adicionalmente, se utilizó urea como agente de precipitación homogénea. La urea se descompone uniformemente entre 80° y 95°C, liberando NH3(aq) al medio de reacción, lo que a su vez, aumenta el pH del sistema hasta alcanzar valores en los cuales HA no es soluble, y por lo tanto, precipita. La urea es muy soluble en agua, y la cinética de su descomposición térmica depende de la temperatura y del tiempo que se caliente la mezcla. Por lo tanto, mediante la variación juiciosa de las condiciones del tratamiento hidrotermal, fue posible controlar el pH del medio de reacción y, en consecuencia, sintetizar monocristales de HA con diversas morfologías. –. Biocomposites: Se obtuvieron nuevos biocomposites mediante la combinación del CPC obtenido (matriz) y cristales de HA con diversas morfologías (fase de refuerzo bioactiva). La idea principal es que la presencia de una inclusión mecánicamente más fuerte que la matriz puede prevenir la propagación de fisuras y mejorar las propiedades mecánicas de dichos biocomposites frente al CPC puro. Ambas situaciones son de extrema importancia en el desarrollo de nuevos biomateriales para reparación de huesos que soportan carga. La influencia de la morfología, tamaño y fracción volumétrica de las inclusiones de HA en las propiedades mecánicas finales de los correspondientes biocomposites fue estudiada y optimizada para obtener un biomaterial funcional. Resumen vii Resultados y discusión • Obtención y caracterización de los cementos de fosfato cálcico Para la obtención de los cementos de fosfato cálcico se eligió la reacción de disolución-precipitación entre fosfato dicálcico (DCPD, básico) y fosfato tetracálcico (TTCP, ácido) para preparar CPCs con HA Ca10(PO4)6(OH)2 como producto final. . El trabajar con un menor tamaño de partícula de los precursores conlleva un incremento en su solubilidad, lo que a su vez, implica un aumento de la velocidad de reacción de formación del CPC. Por ello, se realizó la síntesis de los CPCs mediante varias etapas. La etapa inicial supuso la síntesis de DCPD y TTCP puros (muestras D-raw y T-raw respectivamente). A continuación, estos precursores fueron molidos mediante un molino de bolas con el fin de disminuir su tamaño de partícula. Se emplearon tres tamaños de bolas distintos con el fin de obtener distintos tamaños de partícula. Mediante combinación de los distintos precursores obtenidos tras la molienda, se obtuvieron nueve CPC, cuyas propiedades y características fueron estudiadas y comparadas con el fin de seleccionar el CPC óptimo con el que preparar los biocomposites. El producto D-raw se obtuvo mediante una reacción de precipitación, mientras que T-raw fue sintetizado mediante una reacción a alta temperatura (1500°C). El análisis de las muestras D-raw y T-raw mediante microscopía electrónica de barrido (SEM) y dispersión de luz dinámica (DLS) muestra que ambos están compuestos por cristales y agregados de cristales de varias micras (~6 y ~7 μm respectivamente), lo que probablemente implique una velocidad de disolución baja en el caso de ambas muestras. Con la idea de mejorar la reactividad y solubilidad de los precursores, D-raw y T-raw fueron molidos empleando tres tamaños de bolas distintos (3.6, 4.8 y 9.6 mm). La pureza de las muestras tras la molienda fue comprobada mediante difracción de rayos X (XRD). El estudio de las muestras obtenidas mediante SEM y DLS confirma que la molienda de D-raw y T-raw reduce de modo efectivo el tamaño de partícula y los agregados de los compuestos de partida. La distribución del tamaño de partícula indica que el tamaño de las bolas empleadas en la molienda influye en el tamaño de partícula obtenido: las bolas más pequeñas proporcionan muestras más finas y con una menor distribución de tamaño de partícula (PSD). Así, DCPD y TTCP se obtienen con un tamaño medio de partícula de 3.66 y 2.96 μm, 1.69 y 1.52 μm o 1.13 y 1.33 μm según las bolas empleadas sean de 9.6, 4.8 o 3.6 mm respectivamente. Resumen viii Los tres DCPD y los tres TTCP obtenidos tras la molienda fueron combinados entre si y se obtuvieron nueve CPCs. En contraste con los datos publicados hasta la fecha para un CPC formado por DCPD y TTCP, el análisis mediante XRD muestra que, tras 24 horas de reacción, todos los CPC obtenidos consisten únicamente en HA puro. Muchos grupos han investigado la preparación de CPC empleando DCPD y TTCP como precursores, sin embargo, en la mayoría de los casos, tras 24 horas, la mezcla de reacción todavía contiene TTCP sin reaccionar, lo cual puede generar problemas de biocompatibilidad, ya que el TTCP se hidroliza en agua, generando iones OH– que hacen aumentar el pH del medio a valores muy básicos, tóxicos para las células óseas: −−+ ++⎯⎯→⎯+ )( )( 3 4 2 )()(2)(244 224)( 2 aq aq aq OH aqs OHPOCaOHOPOCa Se cree que el TTCP remanente está relacionado con un mayor tamaño de partícula de este compuesto comparado con DCPD (~10 veces mayor) y con la formación de una capa de HA en la superficie de los reactantes que actúa como barrera y dificulta su disolución. Hasta la fecha, solo fue posible obtener un CPC basado en DCPD y TTCP que se convierte totalmente en HA tras 24 horas empleando una fase liquida saturada con iones fosfato. Sin embargo, dicho CPC presenta un alto valor de pH durante aproximadamente siete días, lo que provocaría daños celulares en el tejido circundante al emplearlo in vivo. La determinación del tiempo de fraguado y los estudios de porcentaje de conversión a HA realizados para los nueve CPC preparados, indican que es el cemento C-D36/T48 formado por DCPD y TTCP con un tamaño medio de partícula de 1.13 y 1.52 μm respectivamente el que presenta un menor tiempo de fraguado (~22 minutos) y una mayor conversión a HA tras tres horas de reacción (~33%). Es importante destacar que este tiempo de fraguado cumple los requisitos médicos para CPCs y que se obtiene usando sólo agua como fase liquida, es decir, sin emplear aceleradores de reacción que pueden presentar potenciales problemas de biocompatibilidad. Los resultados obtenidos mediante el estudio de la evolución de la reacción empleando XRD, indican que C-D36/T48 se transforma totalmente a HA tras sólo seis horas de reacción, lo cual es una clara ventaja comparado con cementos similares que tras 24 horas todavía no han reaccionado totalmente. Cuanto más rápido se produzca la reacción de fraguado del cemento, antes alcanzará éste sus propiedades mecánicas finales. El estudio de la evolución de la reacción de fraguado también muestra que ésta sigue un proceso de nucleación-crecimiento de cristales. Resumen ix El análisis mediante espectroscopía de dispersión de rayos X (EDX) confirma que el cemento C-D36/T48 es deficiente en calcio, ya que obtenemos una relación Ca/P ~1.49, cuando el valor del HA estequiométrico es 1.67. El análisis mediante espectroscopía de infrarrojo (IR), Raman y análisis elemental muestra la existencia de aniones HPO42− y CO32– sustituyendo parcialmente grupos PO43− en la estructura cristalina del HA. La espectroscopia de IR muestra también la incorporación de moléculas de agua en la estructura del cemento CD36/T48. Según estos resultados, la fórmula del cemento puede representarse como Ca10-x-y(HPO4)x(CO3)y(PO4)6-x-y(OH)2-x-y(H2O)z. Es importante señalar que el HA biológico presente en los huesos y dientes de los vertebrados, es deficiente en calcio y además, contiene ~3-5 % en peso de grupos carbonato. Por lo tanto, la obtención de un CPC como el C-D36/T48 con grupos HPO42− (presentes en la HA deficiente en calcio) y CO32– en su estructura no representa ningún problema desde un punto de vista biocompatible. Por otro lado, la determinación del área de superficie y de la porosidad indica que el C-D36/T48 posee un alta área superficial (~169 m2/g) y una porosidad de aproximadamente el 35% producida por mesoporos desordenados. Estudios detallados mediante microscopía de transmisión electrónica (TEM) del cemento C-D36/T48 tras 24 horas de reacción señalan que la muestra esta compuesta únicamente por HA, sin presencia de otras fases como impurezas. TEM confirma además que el CPC está formado por nanoláminas (10-15 nm) fuertemente entrelazadas entre si. Dicho entrelazamiento es el responsable de la solidez del CPC tras la reacción de fraguado. Estos nanocristales crecen según la dirección del eje c. Estudios in vitro muestran la potencial capacidad de C-D36/T48 para formar un enlace químico con el tejido óseo circundante. Es decir, el test in vitro confirma la bioactividad de C-D36/T48. La resistencia a la compresión de este cemento fue establecida como 25 ± 3 MPa, un valor que entra en el rango de valores previamente establecidos para cementos basados en DCPD y TTCP (20-170 MPa). El módulo elástico (relacionado con las propiedades elásticas) y la dureza (relacionada con las propiedades plásticas) fueron determinados mediante nanoindentación. Los valores hallados (23.1 ± 2.6 y 0.73 ± 0.2 GPa respectivamente) son ligeramente superiores a los calculados para el hueso humano trabecular (15.0-19.4 GPa) y están en el rango de los hallados para el hueso cortical (20.0-25.8 GPa). La dureza de Resumen x C-D36/T48 sigue la misma tendencia, y es ligeramente superior a la del hueso trabecular (0.52-0.62 GPa) y está dentro de los valores del hueso cortical (0.62-0.74 GPa). Estos resultados muestran una gran similitud mecánica (elástica y plástica) entre C-D36/T48 y los huesos, presentando el cemento preparado una gran compatibilidad biomecánica con los tejidos óseos. Dicha compatibilidad es de gran importancia en los biomateriales. • Síntesis de monocristrales de hidroxiapatito con distintas morfologías La técnica hidrotermal es ampliamente empleada hoy en día en la síntesis de cristales de HA. En la mayor parte de los casos, los cristales obtenidos presentan una morfología fibrosa (whiskers, fibras, agujas, etc.) o de lámina. Adicionalmente, también existen trabajos en los que las partículas sintetizadas tienen forma de prisma hexagonal, elíptica o de lazo. A pesar de estas diferencias morfológicas, los cristales obtenidos crecen según el eje c, la tendencia natural del HA. En este trabajo, los cristales de HA fueron sintetizados empleando Ca(NO3)2·4H2O y (NH4)2HPO4 como reactivos y urea como agente de precipitación. La urea se descompone entre 80° y 95°C liberando NH3 al medio de reacción, elevando el pH del medio hasta valores a los cuales el HA es la fase menos soluble. La descomposición de la urea es dependiente de la temperatura, por lo cual es posible controlar su velocidad de descomposición aplicando distintos ciclos de temperatura durante el tratamiento hidrotermal (Tmáx. = 90°C). Todos los productos obtenidos son HA puro (XRD y Raman), ligeramente no estequiométrico (EDX, Ca/P ~1.75). Los análisis mediante IR muestra la parcial sustitución de iones OH− y PO43− por grupos CO32− (~ 0.75% en peso según análisis elemental) en la estructura cristalina de los productos. La existencia de grupos HPO42−, indicativos de una deficiencia en calcio del HA obtenido, no fue observada mediante IR, Raman o análisis termogravimétrico y diferencial (TGA/DTA), en concordancia con los resultados obtenidos mediante EDX. Los estudios realizados mediante SEM muestran que las partículas sintetizadas presentan distinta morfología en función del ciclo de temperatura empleado durante el tratamiento hidrotermal y de la concentración de reactivos empleada. Así, si la síntesis se realiza a una temperatura constante de 90°C, la urea se descompone de modo intensivo, proporcionando una importante supersaturación del medio de reacción y, consecuentemente, deberían formarse un gran número de núcleos de cristalización. Sin embargo, en nuestro caso las partículas obtenidas muestran forma de lámina (muestra PT), probablemente debido a un crecimiento epitaxial de los cristales de HA sobre cristales de Resumen xi fosfato octacálcico Ca8H2(PO4)6·5H2O, precursor de los cristales de HA. Por tanto, se supone que la morfología de lámina de la muestra PT se debe a un predominio de efectos cinéticos sobre los efectos termodinámicos de la reacción. Si el ciclo de temperatura empleado no se mantiene constante a 90°C sino que presenta varias etapas en las que se disminuye la temperatura a 70°C (por debajo del rango de descomposición de la urea), los cristales sintetizados presentan morfología de prisma hexagonal (muestra HX). Este resultado indica que con estas condiciones de reacción, los efectos termodinámicos son predominantes. Si la descomposición de la urea se realiza lentamente, las partículas sintetizadas presentan forma de aguja (muestra ND). Dicha morfología concuerda con un predominio de los efectos termodinámicos sobre los cinéticos, ya que se forman un gran número de núcleos de cristalización que al crecer, producen cristales de gran tamaño y pequeña sección. Con un ciclo de temperatura que implica no solo la interrupción de la descomposición de la urea sino que además mantiene dicha interrupción durante una hora, es imposible obtener HA puro si el resto de las condiciones de reacción se mantienen constantes. Sin embargo, si disminuimos la concentración de calcio y fósforo de 0.167 y 0.1 M respectivamente (muestras PT, HX y ND) a 0.083 y 0.05 M los cristales obtenidos presentan, mayoritariamente, forma de laminas finas (muestra FP). También se observa la presencia de partículas casi rectangulares. En general, el tamaño de partícula de esta muestra es mucho menor que el de los productos anteriores. Empleando esta misma concentración de calcio y fósforo (0.083 y 0.05 M respectivamente) pero utilizando cualquiera de los otros ciclos de temperatura mencionados, se obtienen partículas con el mismo tipo de morfología. Estos resultados indican que, para esta concentración, lo que predomina es el efecto de la concentración de reactivos. Por lo tanto, en sistemas diluídos, el control morfológico de los cristales mediante la descomposición controlada de urea está de algún modo limitado. Los estudios de estas muestras mediante TEM confirman que todos las partículas cristalizan en el sistema hexagonal del HA y que son monocristalinas. Además, dichos estudios muestran que, mientras las muestras PT, HX y ND crecen a lo largo del eje c, los cristales de FP crecen según las direcciones (211) en el caso de las partículas con morfología de lámina fina y (102) para las partículas rectangulares. El HA tiene una tendencia natural a crecer orientado según el eje c, por lo que estos resultados son altamente interesantes, pero, al mismo tiempo, sorprendentes. El módulo elástico y dureza de los cristales con morfología hexagonal HX, se determinó mediante nano-indentación. Los resultados obtenidos muestran una anisotropía en Research motivation and background 4 rational synthetic approach to calcium phosphate bone-cement biomaterials with setting time between 10-30 min and a rapid conversion rate to the end product, preferably HA. Owing to the lower fracture toughness parameters of CPCs in comparison to human bone (~0.6-1.5 cf. 2-12 MPa/m2), the clinical applications of CPCs are nowadays limited to areas where bones are free of dynamic load, i.e. for craniofacial and periodontal applications, or as materials for the development of scaffolds in bone tissue engineering. One of the most promising methods to improve the mechanical performance of biomaterials is their reinforcement via physical linking by spatially allocated fibrous materials. For example, HA ceramics can be significantly reinforced by various bioinert fibres like SiC, C, Si3N4, Al2O3 or ZrO2 (bioinert reinforcement). However, utilization of these fibrous materials leads to a decrease in the bioactivity of the respective composite bioceramics. Hence, research related to bioactive reinforcement, wherein hydroxyapatite particles are applied as a reinforcement phase, has recently received much attention. Particles with plate-like morphologies are the most effective at stiffening isotropic composite materials, followed by fibrous shapes, being the least effective geometry the spherical particles. In this way, HA crystals with desired morphological features are applicable for moderated reinforcement of biomaterials for bone repair and substitution, such as the calcium phosphate bone-cements previously mentioned. Consequently, it is of critical importance for reinforcement applications to develop an effective synthetic route to the phase-pure hydroxyapatite single crystals with controlled morphological features. Finally, having a complete set of constituent materials –advanced calcium phosphate bone-cement (matrix) and hydroxyapatite crystals with various particle morphologies (fillers)– it is certainly possible to explore a new family of reinforced biocomposite cements. These materials have the advantage that all their components are biocompatible, bioactive and osteoconductive. Another appealing feature of such composite biomaterials is that the filler is mechanically stronger than the matrix, and therefore, being thoroughly embedded in cement, can halt potential crack propagation when applied as bone graft. The shape, size and volume fraction of the hydroxyapatite filler are key parameters that influence the final mechanical performance of such biocomposites, and need to be optimized. Such optimizations, together with an analysis of the correlation between the material properties of the filler and the mechanical performance of the respective reinforced biocomposite cement, can allow the development of new insights into the design of composite biomaterials and allow for new biomedical applications reinforced biocomposite cements. Research motivation and background 5 1.2.- BACKGROUND The Instituto de Cerámica de Galicia (ICG, Galician Institute of Ceramics) is a relatively young institute of the University of Santiago de Compostela in Spain, which is well known worldwide for its traditionally strong education in the fields of natural sciences. Despite its young “age” ICG is one of the leading institutions involved in ceramics research in Spain. The investigation into original research topics at the institute is realised in close cooperation between the Chemistry, Physics, Materials Chemistry, and Biology departments; a special characteristic of the institute. A remarkable proportion of the research projects at the ICG are devoted to the field of the biomaterials. Since the first study in this research area in 1986, members of the institute have published a series of papers on biomaterials, which have been highly acclaimed within the scientific community. Accordingly, several research projects have been successful carried out on various biomaterials. Specifically, in 1998, the project entitled “Biocompatibility and design of ceramic materials for their application in clinical orthopaedics” (1988-89), was commenced in collaboration with J. Couceiro, M.D., Ph.D, Head of the Orthopaedic Surgery Department of the General Hospital of Galicia and the company Cerámica de Sargadelos. From then on the ICG has been continuously involved in the design and characterization of calcium phosphate ceramics, mostly hydroxyapatite and tricalcium phosphates in projects such as “Biocompatible materials of hydroxyapatite for their application in clinical orthopaedic” (1990-91), where the syntheses of hydroxyapatite and tricalcium phosphate of high purity were investigated and optimized to obtain dense blocks of different sizes. As a result of these investigations a new project was started, “Design and construction of a pilot plant for the industrial manufacture of biocompatible hydroxyapatite” (1989-91). This research led to the fabrication of these materials at a pilot scale in the factory Cerámica de Sargadelos in Cervo, Lugo (Spain). In the following years (1991-98), the research carried out by the biomaterials group in the ICG widened to include study of “Bioactive, structural materials of wollastonite-zircon” (1991-93), where bioactivity of a phosphate-free polycrystalline ceramic material, wollastonite, was established for the first time. This was followed up with another project: “Design of wollastonite biomaterials and wollastonite-containing systems” (1995-98). A new material with eutectic structure was developed within the wollastonite-tricalcium phosphate system that transforms in situ into a porous HA structure. This new material was named Bioeutectic® and was registered as an international trade mark. These results led to a new Research motivation and background 6 project titled “Design and development of Bioeutectic materials. In vitro and in vivo studies” (1997-2000)”. Simultaneously, and continuing with the study of polycrystalline bioceramics, the project “Study, design and fabrication of tricalcium phosphate implants for neurology, dentistry and orthopaedic surgery” (1999-2001) was carried out, where a procedure for the fabrication of tricalcium phosphate implants for those applications was optimized. In 1999 researchers from the ICG with a group of industrial sponsors founded KERAMAT, Ltd., a company dedicated to the manufacture and marketing of bioceramics of calcium phosphates. This company is part of UNINOVA, a spin-off created by the University of Santiago de Compostela and the City Council to launch innovative academic research projects onto an industrial level. Upon its creation, KERAMAT Ltd. signed with the University of Santiago de Compostela, and in particular, with the Instituto de Cerámica de Galicia, a series of agreements for the transfer of some of the technologies developed in the ICG. Due to the emergent interest and as a logical consequence to the research in calcium phosphate compounds started in 1986, the ICG initiated in 2000, research on calcium phosphate cements which was soon followed with the project “Design, obtaining and properties of bioceramic mortars for its use in orthopaedic surgery, neurosurgery and maxillofacial surgery” (2003-2007, Spanish Ministry of Education and Science, project MAT2002-03857), in which this dissertation is included. An important part of the research reflected in this thesis was carried out at the Tokyo Institute of Technology (Materials and Structures Department, Tokyo-Japan) and the Max Planck Institute of Colloids and Interfaces (Biomaterials Department, Potsdam-Germany) which started an ongoing collaboration with the Instituto de Cerámica de Galicia. II INTRODUCTION 2.1.- BONES .........................................................................................................................9 2.1.1. Bone constituents and structural hierarchy .........................................................10 2.1.2. Microarchitecture ................................................................................................ 14 2.1.3. Bone shape and function ..................................................................................... 17 2.1.4. Bone remodelling................................................................................................ 19 2.1.5. Bone microdamage. Microcracks........................................................................20 2.1.6. Bone fracture....................................................................................................... 20 2.2.- TEETH........................................................................................................................ 21 2.2.1. Structure .............................................................................................................. 21 2.2.2. Mechanical properties .........................................................................................24 2.2.3. Hydroxyapatite in teeth....................................................................................... 25 2.3.- BONE REPAIR AND SUBSTITUTION...................................................................26 2.3.1. Bone defects........................................................................................................26 2.3.2. Natural bone grafts. Autogeneous, allogeneous and xenogeneous implants ......26 2.3.3. Synthetic bone grafts: Biomaterials ....................................................................28 2.3.4. Metallic biomaterials...........................................................................................32 2.3.5. Polymeric biomaterials........................................................................................ 34 2.3.6. Ceramic biomaterials........................................................................................... 36 2.4.- CALCIUM ORTHOPHOSPHATES..........................................................................42 2.4.1. Solubility phase diagram.....................................................................................44 2.4.2. Biological occurrence and medical uses .............................................................46 2.5.- CALCIUM PHOSPHATE BONE-CEMENTS..........................................................49 2.5.1. Types of calcium phosphate cements.................................................................. 52 2.5.2. Control of the setting reaction. Cohesion and setting time .................................54 2.5.3. Future developments of calcium phosphate bone-cements................................. 59 2.6.- REINFORCEMENT OF BIOCERAMICS AND CPCs.............................................60 2.6.1. What is a composite?. Biocomposites................................................................. 61 2.6.2. Fibrous reinforcement concept............................................................................ 61 2.6.3. Fibrous reinforcement of calcium orthophosphates............................................63 2.6.4. Bioinert reinforcements....................................................................................... 64 2.6.5. Bioactive reinforcements..................................................................................... 65 2.7.- HYDROXYAPATITE CRYSTALS .......................................................................... 66 2.7.1. Apatite 66 2.7.2. Hydroxyapatite.................................................................................................... 68 2.7.3. Hydrothermal approaches to hydroxyapatite crystals......................................... 74 2.8.- OBJECTIVES AND STRATEGIES .......................................................................... 78 2.8.1. Research objectives............................................................................... 78 2.8.2. Research strategies................................................................................79 Introduction: Bones 9 2.1.- BONES1-5 Bone is one of the mineralized tissues found in vertebrates, together with cartilage, dentine, enamel and cementum. Normally, we think of cartilage and bone as skeletal tissues, while enamel, dentine and cementum are considered as dental tissues, since they form part of the endoskeleton or the teeth of vertebrates respectively. The function of bone is to move, to support, and to protect the various organs of the body; to produce red and white blood cells;i and to store minerals –approximately 99% of calcium, ~85% of phosphorous and between 40 to 60% of sodium and magnesium in the human body are stored in bones. Bone is a composite material consisting of an organic matrix (collagen, ~1/3 of the dry weight) in which are embedded mineral crystals of hydroxyapatite (HA, Ca10(PO4)6(OH)2), a calcium orthophosphate. The volume fraction distribution between organic and mineral phase is ~60/40. Its mechanical, chemical and biological properties as composite are considerably different from the characteristics of either component separately. Bone combines the optimal properties of both, stiffnessii and toughness.iii While HA bone mineral is stiff and brittle,iv the protein (collagen) is much softer and tougher, and hence, also more fracture resistant. Additionally, water, bone cells, polysaccharides, proteins and blood vessels are also present in the bone. Water corresponds to 15-25 vol.-% of the bone in mammals.6 Bone is considered as a living material. Throughout one’s lifetime, bone undergoes constant remodelling to repair microfractures and replace old tissue. Over the course of a ten year period, the entire skeleton is completely reformed. Damaged or old bone is resorbed by osteoclastsv and new bone is formed in the resulting cavity by osteoblasts.vi As long as the rates of bone resorption and bone formation are balanced, there is no net loss or gain in the amount of bone, and its structural and material properties are maintained. The properties of the Eiffel tower (constructed in 1889) hold many parallels to those of bone. The structure was designed to withstand wind force and has an open-lattice design (Fig.2.1). The tower was built of iron to be flexible but strong and to resist collapsing. Due to i Red blood cells are the most common type of blood cell and the vertebrate body's principal means of delivering oxygen from the lungs or gills to body tissues via the blood. White blood cells, or leukocytes, are cells of the immune system defending the body against both infectious disease and foreign materials. ii Resistance of a body to elastic deformation by the action of a force. iii Fracture resistance of a material when a force is applied. iv A material is brittle if it is liable to fracture when subjected to a force. Opposite to toughness. v An osteoclast (from the Greek words for "bone" and "broken") is a type of bone cell, located on the bone surface, that removes bone tissue by removing its mineralized matrix. This process is known as bone resorption. vi An osteoblast (from the Greek "bone" and "germ" or embryonic) is a bone cell, present on the bone surface, responsible for bone formation. Introduction: Bones 10 the continued repair of weakened material, the Eiffel tower still stands strong over more than a century after its construction. Fig. 2.1. View of the Eiffel tower (Paris, France) Nature achieves this task by designing bones that are adapted to the different functions for different parts of the skeleton. According to the laws of physics, the strength of a material is determined by its material and structural properties and is relative to the specific direction of forces applied to it. This means that the strength of a bone is not just determined by its mass, but depends on the materials of which it is made and how they are built up into the overall structure of the bone. In order to serve the differing biomechanical functions of the skeleton, bone must have the following qualities: K light (for speed of movement) K strong (for load-bearing) K stiff (to resist bending under load; to move the body against gravity) K flexible (to absorb the energy of impacts without fracturing) 2.1.1. Bone constituents and structural hierarchy 2.1.1.1. Bone structure As discussed above, bone properties are a list of apparent contradictions, strong but not brittle, rigid but flexible, light-weight but solid enough to support tissues, mechanically strong but porous, stable but capable of remodelling, etc. Introduction: Bones 11 Table 2.1. Mechanical properties of human bones. Property Values Elastic modulusi (GPa)1 12-20 Hardnessii (GPa)7 0.52-0.62 Compressive strengthiii (MPa)120-205 Tensile strengthiv (MPa)1 53-133 Fracture toughnessv (MPa/m2)82-12 Bone is able to fulfil these opposing requirements by a hierarchical structure going from molecular to macroscopic scale (Fig 2.2). Therefore, its mechanical properties are determined by its structure. Fig. 2.2. Bone hierarchical structural organization as follows: (a) cortical and cancellous bone, (b) osteons with Haversian systems, (c) lamella, (d) collagen fibres assemblies of collagen fibrils, (e) hydroxyapatite nanocrystals, collagen molecules and non-collagen proteins (after Rho et al.).2 i It is the mathematical description of an object or substance's tendency to be deformed elastically under the action of a load. It is a measure of the stiffness of a material. ii It refers to various properties of matter in the solid phase that gives it high resistance to permanent shape changes when a force is applied. It is related with plastic deformations. In bones, hardness reflects the degree and quality of mineralization of the bone. iii It is the capacity of a material to bear axially directed pushing forces. When the limit of compressive strength is reached, the material fractures. iv It is the resistance of a material to a force tending to tear it apart, measured as the maximum tension the material can withstand without tearing. v It is the ability of a material containing a crack to resist fracture. Introduction: Bones 12 The building block of the organic component is the collagen, which is a triple helix with a diameter of ~1.2 nm. These tropocollagens (or collagen molecules) present a periodicity of 67 nm –distance by which adjacent tropocollagens are staggered– and holes (or gaps) of 40 nm between their ends which become filled up with HA nanocrystals (Fig. 2.3A). Therefore, collagen molecules are intercalated with the mineral phase HA forming fibrils that are stabilized by cross-linking (Fig. 2.2, 2.3A). The collagen fibrils are arranged in concentric layers with the biological HA embedded into them (Fig. 2.3A, inset). Many collagen fibrils together make up a collagen fibre, also stabilized by cross-linking. The collagen fibres are in turn organized into sheets, either stacked in parallel arrays (lamella) or wrap in concentric layers (3-8 lamellae) around a central canal to form what is known as an osteon or a Haversian system (Fig. 2.2). The osteon looks like a cylinder ~200-250 μm in diameter running roughly parallel to the long axis of the bone. Each layer of the osteon has its constituent fibrils oriented in alternate directions similar to plywood (Fig. 2.2). On a longer scale, the osteons are grouped together into long bundles that are the basic building block of cortical bone (Fig. 2.2, 2.3B, 2.3C). Finally, at the macroscopic level, each bone has a specific shape and structure so that the skeleton can work as an integrated whole (Fig. 2.2). Fig. 2.3. (A) Schematic diagram showing the assembly of collagen fibrils and fibres and hydroxyapatite nanocrystals. The 67 nm periodic pattern results from the presence of adjacent holes (40 nm) and overlaps (27 nm) regions of the assembled molecules (after Rho et al.).2 SEM (B) and polarized light microscope (C) images of a cortical bone showing the grouping of several osteons. The concentric lamellaecan be seen in (C). Introduction: Bones 13 2.1.1.2. Bone composition: collagen and hydroxyapatite. Woven and lamellar bone. The characteristics of HA crystals and of collagen (composition and cross-linking) and the way both component interact is important for bone strength. E Collagen in bone: About 90% of the collagen present in bone is type I. Bone derives its toughness from collagen, which can stretch like an elastic band and allows bone to deform to a certain extent without fracturing. Bone strength is influenced by the orientation and amount of collagen present and the type and number of cross-links. The degree of longitudinal orientation of collagen fibres is a strong determinant of bending strength and stiffness. Disorientation and random distribution of collagen fibres can be detrimental to bone quality. Therefore, above the level of the collagen fibrils and its associated mineral content, mammalian bone exhibits two distinctive structural forms: woven bone and lamellar bone. Woven bone is weaker with respect to lamellar bone but is formed rapidly. It is usually laid down very quickly (> 4 µm per day), most characteristically in the foetus and in the callus that is produced during fracture repair. The collagen is variable and oriented almost randomly. Woven bone is replaced by lamellar bone, highly organized, filled with many collagen fibres parallel to other fibres in the same layer. Lamellar bone laid down much more slowly than woven bone (< 1 µm per day). The fibres run in opposite directions in alternating layers to form a ply-like structure (Fig. 2.2), assisting in the bone's ability to resist torsion forces. E Hydroxyapatite characteristics: Controlled HA nucleation and growth occurs within the gaps formed by the collagen matrix (Fig. 2.3A). Therefore, the collagen matrix works like a net defining the space of the mineral crystals. The gaps localize a microenvironment containing free mineral ions and bound side chain groups, with a molecular periodicity that serves to nucleate the mineral phase heterogeneously. The nucleation of the HA nanocrystals within the gaps between the collagen fibrils (Fig. 2.3A) forces the crystals to be discrete and discontinues. HA crystals grow with a specific crystalline orientation – the c-axes of the crystals are mostly arranged parallel to each other and to the long axes of the collagen fibrils. On a larger scale, the crystals are aligned in parallel arrays across individual fibrils and sometimes this arrangement is even coherent across several adjacent fibrils. Such a long-range ordering of the collagen matrix and HA crystallites could be responsible for the unusual fracture properties of bone. Introduction: Bones 20 at the same time in their bone.10Secondary mineralization includes an increase in the size of mineral crystals (Fig. 2.4) and changes to their composition and crystal structure. An imbalance in remodelling in favour of resorption can occur with age, menopause and disease (i.e., as in osteoporosisi) Remodelling allows bone to fulfil its structural and metabolic functions, since it permits the repair of damage and the maintenance of the structural integrity of the skeleton (targeted remodelling); and the storage and release of calcium, phosphorous, etc., to maintain the acid-base balance (non-targeted remodelling). It has been estimated that ~30% of remodelling is targeted and that ~70% is non-targeted. 2.1.5. Bone microdamage. Microcracks. Axial loads able to deform one of the bone dimensions ~1-1.5% causes microdamage in the form of microcracks, which are a means of absorbing energy from the load without causing complete fracture of the bone. The presence of microcracks stimulated remodelling, and under normal conditions, they are repaired as fast as they are produced. If, however, microcracks form faster than they are repaired, bone strength may be adversely affected, and so, the rate at which microcracks are repaired influences strength. Microcracks accumulate with age in both cortical and trabecular bone, which may be related to the increased fracture risk in elderly persons. 2.1.6. Bone fracture From a mechanical perspective, fractures represent a structural failure of the bone when the forces applied to the bone exceed its load-bearing capacity or the whole-bone strength. Bone adapts its size and shape to accommodate the loads to which it is exposed by bone modelling and remodelling (Sections 2.1.3 and 2.1.4), but too high forces or in unusual directions, might result in bone fracture. The ability of a bone to resist fracture depends on the amount of bone (i.e., mass), the spatial distribution of the bone mass (shape and microarchitecture) and the intrinsic properties of the materials that comprise the bone (collagen, HA). Therefore, although many fractures are the result of high force impact or stress, bone fracture can also occur as a result of certain medical conditions that weaken the i In osteoporosis the bone mineral density (a measure of the amount of minerals, mainly calcium, contained in a certain volume of bone) is reduced, bone microarchitecture is disrupted, and the amount and variety of noncollagenous proteins in bone is altered. It makes bones weak and more likely to break. Introduction: Teeth 21 bones, such as osteoporosis, certain types of cancer or osteogenesis imperfectai or to age due to microcracks accumulation Also, the emergence of bone fragility with age is the result of the failure of modelling and remodelling adaptive mechanisms, with remodelling imbalance in each basic structural unit producing a decline in the volume of bone formed and an increase in the volume of bone resorbed. Therefore, the rate of remodelling affects the degree of mineralization of bone and microdamage repair. Imbalances in bone resorption and formation, mainly due to the detriment of formation, lead to a reduction in the amount of bone formed and to net bone loss. This has an adverse effect on trabecular and cortical microarchitecture and on bone strength. The size of the mineral crystals themselves also affects bone strength. The bones of older animals or form osteoporic animal models contain a greater portion of large HA crystals. These bones tend to fracture more easily than bones from younger animals that contain a greater mixture of smaller and larger crystals. 2.2.- TEETH6, 9, 11 Teeth consist of a unique composition of the three dental mineralized tissues, namely dentine, enamel and cementum. The three form a complex 3D structure designed to ensure that the tooth as a system can function for a time span of about seven decades, since, in contrast with bone which can undergo self-repair and remodelling, enamel, dentine and cementum are static and cannot be repair by remodelling. 2.2.1. Structure Innervated soft tissue (dental pulp or tooth’s nerve) is totally covered by dentine, which, in turn, forms the core of a tooth. In the upper area of a tooth (crown) the dentine is covered by enamel, harder and more mineralized than dentine. In the root area the outer layer of the dentine is coved by cementum (Fig. 2.10), which anchors the tooth in its socket by inserting Sharpey's fibres. ii This combination makes possible that teeth can withstand mechanical forces such as chewing forces up to 800 N, wear (abrasion), chemical or physical interactions such as changes in the pH value or variations in temperature, and biological effects such as bacteria and toxins. i People with this genetic bone disorder either have less collagen than normal or the quality is poorer than normal. ii There are matrixes of connective tissue consisting of bundles of strong collagenous fibres. In the teeth, Sharpey's fibres (also called perforating fibres) are the ends of principal fibres that insert into the cementum. Introduction: Teeth 22 Fig. 2.10. A schematic picture of a tooth (after Dorozhkin).12 2.2.1.1. Enamel Enamel purpose is to protect the tooth from the dangers posed to the teeth by the oral environment. In general, it is vulnerable only to acid attack from excess sugar (dental caries), generalized trauma such as a blow from a hard object, and serious bruxing (the habit of clenching and grinding the teeth) with associated attrition. It is white, but somewhat translucent and allows the colour of the underlying yellow dentine to shine through to a certain extent. Unlike bone, enamel does not contain collagen. Ninety-six percent of enamel consists of mineral, the mineral being hydroxyapatite rods. Water and organic material compose the rest. Enamel varies in thickness over the surface of the tooth and is often thickest at the cusp, up to 2.5 mm, and thinnest at its border. An enamel rod is the basic unit of enamel. Measuring ∼5 μm wide to ∼8 μm high, an enamel rod is a tightly packed, highly organized mass of hydroxyapatite crystals. Enamel rods are found in row along the tooth (Fig. 2.11A, 2.11B). Within each row, the long axis of the enamel rod generally is perpendicular to the underlying dentine. The arrangement of HA crystals within each enamel rod is highly complex. For the most part, the crystals are oriented parallel to the long axis of the rod. The further away the crystals are from the central axis, the more their own orientation diverges. Each enamel rod is attached to the dentine underneath it. For this reason, cracks in the enamel penetrate only as far as the dentine. This method of attachment makes it impossible for the enamel to separate from the tooth. Introduction: Teeth 23 2.2.1.2. Dentine Dentine –which is necessary for the support of the enamel– structure is fairly similar with that of bone but simpler. It contains 30 vol.-% of collagen (90% of which is type I) and 25 vol.-% water, the remainder being HA. One of the major features of dentine is the tubules (microchannels) (Fig. 2.11C, 2.11D). Tubules are originate from the inner surface of the dental pulp and they run approximately parallel to each other and perpendicular from their point of origin throughout the structure of the dentine. They terminate at the under-surface of the enamel. The diameter and density of the tubules are greatest near the pulp. As a result, dentine is a porous material (Fig. 2.11C, 2.11D). Dentine tubules contain cellular structures and a fluid that carries nourishment to the inside. They are surrounded by HA crystals (~0.5-1 μm diameter) arranged in a random fashion. These tubular units are, on their turn, embedded in a composite consisting of a collagen matrix reinforced with HA. This is called the intertubular region. 2.2.1.3. Cementum Cementum is a specialized calcified substance covering the anatomic root of the tooth. The principal role of cementum within the tooth is to serve as a medium by which the periodontal ligaments can attach to the tooth for stability. Hence, its bottom surface is tangent to the periodontal ligaments running through the jaw (via collagen fibres), and the upper portion of the surface is firmly cemented to the dentin of the tooth. From a chemical point of view, cementum is similar to dentine and bone –it is formed by approximately 45 vol.-% of inorganic material (mainly HA), 33 vol.-% of collagen (mostly type-I) and 22 vol.-% of water. Its coloration is yellowish and, since is less mineralized, softer than either dentin or enamel. Introduction: Teeth 24 Fig. 2.11. SEM images of the column-like structure of the enamel (A), detailed of the enamel structure showing ordered rods of HA (after Dorozhkin) (B),12 general view of the dentine (C) and detail of the dentine tubules (D). 2.2.2. Mechanical properties The toughness of enamel is lower than that of the dentine, hence, dentine possess a higher resistance to fracture under the action of a force than enamel. The converse is the case for the hardness (Table 2.3). Enamel hardness is indeed significant: it is the hardest material in vertebrates (Table 2.3). The hardness of the enamel layer is due to its high degree of mineralization, however, the large amount of minerals in enamel accounts not only for its strength but also for its brittleness. Dentine, less mineralized than enamel, is also less brittle. Table 2.3. Mechanical properties of teeth.6 Property Enamel Dentine Fracture toughness (MPa/m2) 0.7-1.3 1-2 Elastic modulus (GPa) 60 Perpendicular to tubules: 5-6 Parallel to tubules: 13-17 Hardness (GPa) 4 0.5 The presence of tubules in the dentine confers a considerable degree of anisotropy to its structure, and, therefore, its mechanical properties in general, and toughness in particular, are quite anisotropic (Table 2.3). Comparing with bone (Table 2.1), teeth are harder and present a higher elastic modulus, thus, they support better elastic deformations. Teeth’s fracture toughness is lower than the found for bones. Introduction: Teeth 25 2.2.3. Hydroxyapatite in teeth Like in bones (Section 2.1.1.2), HA present in teeth is not pure, and it contains several ions in the crystal structure partially substituting calcium, phosphate groups or carbonate anions. Additionally, although the mineral in all the three dental tissues is basically HA, it chemical composition is different depending on the phase considered (Table 2.4). The carbonate contain in dentine is similar with the amount present in bone (∼5 wt.-%) and significantly higher than that of enamel (∼3 wt.-%). The carbonate content of human enamel increases from the surface going inwards. The magnesium content on dentine is also higher than in enamel (∼1 and 0.2 wt.-% respectively). Also, dentine is less calcium-deficient than enamel, with a Ca/P molar ratio ∼1.66 (very close to the Ca/P ratio for stoichiometric HA, 1.67) and ∼1.59 respectively. In enamel, the F– ion content (∼0.01 vs. ∼0.07 wt.-% in dentine) is sharply peaked at the surface. Table 2.4. Typical composition of HA in enamel and dentinei in wt.-%.9, 12 Ca2+ PO43– CO32– Na+ K+ Mg2+ Sr2+ Cl– F– P2O74– H2OCa/P ratio Enamel 37.6 18.3 3.0 0.7 0.05 0.2 0.03 0.4 0.01 0.022 1.5 1.59 Dentine 40.3 18.6 4.8 0.1 0.07 1.1 0.04 0.27 0.07 0.10 10 1.66 When the OH– group of the HA structure is fully replace by F–, we speak about fluorapatite (FAp, Ca10(PO4)6F2) instead of hydroxyapatite. FAp is the most stable and least soluble compound among all calcium orthophosphates. FAp easily forms solid solutions with HA with any desired F/OH molar ratio. Such compounds are called fluorhydroxyapatites (FHA) or hydroxyfluorapatites (HFA) and described with the chemical formula Ca10(PO4)6(OH)2-xFx (0 < x < 2). Pure FA never occurs in biological systems. However, due to the ability to form FHA and HFA, an exposure to fluoride (like that contained in toothpaste and city tap water) to turn some of the HA into FHA/HFA (more resistant to damage caused by acids) is probably the most effective cavity prevention treatment available today. i Numerical values were not found in the literature for the cementum, but they should be similar to those for dentin. Introduction: Bone repair and substitution 26 2.3.- BONE REPAIR AND SUBSTITUTION 2.3.1. Bone defects Although bones repair themselves when they fracture through a natural remodelling process (Section 2.1.4), there are some circumstances when bone defects do not heal spontaneously. These situations frequently result from trauma, an increasingly aged population, pathological degeneration, congenital abnormities, infections or tumour resection of the bone tissue.13 Hence, filling of the resulting defect is a common practice with an increasing value in the re-establishment of the musculoskeletal system to promote bone healing,13 and, in fact, bone is the second most frequently transplanted tissue in humans.14 Currently, repairing osseous deficiencies involves various medical surgical techniques, including autogenous grafts, allografts, internal and external fixation devices, electrical stimulation, alloplastic implants,i etc.13 The goal of bone replacement is to bridge a bone defect over healing procedures stable and durable without having to accept new problems or complications in order to restore form and function to patients. In general, bone-grafting materials can be divided into two main groups, natural and synthetic bone grafts. 2.3.2. Natural bone grafts. Autogeneous, allogeneous and xenogeneous implants Natural materials used as bone graft materials can be divided into autograft, allograft and xenograft.13, 15  Autogenous grafting involves harvesting healthy bone from one anatomical site and implanting the graft material in a defect site of the same individual. The tissue is usually collected from the iliac crest, but also from the distal femur and proximal tibia.  Allografts are taken from donors or cadavers and implanted into another individual of the same specie.  Xenografts refer to bone tissue collected from one species and implanted into a different one, i.e. ivory or bovine bone. Since von Walter described as early as 1821 the successful clinical application of an autogenous bone transplant for the first time in Germany,13 autogenous bone grafting became the 20th century standard for bone repair and regeneration as its additional functions of serving as a mechanical bridge and providing a scaffold for the ingrowth of new bone became clear.16 i It refers to inorganic materials implanted in living tissue. Introduction: Bone repair and substitution 27 Essentially, autografts possess the advantages of an absence of immunogenic reactioni after surgery and very good biological performance17 in terms of osteogenicity, ii osteoinductivityiii and osteoconductivity.iv However, although autologous bone graft meets many requirements for successful bone regeneration, their use is also associated with some important drawbacks. Harvesting the graft requires an additional invasive surgical procedure that may lead to donor site morbidity (it occurs in ~20% of all cases),15 chronic post-operative pain, extra blood loss from the donor site, infection and cost.17, 18 Furthermore, autografting is ineffective when the defect volume exceeds the volume of healthy graft material. Additionally, a second invasive procedure is contraindicated for geriatric or pediatric patients in a compromise state.18 Another important drawback of the use of autografts is limited availability. Unlike autologous bone, allogeneic and xenogeneic grafts are widely available and do not require additional surgery on the patient. However, nowadays xenogeneic tissues find rather infrequent application in bone grafting owing to concerns with immunogenicity and disease transmission.18 Allogeneic bone has to undergo processing techniques such as lyophilization, irradiation or freeze-drying to remove all immunogenic proteins, in order to avoid any risk of immunogenic reaction. In turn, these processing techniques have a negative effect on the osteoinductive and osteoconductive potential of the allografts, which consequently decreases their biological performance as compared to autografts.17 However, despite the extensive testing of the donor and the bone obtained, a risk of both bacterial and viral transmission from the donor material to the host remains.15, 19 Also in some areas of the world, the practice of cadaveric bone transplantation is culturally unacceptable. Allografts fail more frequently than autografts.18 Nevertheless, use of the various forms of allogenic bone grafts has become commonplace in a wide range of orthopaedic procedures. In fact, the demand for these materials has risen dramatically over the past decade and it is believe that the supply of cadaveric bone graft will become incapable of meeting the epidemiologically driven demand.16 i Immunogenicity is the ability of a particular substance (antigen), to provoke an immune response; a collection of mechanisms within an organism that protects against disease by identifying and killing pathogens. ii It is the supply of bone-forming cells by the marrow of the harvested bone. iii Associated to the process whereby implanted proteins and growth factors induce new bone to grow. iv Osteoconduction refers to the ability of some materials to serve as a scaffold on which bone cells can attach, grow and divide. In this way, the bone healing response is "conducted" through the graft site. They promotes bone apposition on its surface or down into pores or channels functioning in part as a receptive scaffold to facilitate enhanced bone formation. Introduction: Bone repair and substitution 28 Because of the above-mentioned drawbacks of natural bone grafts, large number of synthetic grafts has been developed as a promising alternative to traditional bone grafting.15-18 2.3.3. Synthetic bone grafts: Biomaterials 2.3.3.1. What is a biomaterial? Despite the benefits of both autografts and allografts, the limitations of each have necessitated the development of synthetic materials for bone repairs and replacements. We speak, then, of biomaterials used as bone graft. A biomaterial is an artificial material capable of being implanted in the body. They are used to make devices to replace a part or a function of the body, in safe, reliable, economic and physiologically acceptable manner.20 A variety of devices and materials used in the treatment of disease or injury include such commonplace items as sutures, needles, catheters, plates, tooth fillings, etc. Over the years, various definitions of the term biomaterials have been proposed. For example, a biomaterial can be simply defined as a synthetic material used to replace part of a living system or to function in intimate contact with living tissue.21 Williams defined biomaterials in 1987 as “a nonviable material used in a medical device, intended to interact with biological systems”.22 Other definitions have included “materials of synthetic as well as of natural origin in contact with tissue, blood and biological fluids, and intended for use for prosthetic, diagnostic, therapeutic and storage applications without adversely affecting the living organism and its components”23 and “any substance (other than drugs) or combination of substances, synthetic or natural in origin, which can be used for any period of time, as a whole or as a part of a system which treats, augments, or replaces any tissue, organ or function of the body”.24 By contrast, a biological material is a material, such as skin or artery, produced by a biological system. Artificial materials that simply are in contact with the skin, such as hearing aids and wearable artificial limbs are not included in the biomaterials definition since the skin acts as a barrier with the external world. 2.3.3.2. Hystorical background In 1931, it was found at the Ulúa valley (Honduras) a fragment of a Mayan girl jaw dated from 600 A.D. where three pieces of a shell shaped as teeth were replacing three lost lower incisisors.25 This fact shows that the use of foreigner materials for the replacement and care of damaged parts of the human body is not new. Introduction: Bone repair and substitution 29 The use of biomaterials did not become practical until the advent of an aseptic surgical technique developed by Dr. J. Lister in the 1869s.21 Earlier surgical procedures, whether they involved biomaterials or not, were generally unsuccessful as a result of infection. The earliest successful implants, as well as a large fraction of modern ones, were in the skeletal system. The first documented use of synthetic bone grafts was reported in 1892 by Dreesman, who treated large-bone defects with calcium sulphate hemihydrate or gypsum (CaSO4·0.5H2O, also called plaster of Paris).16 Metallic bone plates were introduced in the early 1900s to aid in the fixation of long bone fractures, but many of these early plates broke as a result of unsophisticated mechanical design. Also, materials such as vanadium steel corroded rapidly in the body and caused adverse effects on the natural healing process.21 Better designs and materials soon followed, as the use of stainless steels and cobalt-chromium alloys in the 1930s, when the first joint replacement surgeries were performed.21 During the 1940s, polymethylmetacrylate was first used for corneal or damaged skull bones replacement. Following further advances in materials and surgical technique, blood vessel replacements were tried in the 1950s and heart valve replacements and cemented joint replacements in the 1960s.21 Dr. W. J. Kolff implanted the first artificial heart in 1982. The role of biomaterials has been influenced considerably by advances in many areas of biotechnology and science. 2.3.3.3. Requirements of biomaterials. Biocompatibility, bioactivity and bioresorbability As discussed before, biomaterials function physically and/or biologically in intimate contact with living tissue. Therefore, the success of biomaterials in the body depends on factors such as the material properties, design and biocompatibility of the material used. Other factors include the technique used by the surgeon and the health, activities and condition of the patient.21 Biocompatibility involves the acceptance of an artificial implant by the surrounding tissues and by the body as a whole. Biocompatible materials do not irritate the surrounding structure, do not provoke a sustained inflammatory or toxic response, do not incite allergic or immunologic reactions and do not cause cancer.21 Other compatibility characteristics which may be important in function of an implant device made of biomaterials include (i) and adequate mechanical properties such as strength, stiffness and fatigue properties, (ii) appropriate optical properties of the material is to be used in the eye, skin or tooth and (iii) appropriate density. Sterilizability and appropriate design are also to be considered.21 Also, it is now known that clinical success of a biomaterial implant requires the simultaneous achievement of a stable interface with connective tissue (bone bonding).8 Introduction: Bone repair and substitution 36 structures and (iii) adopting combinatorial and computational approaches in biomaterial design to accelerate the discovery of novel resorbable polymers.30 Polymeric biomaterials are less stiff and hard than bones, and therefore, they cannot be used in load-bearing implants. In order to increase their mechanical properties, polymers can be mixed with a reinforcement phase such as, i.e., carbon fibres41 or ceramics granules such as HA42 or bioactive glass.43 Fig. 2.14. A poly(glycolide-co-lactide) bone screw (after Jan and Grzegorz).41 2.3.6. Ceramic biomaterials Unlike metals and polymers, ceramics are difficult to shear plastically due to the ionic or covalent nature of the bonding and minimum number of slip systems.i These characteristics make the ceramics non-ductile ii and are responsible for almost zero creep iii at room temperature.44 Consequently ceramics are very susceptible to notches or microcrakcs because instead of undergoing plastic deformation they fracture elastically on initiation of a crack, thus, they are brittle. This is also the reason ceramics have low tensile strength compared to compressive strength. Ceramic are generally hard, i.e., alumina (Al2O3) shows a hardness of 9 in the Mohs’ scale.iv Other characteristics of ceramics are their high melting temperatures and low conductivity of electricity and heat due to the nature of their bonding. In recent years, humans have realized that ceramics and their composites can be used also to augment or replace various parts of the body, particularly bone. Thus, the ceramics used for the latter purposes are classified as bioceramics. Also, their relative inertness to the body fluid, high compressive strength and esthetical pleasing appearance led to the use of ceramics in dentistry as dental crows. Some carbons have found use as implants especially for blood interfacing applications such as heart valves. Due to their high specific strength as fibres, ceramics are also being used as reinforcing components of composite implant materials41-43 and for tensile loading applications such as artificial tendons and ligaments.44 i Slip is the process by which plastic deformation is produced by a dislocation (a crystallographic defect, or irregularity, within a crystal structure motion). By an external force, parts of the crystal lattice glide along each other, resulting in a changed geometry of the material. ii A ductile material is capable of withstanding a certain amount of force, by changing its form, before fracturing or breaking (related with plastic behaviour). iii It is a deformation under a constant load. iv At the Mohs scale, talc Mg3Si4O10(OH)2 is the softest ceramic with a Mohs hardness of 1 and diamond is the hardest (value of 10). Apatite has a Mohs value of 5. Introduction: Bone repair and substitution 37 Other applications of bioceramics are as carriers for enzymes, antibodies and antigens –non-resorbable porous glasses– since they have several advantages, notably resistance to microbial attack, pH changes, solvent conditions, temperature, etc.8 However, due to bioceramics’ inherent brittleness, susceptibility to microcracks, low tensile and impact strength and poor fracture toughness, their use for applications requiring significant load-bearing, torsion bending or shear stress, etc., seems impracticable at present.8 From a structural point of view, biocompatible ceramic materials used in fabricating implants can be classified as:  crystalline solids (ceramics),  amorphous solids (glasses),  amorphous solids with crystallization nuclei (glass-ceramics), which can in turn be considered as (Section 2.3.3.3):8 • non-resorbable (relatively bioinert): not able to bond with the host tissue, • bioactive or surface reactive (semi-inert): able to physically bond directly to the living host bed, • resorbable (non-inert) bioceramics : able to bond with living tissues and to be replaced by the natural host tissue. 2.3.6.1. Non-resorbable bioceramics Non-resorbable bioceramics have almost no influence in the surrounding living tissue and maintain their physical and mechanical properties while in the host. They resist corrosion and wear.i Examples are dense and porous alumina (Al2O3), zirconia (ZrO2) and titania (TiO2) ceramics and single-phased calcium aluminates (xCaO.yAl2O3). Because they do not bond to bone, their application is limited to repairs that will not encounter sheering or torqueii forces.16 They are typically used as dental implants (Fig. 2.15A) (because of its combination of excellent corrosion and high wear resistance), structural-support implants (i.e., bone plates, bone screws, Fig. 2.15B), or for long-bone defects (because of their excellent compressive strength). Examples of non structural support uses are ventilation tubes, sterilization devices and drug delivery devices.45 i In materials science, wear is the erosion of material from a solid surface by the action of another solid. ii Torque is a vector that measures how much a force acting on an object causes that object to rotate. Introduction: Bone repair and substitution 38 Fig. 2.15. Images showing a bioceramic dental implant (A) and a bone screw (B). 2.3.6.2. Bioactive bioceramics Upon implantation in the host, bioactive bioceramics form strong bonds with adjacent tissue. Surface reactive bioceramics are, i.e., HA ceramics, hydroxyapatite/wollastonitei glassceramics and certain compositions of silicate-based glasses (bioactive glasses).8 Bioactive glasses developed for implantation are, mainly SiO2-CaO-Na2O-P2O5 and Li2-ZnO-SiO2 systems.44 Bioactive ceramics have many applications, primarily in the areas of bone repair and bone regeneration via tissue engineering, such as:  synthetic bone graft materials for general orthopaedic, craniofacial, maxillofacial and periodontal repair (Fig. 2.16),  cochlear implants, or  bone tissue engineering scaffolds (porous bioceramics). Fig. 2.16. X-ray image of a dog’s periodontal bone defect before (A) and after (B) filling with a bioactive glass. i CaSiO3 Introduction: Bone repair and substitution 39 Also, bioactive ceramics are employed to obtain bioactive composites. As an example, bioactive glass has been used in combination with PMMA to form a bioactive bone cement, and HA and glasses are used as a coating to form a layer on metallic implants’ surface that eases the chemical bonding of the implant to the surrounding bone tissue (Section 2.3.4).8, 15 A common characteristic of bioactive glasses and bioactive ceramics is a time-dependent, kinetic modification of the surface that occurs upon implantation into living bone tissue. The surface forms a biologically active carbonated HA layer, which provides the bonding interface with osseous tissues (bone bonding), and is chemically and structurally equivalent to the mineral phase in bone. It is that equivalence which is responsible for interfacial bonding (Section 2.3.3.3).8 2.3.6.3. Resorbable bioceramics Resorbable bioceramics degrade upon implantation in the host and the resorbed material is replaced by endogenous tissues,i being this its principal advantage. The rate of degradation varies from material to material. Almost all bioresorbable ceramics, except biocoral (CaCO3) and gypsum (CaSO4·0.5H2O), are variations of calcium orthophosphates, including aluminium-calcium-phosphorous oxide, ferric-calcium-phosphorous oxide, zinc-calcium-phosphorous oxides, zinc-sulphate-calcium-phosphorous oxides, tricalcium phosphate (TCP, Ca3(PO4)2) and HA.44 Fig. 2.17. Examples of calcium orthophosphate-based bone substitution materials (granules and scaffolds) (after Dorozhkin).12 As for non-resorbable or bioactive bioceramics, the mechanical behaviour of resorbable bioceramics strongly influences their use as implants, and therefore, they must be shielded from loading forces until bone ingrowth has occurred. Rigid stabilization of i Endogenous substances are those that originate from within an organism, tissue, or cell. It is the opposite of exogeneous. Introduction: Bone repair and substitution 40 surrounding bone and non-weight bearing are required during this period because ceramics can only tolerate low bending and torque load before failing, unless they are used in sites of relatively low mechanical stress or when forces are basically compressive.16 Applications include dental implants, percutaneous devices, periodontal treatment, drug delivery devices, alveolar ridge augmentation, orthopaedic and maxillofacial surgery, otolaryngology, spinal surgery, bone defect filling, excised tumour and diseased bone loss, scaffolds (Fig. 2.17), etc. Although gypsum –or plaster of Paris– was used already in 1892 as bone substitute,16 the concept of using synthetic resorbable ceramics as bone grafts was introduced in 1969.46 Gypsum biocompatibility, bioactivity and resorbability make it a desirable candidate for bone graft substitutes, but its degradation time –between 30 and 60 days– is too short for bone repair. Also, there is significant loss of mechanical properties upon degradation, so it does not provide internal strength or support. Therefore it can only be used to fill small bone defects such as those resulting from cyst curettage.15, 16 Nevertheless, calcium sulphate has recently seen resurgence in use with the recent marketing of this material in the form of tablets for use in filling osseous defects.16 Calcium orthophosphates make up the majority of ceramic-based bone graft substitutes currently available on the market, namely TCP, hydroxyapatite or the combination of both.15, 28 Calcium orthophosphate-based bioceramics have been in use in medicine and dentistry for nearly forty years, mainly with bone repair purposes.8 The first in vivo use of calcium orthophosphates was performed in 1920; that time the researchers implanted TCP into animals to test its efficacy as a bone substitute.47 In 1951, HA was tested in vivo for the first time. However, it was already the 1970s, when other calcium orthophosphates were synthesized, characterized, investigated, and tried in medicine.47 An advantageous property of these ceramics, particularly HA, is that they are osteoconductive and they bond well to the host bone (osseointegration). i Additionally, although these ceramics do not contain intrinsic osteoinductive properties, the bond between host and graft provides sufficient affinity for local or exogenousii growth factorsiii that serve in the regeneration process.16 Like HA, TCP is biocompatible, bioactive and bioresorbable, but its resorption rate is 10-20 times faster than the HA one.48 When a biphasic calcium i It is the direct structural and functional connection between living bone and the surface of an artificial implant. ii In biology, an exogenous factor is any factor that is present and active in an individual organism or living cell but that originated outside of that organism. It is the opposite of endogenous. iii Growth factor refers to a naturally occurring protein capable of stimulating cellular differentiation and maturation. For example, bone morphogenic proteins stimulate bone cell differentiation. Introduction: Bone repair and substitution 41 orthophosphate ceramic, consisting of HA and β-TCP, was compared with a pure HA ceramic with similar macroand microstructure, more bone was found in the TCP more soluble phase. This suggests that the more soluble calcium orthophosphate phases are more osteoinductive than less soluble ones.17 However, implantation of highly soluble carbonated apatite ceramic did not result in bone induction, suggesting that a relatively stable surface is needed for facilitation of bone formation.17 Therefore, a combination of a mixture of both HA and TCP phases might be clinically favorable.16 HA and TCP bioceramics are obtained by sintering calcium orthophosphate salts at high temperatures (1000-1500°C) to produce a powder that can then be moulded into the desired shape by high-pressure compaction.8, 16 There is little control over the porosity of this material, and hence, the ceramic remains stable (not resorbed) for long periods of time due to its high density –radiographic findings demonstrate a continued presence of the calcium orthophosphate ceramic for a prolonged period of time due to the failure of complete remodelling.16 Osteoconduction is a 3D process that is observed when porous structures are implanted into or adjacent to bone. Porosity alone, however, is not adequate for bone ingrowth. Open porosity –connected to the outside surface– with interconnectivity is the most essential prerequisite. This is based on the 3D interconnections between the lacunae in the bone that provide intercellular communication (Section 2.1.2.1).35 Hence, open porosity is critical to the integration of tissue into the ceramics, and porous calcium orthophosphate bioceramics are widely studied.8, 28, 44 The consensus of research indicates that the requisite pore size for bone ingrowth into porous implants is 100 to 500 μm and the interconnections must be larger than 100 μm.8, 16, 35 However, when designing a porous resorbable bioceramic, it should be take into account that, with increased porosity, the graft exhibitss significantly less compressive strength and fatigue resistance.8 Recently, it was predict that calcium orthophosphate bioceramics designed to mimic the bone mineral and intended for use as implants in vivo should not possess steps of heating/firing/calcination at or above 650°C in any phase of their processing, manufacturing and/or shaping operations.29 The reason is that at, or above, this temperature, carbonate ions which may be present in the HA structure are opt to readily leave the materials. The same also applies to the case of HPO42– ions. Both CO32– and HPO42– ions are present in biological HA (Section 2.1.1.2), and therefore, researches should incorporate them when preparing bioceramics grafts in order to prepare materials that will resemble biological HA to the most possible extent.29 Na+, K+ and Mg2+, which altogether amount to a value greater than 1 wt.-% Introduction: Calcium orthophosphates 42 in the bone HA (Section 2.1.1.2), must also be consider in preparing synthetic bioceramic bone substitute materials.49 In sharp contrast to the aforementioned sintered bioceramics, calcium orthophosphates are also used as biomaterials for bone repair and substitution in the form of calcium phosphate bone-cements –a mixture of calcium orthophosphates that reacts in an aqueous medium at room temperature to form HA or dicalcium phosphate dihydrate (DCPD, CaHPO4·2H2O). They are biocompatible, bioactive and resorbable. 2.4.- CALCIUM ORTHOPHOSPHATES Diverse combinations of oxides of calcium and phosphorus (both in the presence of water and without it) provide a large variety of compounds which are distinguished by the type of the phosphate anion: ortho- (PO43–), meta- (PO3– ), pyro- (P2O74–) or polyphosphate ((PO3)nn–). In the case of multi-charged anions (orthophosphates and pyrophosphates), calcium phosphates are also differentiated by the number of hydrogen ions present in the anion. Examples include mono- (Ca(H2PO4)2·H2O), di- (CaHPO4), tricalcium phosphate (Ca3(PO4)2) or calcium pyrosphosphate (Ca2P2O7). All these varied compounds form the family of the calcium phosphates. With the exception of calcium pyrophosphate, calcium phosphates used in medical applications are namely calcium orthophosphates.48 By definition, all calcium orthophosphates consist of three major chemical elements: calcium (oxidation state +2), phosphorus (+5) and oxygen (–2). In addition, the chemical composition of many calcium orthophosphates includes hydrogen, either as an acidic orthophosphate anion (for example, HPO42– or H2PO4–), and/or as incorporated water (for example, CaHPO4·2H2O). In the ternary system Ca(OH)2-H3PO4-H2O there are eleven known non-ion-substituted calcium orthophosphates with the Ca/P molar ratio within 0.5 and 2.0 (Table 2.5). It should be noted that in literature occasionally one could find brief notes on the 12th calcium orthophosphate, namely oxyapatite (Ca10(PO4)6O). A mixture of oxyapatite and HA might be prepared by dehydration of HA at temperatures exceeding ~900°C (i.e., during plasma spray of HA) only in the absence of water vapour. It also might be crystallized in glass-ceramics. Oxyapatite is very reactive and transforms to HA in contact with water vapour. It is still very poorly known; however, data on the solubility constant and crystal structure are available.12 Introduction: Calcium orthophsophates 43 Table 2.5. The members of the calcium orthophosphate family with their respective Ca/P ratio and pH stability range in aqueous solutions at 25°C. Common name Acronym Formula Ca/P pH range12 Monocalcium phosphate anhydrous MCPA Ca(H2PO4)2 0.5 e Monocalcium phosphate monohydrate MCPM Ca(H2PO4)2·H2O 0.5 0.0-2.0 Dicalcium phosphate anhydrousa DCPA CaHPO4 1.0 e Dicalcium phosphate dihydrateb DCPD CaHPO4·2H2O 1.0 2.0-6.0 Amorphous calcium phosphate ACP CaxHy(PO4)z·nH2O n = 3-4.5; 15-20% H2O 1.2-2.2 ~5-12f Octacalcium phosphate OCP Ca8(HPO4)2(PO4)4·5H2O 1.33 5.5-7.0 α-Tricalcium phosphate α-TCP α-Ca3(PO4)2 1.5 g β-Tricalcium phosphatec β -TCP β -Ca3(PO4)2 1.5 g Calcium deficient hydroxyapatite CDHA Ca10-x(HPO4)x(PO4)6-x(OH)2-x (0 < x < 1) 1.5-1.67 6.5-9.5 Hydroxyapatite HA Ca10(PO4)6(OH)2 1.67 9.5-12 Tetracalcium phosphated TTCP Ca4(PO4)2O 2.0 g a– mineral: monetite; b– mineral: brushite; c– mineral: whitlockite; d– mineral: hilgenstockite; e– stable at temperature above 100°C; f– always metastable; g– cannot precipitate from aqueous solutions. The most important parameters at the calcium orthophosphates are the molar Ca/P ratio, basicity/acidity and solubility. These parameters strongly correlate with the solution pH. The lower the Ca/P molar ratio is, the more acidic and water-soluble the calcium orthophosphate is.9 Hence, MCPM is both the most acidic and water-soluble calcium orthophosphate and TTCP is the most basic. However, TTCP solubility in water is higher than that of HA (log(Ks)25°C = 38-44 vs. 116.8, respectively).12 TTCP is not very stable in aqueous solutions –it hydrolyses50 to HA and Ca(OH)2. The vast majority of calcium orthophosphates are sparingly soluble in water;9 however, all of them are easily soluble in acids but insoluble in alkaline solutions. Chemically pure calcium orthophosphates are crystals of white colour and moderate hardness. However, natural minerals of calcium orthophosphates are always coloured due to impurities. The atomic arrangement of calcium orthophosphates is built up around a network of PO43– groups, which gives stability to the whole structure.51 Introduction: Calcium orthophosphates 44 2.4.1. Solubility phase diagram As mentioned above, calcium orthophosphates comprises eleven known non-ion-substituted compounds. Each of them can be synthesized at different conditions, i.e. those with H2PO4– ions only form under rather acidic conditions.9 Due to the triprotic (three protons) equilibrium that exists within orthophosphate-containing solutions (Eq. 2.1), variations in the pH solution alter the relative concentrations of the four orthophosphoric acid polymorphs and, thus, both the chemical composition and the amount of the calcium orthophosphates that forms by direct precipitation.52 −−− ⎯→←⎯→←⎯→←3 4 2 44243 3 21 POHPOPOHPOH k kk (Eq. 2.1) where12 k1 = 7.5·10– 3; k2 = 6.2·10–8 and k3 = 1.7·10–12. Solubility diagrams show the thermodynamically stable phases at given conditions, and provides an indication of the likely conditions required for the synthesis of a determinate calcium orthophosphate. The solubility diagram53 for the ternary system Ca(OH)2-H3PO4-H2O at 25°C is given in Fig. 2.18. Fig. 2.18. Diagram of solubility isotherms of calcium orthophosphates phases in the ternary system Ca(OH)2-H3PO4-H2O at 25°C.53 Gibbs' phase rule (Eq. 2.2) is the fundamental rule on which phase diagrams are based. It provides the number of degrees of freedom of a system for a given thermodynamic condition, that is, how many control variables (pressure, temperature, composition, etc.) can be altered while maintaining this condition. Mathematically, this number is: F = 2 − P + C (Eq. 2.2) where F is the degrees of freedom, P is the number of phases that coexist in the equilibrium (solid, liquid, gas, etc.) and C is the number of components in the system. The number two Introduction: Calcium orthophsophates 45 arises from the two main thermodynamic control parameters, usually temperature and pressure. If other variables such as magnetic, electric or gravitational fields are considered, the numerical value two will increase by the number of new variables considered. According to Gibbs’ phase rule, a ternary system as Ca(OH)2-H3PO4-H2O with a single phase, i.e. liquid (solution), has two degrees of freedom at a fixed temperature and pressure; that is, two variables (pH, concentration, composition, etc.) can be altered while keeping this condition. Hence, the composition of any solution in this system is defined by fixing, i.e., pH and calcium concentration [Ca], or pH and [P], or [Ca] and [P]. The solubility diagram at Fig. 2.18 represents the pH vs. [Ca] diagram of the calcium orthophosphates at a fixed pressure and temperature. According to this diagram, HA is the phase thermodynamically stable at pH ~4.5-10, and therefore, it will be the calcium orthophosphate precipitated in this pH range. However, it should be take into account that the actual phase that forms under any given conditions is often also influenced by kinetic considerations.9 The solubility diagram also reflects the fact that all calcium orthophosphate compounds are more soluble as the pH decreases. The slope of the isotherm is an indication of how fast the solubility of the salt increases with decreasing pH. Since, for a given drop in pH, the solubility of a basic salt would increase more than would an acidic salt, the slope of the isotherm is related to the alkalinity of the salt. Consequently, the more acidic salts, DCPD and DCPA, have smaller negative slopes than do the more basic salts TTCP, HA and the two TCPs. At the invariant or singular points (point where two solubility isotherms across), three phases coexist in equilibrium. As an example, DCPD-TTCP invariant point occurs at ~8.5 pH. This equilibrium is characterized by a continuous saturationi of the liquid phase with respect to DCPD and TTCP, which are equally stable at this solution composition. However, their relative stabilities are different –DCPD is more stable than TTCP at pH below 8.5 and vice versa. At this invariant point, the solution is supersaturatedii with regards to HA (the thermodynamically stable phase at this pH). Thus, HA will precipitate under these conditions while DCPD and TTCP will keep in solution. As long as DCPD and TTCP are present in excess, the solution pH will remain close to the DCPD-TTCP singular point. Therefore, the i Is the point at which a solution of a substance can dissolve no more of that substance. ii It is a solution that contains more of the dissolved material than could be dissolved by the solvent under normal circumstances. Introduction: Calcium phosphate bone-cements 52 Hence, CPCs are becoming of increasingly great importance in the field of biomaterials, in particular, as bone substitutes and repair (Fig. 2.19).68 CPCs also have prospective applications in the field of drug delivery devices.67 Because of their remarkable insertion properties, CPCs and related biomaterials have been studied intensively over the last twenty five years and, as a result, a number of different formulations have been established that are currently used in medical practice.69 Consequently, several reviews covering the preparation, properties and applications of CPC biomaterials are available,47, 67-72 as well as more than a dozen commercial CPC formulations.69 However, it should be emphasized that owing to the lower fracture toughness parameters of CPCs73 in comparison to human bone8 (~0.6-1.5 cf. 2-12 MPa/m2), the clinical applications of CPCs are limited to areas where bones are free of dynamic load, i.e. for non-load bearing as well as craniofacial and periodontal applications,74 or as materials for the development of scaffolds in bone tissue engineering.75 (Table 2.7) Fig. 2.19. SEM image illustrating CPCs application for bone repair (after Ni et al.)76 2.5.1. Types of calcium phosphate cements As mentioned before, generally, all CPCs can be generated by mixing one or several reactive calcium orthophosphate precursors in an aqueous medium. After mixing, a putty-like material/injectable liquid is formed, which is able to set and harden in vivo. Consequently, CPCs are widely applicable in clinical practice, i.e., for filling bone defects during surgery, resulting in reduced patient discomfort in comparison to other techniques currently employed.77 Introduction: Calcium phosphate bone-cements 53 In 1990s, it was established that there were about fifteen different binary combinations of calcium orthophosphates which gave pastes upon mixing with water or aqueous solutions, so that the pastes set at room or body temperature into a solid cement.47 From these basic systems, secondary formulations could be derived containing additional or even non-reactive compounds, but still setting like cements.55, 57 Most of the cement formulations after the setting reaction result in the formation of CDHA, while the remaining cement formulations give DCPD (brushite) as the end product,70 since at body temperature only these two calcium orthophosphates are stable in contact with aqueous media such as body fluids.i DCPD is the stable phase a pH < 4.2, whereas at pH > 4.2 HA is the stable phase (Table 2.5, Fig. 2.18).53 A striking difference between these two types of CPCs (called apatitic or brushite CPCs) is that the solubility at physiological pH of the brushite cements is approximately two times higher compared with the HA cements,78 resulting in their enhanced resorption by biological serum.57 However, DCPD is a metastable phase which can hydrolyses to apatite in vivo and thus may potentially generate biocompatibility problems.79 In addition, upon examination of the literature related to brushite CPCs, it is apparent that overall they exhibit a lower mechanical strength compared to the apatite CPCs.80 Other advantages of the apatite CPCs are: thermodynamic stability,71 reasonable setting time with respect to surgeons’ requirements70 and the formation of CDHA as the end product72 –the inorganic component of mineralized bone tissues. As-produced apatite CPCs are chemically and structurally very similar to biological apatites, exhibiting excellent biological and physiological characteristics. For the majority of apatite cements, water is not a reactant in the setting reaction. Therefore, the quantity of water actually needed for setting of apatite cements is very small.81 However, for brushite cements, water always participates in the chemical transformations because it is necessary for DCPD formation. Due to this reason, brushite cements are always hydraulic,ii while usually this term is not associated with apatite cements. Despite their very large number, the most studied CPC formulations are based on the following reactions (Eq. 2.3 to 2.5):70 β-Ca3(PO4)2 + Ca(H2PO4)2·H2O + 7H2O → 4CaHPO4·2H2O (Eq. 2.3) i ACP and OCP appear in humans as precursor in the initial phase of enamel and bone mineral formation. However, they appear as an unstable transient intermediate phase (Section 2.4.1.1). ii Hydraulic cements are materials that set and harden, after being combined with water, as a result of chemical reactions with the mixing water. Introduction: Calcium phosphate bone-cements 54 It is an acid-base reaction, where β-TCP acts as base an MCPM as acid. They react to form DCPD, a neutral calcium orthophosphate 2Ca4(PO4)2O + 2CaHPO4 → Ca10(PO4)6(OH)2 (Eq. 2.4) In this acid-base reaction TTCP (basic character) reacts with DCPA (neutral) to form HA (slightly basic). 3α-Ca3(PO4)2 + H2O → Ca9(HPO4)(PO4)5(OH) (Eq. 2.5) The slightly basic α-TCP hydrolyzes to CDHA (slightly basic) via a dissolutionprecipitation process without pH change. Therefore, all existent CPC formulation reacts, well by an acid-base reaction (Eq. 2.3 and 2.4), well by a conversion (Eq. 2.5) to give as the end-product either DCPD or HA (Ca-deficient or stoichiometric). 2.5.2. Control of the setting reaction. Cohesion and setting time Generally, CPC reaction rate must be well controlled: the CPC must react slowly enough to provide enough time to the surgeon for implantation, and fast enough to prevent delaying the operation. In the clinical situation, the surgeon needs to place the CPC paste into the prepared surgical site and then wait for the paste to sufficiently set before closing the wound. Otherwise the soft paste could be deformed easily to lose its geometrical shape by the stress generated in closing the wound, or even be disintegrated and fail to set into a cohesive implant.82 Thus, from a fundamental and an application point of view, it is important to design a rational synthetic route to CPC biomaterials with setting time between 10-30 minutes.47, 70 The setting time of a CPC is divided in initial (I) and final (F) setting time. The initial one marks the onset of hardening after water has been added, while the final one refers to the fact that, after this time, the setting reaction has proceeded to an extent that the manipulation of the CPC will not deteriorate its mechanical properties. Clinically, it states that the cement paste should be implanted before time I and that the wound can be closed after time F (Fig. 2.20). A cement should not be deformed between times I and F because in that stage of the setting process any deformation could induce cracks.57 The following handling requirements (in minutes) have been formulated for CPCs: 3 ≤ I < 8; F ≤ 20.47 Introduction: Calcium phosphate bone-cements 55 Another requirement of CPCs is that the cohesion timei (CT) must be at least 1 min before the initial setting time so that a clinician has at least 1 min to apply and to mold the material. As the mixing in a mortar is about 1 min, the shortest CT that can be allowed is about 2 min, so that a clinician has at least 1 min to collect the paste from the mortar and put it on the pallet knife or in the syringe with which it is to be transferred to the wound after CT and before I (Fig. 2.20).47 Fig. 2.20. Diagram of the setting parameters relevant for a CPC: CT – cohesion time; I – initial setting time; F – final setting time (after Dorozhkin).47 Additionally, it is noteworthy that in many cases the rapid conversion of the CPC precursors to HA not only leads to the rise of the appropriate physiological properties, but also results in the strengthening of the mechanical properties. In particular, rapid conversion accelerates achievement of the final compressive strength value, which is almost linearly dependent on the extent of the CPC setting reaction.83 Therefore, a rapid conversion rate to the end product might be also desirable. The setting reaction of CPCs consists of three stages, namely, (i) dissolution of reactants, (ii) nucleation of crystals, and (iii) growth of crystals. Therefore, the strategies to modify the setting reaction of CPCs have been targeted to these three stages of the setting process.70 Table 2.8 gathers the approaches and the employed tactics to modify CPCs reactivity. 70 i It is the time from which a cement no longer disintegrates when immersed in an aqueous phase. 56 Table 2.8. Strategies and approaches of the employed tactics to modify CPCs’ setting reaction rate.70 Strategy Approach How? Change contact area between reagent and mixing liquid – Milling of reactants – Use nano/micro powders Change solubility in the mixing liquid – Use more/less soluble phase – Change of reaction pH Change mixing liquid saturation Use a liquid phase containing Ca2+ and/or phosphate ions Use dissolution inhibitors in the mixing liquid Dissolution rate Modify reagent surface – Chemical change (pre-reaction) – Physical change (dissolution pits) Use seeds Change the saturation of the reaction product in the mixing liquid – Change of saturation – Change of end-product solubility Nucleation rate Use of nucleation inhibitors Change the saturation of the reaction product in the mixing liquid Change of saturation Growth rate Use crystal growth inhibitors – Change of saturation – Change of end-product solubility Introduction: Calcium phosphate bone-cements 57 2.5.2.1. Dissolution rate of reactants Experimental studies have revealed that the setting reactions based on Eq. 2.4 and 2.5 are initially controlled by surface reactions, and then by diffusion.84, 85 In brushite CPCs (Eq. 2.3), the very strong effect of β-TCP specific surface area on β-TCP–MCPM–water CPCs suggests that a surface reaction is also the limiting factor, at least at the start of the brushite CPCs setting reaction.70 Therefore, the control of the reagent dissolution rate is a very important aspect of the overall CPC setting reaction. The dissolution rate of reactants can be modified by various means (Table 2.8):70 K Increasing the surface area between reagent and liquid phase, K Changing the reagent solubility in the liquid phase, K Changing the saturation of the liquid phase towards the reagent, K Using dissolutions inhibitors, K By surface modification (passivation or activation of the surface).  The surface area between reagent and liquid phase can be increase by milling of the reagents or by using nano/micro powders. Unfortunately, nano-powders tend to agglomerate, and therefore a large amount of water is required to wet the powder.70 On the other hand, by milling, the particle size of the starting powder does not decrease further than a few micrometers.50 Beyond this point, powder amorphization occurs.50, 86  The solubility of the reagent in the mixing liquid can be varied, either by choosing a more (or less) soluble reagent, or by changing the mixing liquid composition. For example, the replacement of β-TCP by the more soluble α-TCP, or the less soluble HA in β-TCP–MCPM–water cements can lead to much shorter or much longer reaction times.87 The dissolution rate of CPC reagents can also be modified with a shift of the pH value of the mixing liquid. However, it should be taken into account that HA and DCPD only form at specific pH values.  The saturation of the liquid phase towards CPC reagents can be significantly increase using a liquid which already contains phosphate ions (i.e. (NH4)H2PO4(aq), NaH2PO4(aq), KH2PO4(aq)).88 However, the use of a liquid phase saturated with phosphate ions is somewhat limited owing to the incorporation of foreign ions (i.e. NH4+, Na+, K+) into the crystal lattice of the end product. Additionally, such CPC systems have an excess of non-participant reactive ions (i.e. phosphate) and the low or high initial pH of the mixture is also prone to forming intermediate phases.71, 72 These effects can potentially create biocompatibility problems with as-prepared CPCs as well as slow their conversion rate. Introduction: Calcium phosphate bone-cements 58  The use of inhibitors that allow controlling the dissolution rate of reagents can also modify CPCs setting reaction. As an example, HA dissolution is inhibited by the presence of compounds such as bisphosphonates.89  Recently it was suggest that the nucleation of dissolution pits at the material surface is a very important factor controlling material dissolution.90 As an example, an approach to reduce the dissolution rate via surface modification is to partially react the powders to form a dissolution barrier.70 2.5.2.2. Nucleation rate The nucleation rate of new HA or DCPD crystals during the CPC setting reaction can be modified by (Table 2.8): K The addition of nuclei (seeds) in the cement paste, K The modification of the saturation of the mixing liquid, K The use of nucleation inhibitors.  By adding crystal nuclei as seeds into the CPC formulation, it is possible to accelerate its nucleation rate.91  By using more soluble reagents or by adding predissolved calcium and phosphate ions in the mixing liquid, nucleation can also be modified. For example, since fluoroapatite is less soluble than HA, the presence of fluoride ions in the mixing liquid of an apatite CPC accelerates the whole setting reaction.70  The use of nucleation inhibitors strongly reduces the nucleation rate (i.e., Mg2+ ions in apatitic CPCs).85, 92 2.5.2.3. Growth rate Since the means to change growth and nucleation rate are the same, the tactics employed to control the growth rate are very similar with the used to modify the nucleation rate. Among all substances affecting apatite crystal growth, several are of particular importance due to their relevance in mineralization, for example Mg2+ ions85, 92 and carbonate ions70 which are known to inhibit hydroxyapatite growth and also prolong the setting reaction of α-TCP-based CPCs. It should be noted that strategies that increase the dissolution rate of CPC reagents can have a negative influence on the rate of crystal nucleation and growth.88 Introduction: Calcium phosphate bone-cements 59 2.5.3. Future developments of calcium phosphate bone-cements47, 70 The discovery of CPCs has already opened up new perspectives in synthesis of bioceramic scaffolds possessing sufficient mechanical properties. In the past, such scaffolds could only be manufactured by the sintering route at elevated temperatures. Therefore, until recently it was impossible to produce resorbable preset low-temperature hydrated 3D ceramics for various applications, i.e. scaffolds and granules from low-temperature calcium orthophosphate phases, such as ACP, DCPA, DCPD, OCP and CDHA. Now, using the appropriate techniques, open macroporous 3D scaffolds consisting of the aforementioned low-temperature phases (currently, excluding ACP and OCP) can be produced via a cementation reaction, thus noticeably widening the application of these calcium orthophosphates as biomaterials. This type of materials could be very promising for tissue engineering applications. Among them, CDHA is of a special interest due to its chemical similarity to bone material and a large specific surface area. As CPCs and cement-based composites represent an intriguing group of new materials for bone augmentation and reconstruction, there is a great potential for further improvement of their properties, in which the ideal characteristics (Table 2.7) should be achieved by manipulations with the chemical composition, powder particle size and distribution, as well as by means of various additives. Few commercial apatite cements are currently approved for clinical use.69 New formulations of both apatite and brushite cements are expected to appear in the market soon. New formulations will include (i) injectable and open macroporous formulations to optimize their osteoconduction, (ii) formulations containing only one calcium orthophosphate (single-phase cement powders), and (iii) drug-loaded and hormone-loaded cements for the treatment of bone diseases. The former two directions deal with both chemistry and material science, while the last direction is more related to tissue engineering and medicine. The most promising direction of the future developments of CPCs is obviously seen in their functionalization by incorporation or impregnation of various hormones, growth factors, drugs or other biorganic compounds; as well as incorporation of living cells and other tiny biological objects. The initial attempts have already been performed but without a great success yet. For example, researchers have already found that unset CPCs might have toxic effects when placed on cell monolayers, while the set cements are biocompatible for the same type of cells. Stability (insolubility) in normal physiological fluid environment and resorbability under acidic conditions produced by osteoclasts are also among the most important in vivo Introduction: Reinforcement of bioceramics and calcium phosphate cements 60 characteristics of modern CPCs. For some clinical applications, such as cranioplasty, a relatively slow resorption and replacement by bone is quite acceptable, whereas in other applications, such as periodontal bone defects repair, sinus lift, etc., the ability of the hardened cement to be replaced quickly by bone is crucial. Experimental results suggest that a number of parameters of CPCs, such as Ca/P ratio, carbonate content, ionic substitution, crystallinity, etc., might affect the dissolution characteristics of the cements in slightly acidic solutions. This gives an opportunity to formulate cements possessing different resorption rates, which are suited for different applications. Finally, besides the aforementioned chemical, material, and biomedical improvements of CPCs, one should not forget on a better design of both the mixing equipment and delivery (injection) techniques. 2.6.- REINFORCEMENT OF BIOCERAMICS AND CALCIUM PHOSPHATE CEMENTS Bones are stiff and tough because they combine the “good” properties of the collagen (tough but not stiff) and HA (stiff but not very tough) present in their structure (Section 2.1). However, compared with bones, bioceramics and CPCs used for their repair and substitution are brittle, with both a low impact resistance and a low-tensile strength, i.e., fracture toughness for synthetic HA73 with respect to bones8 is 0.6-1.5 vs. 2-12 MPa/m2, respectively. They are strong enough only at compression strength. Consequently, the use of HA ceramics and CPCs for clinical practice is at present limited to low-load bearing areas. Therefore, numerous studies have been developed to investigate the reinforcement of bioceramics materials, and improvements in strength have been reported due to the use of various additives –i.e. phosphate88 or chitosan;93 by increasing the solid to liquid ratio in CPCs;94 by a previous thermal treatment of the CPC precursors;95 by using fibres to reinforce the cement matrix;47 etc. The use of fibres also increase the fracture toughness of the final products.96 Increasing of pressure during the sintering step also improve the strength and fracture toughness of HA ceramics.97 When a reinforcement phase is mixed with bioceramics or CPCs, we speak about biocomposites. Introduction: Reinforcement of bioceramics and calcium phosphate cements 61 2.6.1. What is a composite?. Biocomposites Composite materials are solids which contain two or more distinct constituent materials or phases on a scale larger than the atomic and in which properties are significantly altered in comparison with those of a homogeneous material.98 Composite materials offer a variety of advantages in comparison with homogeneous materials, as in the case of bone (a natural composite,i Section 2.1) or HA-coated metallic implants (Section 2.3.4). The properties of a composite material depend very much upon structure, i.e. bone hierarchical structure (Section 2.1.1.1). In particular, they depend upon the shape of the heterogeneities, the volume fraction occupied by them and upon the interface among the constituents.98 The shape of the heterogeneities in a composite material can be (i) particles with no long dimension (spherical, ellipsoidal, polyhedral or irregular), (ii) fibresii and (iii) platelets.iii In isotropic systems,iv stiff platelet inclusions are the most effective in creating a stiff composite, followed by fibres. The least effective geometry for stiff inclusions is the spherical particles.98 On the contrary, when the inclusions are more compliant that the matrix, spherical particles reduce the stiffness the least and platelet particles reduce it the most.98 Soft spherical inclusions are used intentionally as crack stoppers to enhance the toughness of polymers such as polystyrene with a small sacrifice in stiffness.98 Apart from the mechanical properties of the reinforcing material, the orientation of the incorporated fibres and the matrix-fibre interface are also of great importance for the resulting mechanical properties of the composite. In biomaterials, it is important for each constituent of the composite to be biocompatible. Moreover, the interface between constituents should not be degraded by the body environment.98 Composites employed as biomaterials are often called biocomposites. 2.6.2. Fibrous reinforcement concept The idea behind reinforced ceramics is simple: if a strong filler is present in the matrix, it might stop crack propagation. The concept of using fibres as reinforcement for brittle materials is not new. Fibres have been used as reinforcement since ancient times, when straw was used in mud huts to improve structural integrity. Today, this technique has advanced considerably in the i Other natural composites include dentine (Section 2.2.1.2), cartilage, skin and wood. ii Particles with one long dimension. iii Particles with two long dimensions. iv They have identical values of a property in all crystallographic directions. Introduction: Hydroxyapatite crystals 68 Fig. 2.24. Schematic representation of the hydroxyapatite Ca10(PO4)6(OH)2 structure along the [001] axis (hexagonal, a=9.4320 Å, c=6.8810 Å, space group P63/m (No. 176)). Ca atoms: red, P atoms: yellow; O atoms belonging to PO4 tetrahedra: violet; O atoms from hydroxyl groups: blue. of F– ion (Section 2.1.1.2). ClAp, substitutedand non-stoichiometric-ClAp are of interest as model compounds for studying the crystal chemistry of apatites because single crystals are easily grown. Biological apatites also contain Cl– ions (Section 2.1.1.2). 2.7.2. Hydroxyapatite 2.7.2.1. Crystal structure. Differences in surface reactivity. The formula for hydroxyapatitei can be represented as Ca5(PO4)3(OH), but is usually written as Ca10(PO4)6(OH)2 to denote that the crystal unit cell comprises two molecules. Chemically pure, stoichiometric HA crystallizes in the monoclinic space group P21/b with lattice parameters a = 9.4320, b = 2a, c = 6.8810 Å , γ = 120°.9 However, at temperatures above 250°C, there is a monoclinic to hexagonal phase transition in HA9 (space group P63/m, bypiramidal hexagonal, a = b = 9.4320, c = 6.8810 Å, Fig. 2.24).12 Some impurities, such as the partial substitution of hydroxide by fluoride or chloride anions, stabilize the hexagonal structure of HA at ambient temperature. For this reason, the very rare single crystals of natural HA always exhibit a hexagonal space group.12 The hexagonal structure is the form encountered most frequently and involved in bone formation, because it allows for much easier exchange of OH– groups with other anions, such as F–, Cl– and CO32–. Therefore, we will focus on hexagonal HA (Fig. 2.24). However, from the thermodynamic aspect, the monoclinic form is more stable.119 i The chemically correct name would be hydroxylapatite or hydroxidapatite, while hydroxyapatite is the common name employed by the medical and material communities. Introduction: Hydroxyapatite crystals 69 In the monoclinic phase, the OHs occur in columns on the screw axis,i pointing upward and downward in alternate, nearest-neighbour dyads (Fig. 2.25). The hexagonal form of HA can be derived from the monoclinic form by allowing the OHs to become statistically disordered about the associated Ca ion triads, centred on the screw axis. However, within the P63/m space group, is not possible to simulate the hexagonal HA because of the non-physical duplication of each OH– group by the m plane (Fig. 2.25). To avoid this situation, many studies on the HA hexagonal structure reduce the symmetry to a P63 space group that assumes only one OH– group per Ca ion triangle.119, 120 Within the P63 group, all the OH– groups maintain the same alignment in each column within the HA structure (Fig. 2.25); other structural features remain very similar. Fig. 2.25. Scheme showing the different HA structures: monoclinic P21/b (top), experimental hexagonal P63/m (middle) and theoretical adopted hexagonal P63 (bottom) unit cells. P, O and Ca atoms not relevant for the scheme have been omitted for clarity. Ca atoms are at the vertices of triangles around each hydroxyl group (after Corno et al.).119 The optimized structure of hexagonal P63 HA is characterized by two formula units per cell and contains 44 atoms per unit cell. As a consequence of lowering symmetry, there are three types of Ca ions inside the cell instead of two (as in the case of P63/m space group), indicated by Ca1, Ca2 and Ca3, in which Ca1 has three oxygen atoms as first neighbours, Ca2 has six, and Ca3 has four. The OH– group is linked to three Ca3 ions, which form an equilateral triangle in the ab plane with the hydroxyl in the centre (Fig. 2.25). Two ions of i In crystallography, a screw axis is a symmetry operation describing how a combination of rotation about an axis and a translation parallel to that axis leaves a crystal unchanged. Introduction: Hydroxyapatite crystals 70 type Ca1 and two of Ca2 are present in the cell and involved in an octahedral arrangement and, as for phosphates groups, there are six of them inside the unit cell, all equivalent for symmetry. Finally, there are two OH– groups per cell, oriented in the same direction along the c-axis and defining internal channels. HA crystal structure planes are differently charged. The basal planei (or c-surface) is calcium-rich and therefore, positively charged; while the prismatic planeii (or a-surface) is phosphateand hydroxyl-rich and hence, negatively charged (Fig. 2.26).120 Owing to this intrinsically different surface reactivity, HA exhibits selective adsorption of various ions, organic compounds and proteins, and can therefore be used in numerous chromatographic applications.121 Fig. 2.26. A part of the crystal structure of the hexagonal HA is shown in (A), wherein PO4 tetrahedra are omitted for clarity. A unit cell is shifted to (0; 0; ¼). (001) plane contains (1 + 4/2) Ca atoms and only 4/4 hydroxyls’ oxygen atoms, hence having ratio Ca:O = 3:1, and therefore discussed surface will be charged positively. In contrast, (100) plane contains equal amounts of Ca and hydroxyl groups, (1 + 2/2) and (4/4 + 2/2), respectively. Additionally, Ca atoms are slightly shifted out of plane to ± 0.06 Å, resulting in the enrichment of (100) plane with hydroxyl and phosphate groups. Therefore this surface will be charged negatively. Schematic representation of HA crystal growing along c-axis is displayed in (B). Such crystal growth leads to the development of the negatively charged prism-faceted a-surface (aand bsurfaces are equivalent). 2.7.2.2. Hydroxyapatite applications Additionally to its wide uses in chromatography, synthetic HA is used nowadays in several other applications, i.e. as fertilizer,122, 123 in catalysis124, or as a way of effectively sequestering heavy metal species in contaminated soils and/or ground water.125, 126 Also, HA is particularly important in medicine, and is widely used as a biomaterial in clinical i It is the plane which is perpendicular to the c-axis of the hexagonal structure. ii It is the plane which is parallel to the c-axis of the hexagonal structure. Introduction: Hydroxyapatite crystals 71 applications such as a replacement material in bone repair and substitution, as a substrate for the development of scaffolds in bone tissue engineering, as a drug delivery system, etc. HA medical applications derive from the fact that it is chemically similar to the inorganic component of bones and teeth. Nevertheless, as a result of functional irregularities, vertebrates can produce HA pathologically (ectopic calcification), resulting in the formation of renal and bile stones, calcifications of cartilages, basal ganglia, etc.127 Hence, synthetic HA crystals can provide a unique tool towards understanding these biomineralization phenomena.128 In addition to it’s peculiar physico-chemical and physiological properties, the material properties of HA such as: particle size, dimensional anisotropy, morphology, real microstructure, etc., are also of critical importance for its application and can be optimized. There are numerous applications in which material properties HA crystals play a significant role. For example, plate-like and fibrous HA particles exhibit enhanced adsorption properties due to the charging surface efficiency.129 Also, it is well known that crystals with plate-like morphologies are the most effective in stiffening isotropic composite materials, followed by those of fibrous morphology, with the least effective geometry being spherical (Section 2.6.1).98 From this we can conclude that HA crystals with the desired morphological features are candidates for the moderated reinforcement of biomaterials used in bone repair and substitution, such as calcium phosphate bone-cements, biocomposites, etc.114, 130 2.7.2.3. Overview of the synthetic routes to hydroxyapatite crystals The first synthesis of apatite was that of Daubreé131 in 1851 who obtained HA by passing phosphorus trichloride vapour over red hot lime.i Since then, interest in HA continued growing and already in 1951 the first review on methods for the preparation of HA was reported.132 Several approaches have been used to synthesize phase-pure, well-defined HA crystals with controlled materials properties. The techniques employed include, hydrothermal routes, molten salts syntheses, precipitation, and, to a lesser extent, growth in a gel system or liquid-solid-solution synthesis.9, 111, 133-137 Depending upon the technique, materials with various morphologies, stoichiometries, sizes and levels of crystallinity can be obtained (Table 2.9). i A naturally occurring mineral predominately composed of carbonates, oxides and hydroxides of calcium. Introduction: Hydroxyapatite crystals 72 Of all the methods employed to synthesize HA crystals, hydrothermal is one of the most promising and convenient due to it’s successful application for the one-pot synthesis of a desired phase under gentle reaction conditions.111 Remarkably, hydrothermal technique is also well known as an efficient approach to synthesized defect-free single crystals of high crystallinity.138 Because the reactions are solution-mediated, particle size and morphology can be controlled by experimental strategies that regulate nucleation, growth and aging processes.138 The powder prepared by hydrothermal technique is highly pure and not agglomerated. Hydrolysis of TCP, DCPA, DCPD or OCP requires low temperatures (usually below 100°C) and results in HA needles or blades in the micron range. However, in most cases, the hydrolysis product is highly non-stoichiometric (Ca/P ratio in the range of 1.50–1.71).111 Molten salts synthesis (a solid state reaction) usually give a stoichiometric and well-crystallized product, but they require relatively high temperatures and long heat-treatment times.111, 134 Moreover, they incorporate K+ ions from the reactants into the HA crystal lattice and show a sensitive dependence on the preparation conditions.134 In the case of precipitation, where the temperature does not exceed 100°C, nanocrystals are obtained. They have shapes of blades, needles, rods, or equiaxed particles. Their crystallinity and Ca/P ratio depend strongly upon the preparation conditions and are in many cases lower than that of well-crystallized stoichiometric HA.9, 135 HA fibres grown in the gel system do not have high mechanical strength. Additionally, they show a strong dependence on the synthesis conditions and their crystallinity and thermal stability are relatively inferior.111, 136 Liquid-solid-solution syntheses are also dependent on the preparation conditions.137 Introduction: Hydroxyapatite crystals 73 Table 2.9. Preparation techniques for hydroxyapatite crystals. Technique Synthetic conditions Comments Homogeneous precipitation using urea (85-95°C) • High degree of crystallinity. • Ca/P ratio ~1.67. • Controlled morphology. • Controlled particle size (from nm to mm). Hydrothermal111 Homogeneous precipitation using chelates (90-200°C) • High degree of crystallinity. • Ca/P ratio ~1.67. • Controlled morphology. • Controlled particle size (from nm to mm). Hydrolysis of calcium orthophosphates111, 133 ≤ 100°C • Needles or blades. • Micro-particles. • Ca/ P ratio ~ 1.50-1.71. Molten salts synthesis134 850-1190°C • Well crystallized product. • Ca/P ratio ~1.67. • Size of several micrometers. • Single crystals with hexagonal morphology. • Incorporation of K+ to the crystal structure. • Requires high temperatures. Precipitation9, 135 ≤ 100°C • Inhomogeneous nanocrystals. • Needles, blades or equiaxed particles. • Variable crystallinity and Ca/P ratio. Gel system111, 136 Room temperature • Polycrystalline crystals. • Sensitive dependence on the preparation condition. • Inferior thermal stability. Liquid-solid-solution synthesis137 Use of templates (90-180°C) • Sensitive dependence on the preparation condition. • Nanocrystals. Introduction: Hydroxyapatite crystals 74 Hence, hydrothermal routes appear to be well suited for the development of an inexpensive synthesis method for crystalline HA, which focuses on precise control of particle size, morphology and chemical composition. For this reason, hydrothermal is nowadays the most commonly employed method to synthesize HA crystals.111, 139 Fibres prepared by hydrothermal method have a more controlled and better morphology, and are more uniform in size, purity and composition compared to HA produced by other methods.138 Many research groups have prepared HA crystals with various morphologies by hydrothermal routes. In the majority approaches reported, the phase-pure HA products are commonly characterized by either rod-like111, 136, 137 or plate-like137, 140, 141 particle morphology. In addition, hydrothermally-produced HA powders have been reported to exhibit hexagonal prism-like,129, 134, 142, 143 elliptical-like,133, 144 or ribbon-like145 particle shapes. Despite this variety of morphological features, all synthesized crystals growth along the c-axis of the hexagonal structure, the natural tendency of HA. 2.7.3. Hydrothermal approaches to hydroxyapatite crystals 2.7.3.1. Hydrothermal technique. Definition. The term hydrothermal is purely of geological origin. It was first used by the British geologist Sir Roderick Murchison (1792–1871) to describe the action of water, at elevated temperature and pressure, in bringing about changes in the earth’s crust leading to the formation of various rocks and minerals.139 The largest single crystal formed in nature (beryl crystal of >1.000 kg) is of hydrothermal origin. Hydrothermal processing can be defined as any homogeneous (nanoparticles) or heterogeneous (bulk materials) reaction in the presence of aqueous solvents or mineralizers under high pressure and temperature conditions to dissolve and recrystallize materials that are relatively insoluble under ordinary conditions.139 Byrappa and Yoshimura define hydrothermal as any homogeneous or heterogeneous chemical reaction in the presence of a solvent (whether or non-aqueous) above the room temperature and at pressure greater than one atmosphere in a closed system.138 However, chemists prefer to use the term solvothermal when these chemical reactions take place in the presence of organic solvents. There are several other terms like glycothermal, alcothermal, ammonothermal, carbonothermal, etc. used sometimes to further specify the employed non-aqueous solvent. Introduction: Hydroxyapatite crystals 75 In hydrothermal, it is possible to work under supercritical fluid i conditions. Supercritical solvents (the most commonly used are H2O or CO2) are now regularly used to carry out a wide range of chemical reactions replacing organic solvents in a number of chemical processes, including nanoparticles fabrication, extraction, chemical manufacturing, waste treatment, recycling, etc.139 2.7.3.2. Hydrothermal method. Historical background111, 138, 139 In 1839, the German chemist R. W. Bunsen contained aqueous solutions in thick-walled glass tubes at temperatures above 200°C and at pressures above 100 bars. The crystals of barium carbonate and strontium carbonate that he grew under these conditions mark the first use of hydrothermal aqueous solvents as media. Another early report of the hydrothermal crystal growth was by E.T. Schafhäult, who in 1845 prepared fine particles of quartz in a Papin’s digester.ii In 1905 G. Spezzia published reports on the growth of macroscopic crystals. By 1900 more than 150 mineral species had been synthesized by hydrothermal methods, including diamond. Commercial application of the hydrothermal technique began in 1908 when K.J. Bayer leached bauxite mineral under hydrothermal conditions to obtain aluminium. This opened up a new avenue for hydrothermal research in the area of metallurgy. This was subsequently followed by the synthesis of various minerals and, during the 1940s, bulk crystal growth, phase equilibria studies, etc. The focus of hydrothermal technology in the earlier days was exclusively on the growth of bulk single crystals, since it was technologically important to produce big, pure and defect-free single crystals. The earliest material to be grown was quartz,iii followed by a wide range of mineral species in the bulk form such as ruby, corundum and various other silicates, carbonates, phosphates, sulphates, etc. Over 5.000 kg of quartz single crystal can be produced in one experimental run in a single large autoclave. So far no other crystal growth technique can match the size and quantity of quartz produced in a single experimental run using the hydrothermal technique. Today much of the commercial production of bulk single crystals using hydrothermal technology is devoted to quartz, gems –i.e. emerald, corundum, ruby or alexandrite– and other i A supercritical fluid is any substance at a temperature and pressure above its thermodynamic critical point. ii The Papin’s digester (also known as a steam or bone digester) is a high-pressure cooker invented by French physicist Denis Papin in 1679. It is a device for extracting fats from bones in a high-pressure steam environment, rendering them brittle enough to be easily ground into bone meal. It is the forerunner of the autoclave and the domestic pressure cooker. iii With a melting point of 1670 °C and high chemical resistance, quartz crystals would be very difficult to form by either melt or flux techniques. Introduction: Hydroxyapatite crystals 76 single crystals with commercial value. Additionally, nowadays hydrothermal method is regularly used for processing a wide range of materials not only as bulk crystals, but particles ranging from fine crystallites to nanoparticles with controlled size and morphology. Also, hydrothermal methods have proved to be extremely efficient both in the search for new compounds with specific physical properties and in the systematic physico-chemical investigation of intricate multicomponent systems at elevated temperatures and pressures. 2.7.3.3. Hydrothermal advantages in comparison with other conventional techniques The aforementioned hydrothermal methods allow the synthesis of defect-free single crystals of high crystallinity and purity with controlled morphology and narrow particle size distribution. The reaction takes places under gentle conditions in one step. Additionally, this method offers several other advantages cf. other conventional techniques:138 i) compounds with elements in oxidation states that are difficult to obtain, can be synthesized in a closed system by hydrothermal methods (i.e., ferromagnetic chromium(IV) oxide). ii) it is useful for so-called low temperature phases, i.e. α-quartz. iii) it is possible to synthesize metastable compounds, such as subiodides of tellurium, Te2I. Hydrothermal method has a range of other advantages such as: a high reaction rate for powders, homogenous dispersion in the liquid phase, almost pollution free, does not require very expensive and highly sophisticated equipment, energy saving processing and, in many cases, it is possible to produce new phases with no natural analogue inaccessible via other methods.138 Also, since hydrothermal fluids offer higher diffusivity, lower viscosity, facilitate mass transport and higher dissolving power, reaction kinetics can be enhanced cf. conventional techniques.139 Table 2.10. Comparison of various powder synthesis methods.139 Parameter Solid state reactions Co-precipitation Sol-gel Hydrothermal Cost Low to moderate Moderate High Moderate Compositional control Poor Good Excellent Good/excellent Morphology control Poor Moderate Moderate Good Powder reactivity Poor Good Good Good Purity (%) <99.5 >99.5 >99.9 >99.5 Annealing step Yes Yes Yes No Milling step Yes Yes Yes No Introduction: Hydroxyapatite crystals 77 2.7.3.4. Principles of the crystal growth under hydrothermal conditions The grow of a single crystal into the seed can proceed in two ways:138 • recrystallization of the solid substance, including its dissolution in the liquid phase, followed by convectivei mass transfer of the dissolved part of the substance to the growth zone or seed. • dissolution of the mixture of nutrient components with the help of their convective mass transport into the growth zone and interaction of the dissolved components of the seed surface. The methodology of the growth of single crystals on a seed is the establishment of growth conditions in which the process is represented by the sum of macroand micro-processes occurring between the interface boundary of the solution and the crystal. The composition and concentration of the solution, temperature and pressure, hydrodynamic conditions and surface contact of the phases are some of the basic physical and chemical parameters which determine the regime and rate of the dissolution of the nutrient, mass transport and the possibility of the formation of new phases.138 Typically, hydrothermal syntheses mainly result in the formation of the thermodynamically favourable product under certain conditions. However, kinetics also plays an important role. As an example, in the case of ZnO the presence of NH4+ ions in the solution produces a marked increase in the growth rate of the prismatic faces, which under ordinary conditions, does not exceed hundredths of a millimetre per day.138 2.7.3.5. Hydrothermal synthesis of hydroxyapatite crystals Hydrothermal preparation techniques of HA crystals can be divided into two main groups (Table 2.9): i) Methods using urea as precipitation agent146-149 ii) Methods using chelating agents150-152 Both techniques involve a homogeneous precipitation of the HA crystals with a slow reaction rate, a relatively easy procedure for obtaining uniform HA particles. The first method is based on the thermally-induced decomposition of urea (NH2)2CO. Essentially, at the temperature range 80-95°C, (NH2)2CO smoothly decomposes to form carbon dioxide and aqueous ammonia species according to the Eq. 2.6 and 2.7: i Convective transport occurs when a constituent of the fluid (mass, energy, a component in a mixture) is carried along with the fluid. The flux is determined by its concentration and the fluid velocity. Materials and methods: Characterization methods 84 rod, divided into several pieces and allowed to set for 1 hour. Each piece was separately incubated in distilled water at 37°C for a specific period of time. Immediately prior to the XRD analysis, the CPC test sample was very quickly ground in warm acetone using a hot agate mortar and pestle (80°C), in order to remove the rest of the water and, to some extent, suspend the reaction process. Diffraction patterns were collected in the 2 θ range of 8º to 38º, with a scan speed of 4º/min. All rest of the XRD experimental parameters were maintained the same as in the case of the XRD analysis described above. The conversion of CPC precursors to HA was roughly estimated by employing a technique reported by Fukase et al.156 and improved by Ishikawa and co-workers.88 Simply put, this approach is based on the fact that the intensity diffracted by a crystalline phase is essentially proportional to the quantity of the diffracting compound. Although this technique is found to be a reasonable and valid method for the rough estimation of CPC conversion, it is not very accurate owing to the line broadening effect and low intensities distinctive for apatite CPC systems. The extents of the cement setting reactions (%conv) were calculated according to Eq. 3.1:88 100 2 2 )(1)/(1 % )0()()0()( )( )( × − − + − + =∞ TTCPtTTCPDCPDtDCPD HA tHA conv IIII I I (3.1) where IHA(t), IDCPD(t) and ITTCP(t) are intensity values of the (002) diffraction peak of HA (JCPDS No. 72-1243), the sum of the intensities of the (020) and (021) diffraction peaks of DCPD (JCPDS No. 9-77) and the sum of the intensities of the (013) and (040) diffraction peaks of TTCP (JCPDS No. 70-1379), respectively, at specific times of setting reaction t = 0.5, 1, 2, 3, 4, 5, 6, 7, or 8 hours. IDCPD(0) and ITTCP(0) are intensity values of the same characteristic diffraction peaks from DCPD and TTCP, respectively, but for the unreacted CPC specimen at the starting time (t = 0). Since the dissolution-precipitation reaction between the fastest setting combination of CPC precursors resulted in phase-pure HA already after 6 hours and no significant product advantages were observed from longer reaction times, the value IHA( ∞ ) –peak intensity of HA in the fully converted CPC– can be consistently chosen as the intensity of the (002) diffraction peak of HA at t = 8 hours. In a typical data processing step, firstly the background was removed from each XRD spectrum using the PeakFit software program and then the resulting I data (three replicates) were averaged. All used diffraction peaks were free of overlapping. Materials and methods: Characterization methods 85 3.2.2. Spectroscopic product analyses 3.2.2.1 Energy-disperse X-ray spectroscopy Energy-dispersive X-ray spectroscopy investigations (EDX) for semiquantitative Ca and P content determination were performed with EDAX DX-95 spectrometry system using commercial hydroxyapatite powder (HW-003, Mitsubishi Chemicals) as a standard sample (molar ratio Ca/P = 1.67). 3.2.2.2 Diffuse reflectance infrared Fourier-transform spectroscopy The room temperature diffuse reflectance infrared (IR) Fourier-transform spectra were recorded on a Jeol JIR-7000 spectrometer with a resolution of 4 cm−1 and a number of scans 100. The crystalline products (~4 mg) were thoroughly ground with potassium bromide powder (~200 mg) in an agate mortar and pestle to give a fine mixture and subjected to IR analysis. For the background spectrum, a finely ground KBr powder was used. 3.2.2.3 Raman scattering The room temperature Raman scattering measurements were carried out on a confocal Raman microscope WITEC CRM200 with linear polarized laser light (λ = 532 nm) as the excitation source. The laser beam was focussed on the sample surfaces with the aid of an optical microscope unit equipped with a Nikon 10x objective and numerical aperture of 0.25. 3.2.3. Electron microscopy product analyses 3.2.3.1 Scanning electron microscopy The morphology was studied by scanning electron microscopy (SEM) and field-emission environmental scanning electron microscopy (FE-ESEM) using Hitachi S-4500 and FEI Quanta 600F microscopes operating at 15 and 5 kV, respectively. 3.2.3.2 Transmission electron microscopy The fine microstructure and phase purity of the samples were investigated by transmission electron microscopy (TEM), electron diffraction (ED) and high-resolution TEM (HRTEM) using Hitachi H-8100 and Jeol 4000EX microscopes operating at 200 and 400 kV, respectively. The samples for TEM were crushed or dispersed in methanol and deposited on a carbon mesh grid. Computer-simulated HRTEM images for different defocus and different thickness were obtained using the Mac Tempas and Crystal Kit software programs. Materials and methods: Characterization methods 86 3.2.4. Particle size distribution estimation Average particle sizes and particle size distributions (PSD) of the raw DCPD and TTCP powders were determined using sedimentation analysis on a Micrometrics SediGraph 5100 particle size analyser. For the measurement, 2 g of the test sample was dispersed by vigorous stirring for 2 hours in 50 mL of 0.075% aqueous solution of sodium hexametaphosphate (dispersant). Particle size characteristics of the ball-milled DCPD and TTCP samples were monitored using dynamic light-scattering (DLS) technique on a Malvern Instruments Zetasizer Nano-ZS (λ = 633 nm, 3 runs, run duration of 10 seconds). For DLS measurements, a small amount of the powder was fully dispersed in water using high-power ultrasonication and subjected to the DLS analysis. 3.2.5. Thermo-gravimetric and differential thermal analyses Thermo-gravimetric and differential thermal analyses (TGA/DTA) were performed using a Mac Science Type-2020 TG-DTA thermoanalyser. The heating profile was a 10°C/min linear ramp from room temperature to desired temperature usually 1300°C. Measurements were realized in a static air atmosphere using platinum crucibles and alumina powder as reference material. 3.2.6. Elemental analysis Carbon content was determined quantitatively by combustion bulk elemental analysis using a Fisons EA 1108 elemental analyzer. In this technique, a sample is burned in an excess of oxygen, and the weight of traps collect carbon dioxide is used to calculate carbon content. 3.2.7. Porosity evaluation 3.2.7.1 Adsorption-desorption of nitrogen Nitrogen adsorption-desorption measurements were performed at liquid nitrogen temperature on a Quantachrome Autosorb-1 automated gas sorption system. Prior to measurement, the samples were outgassed in vacuum at 150°C for 2 hours. Total surface area was calculated according to the multipoint Brunauer-Emmett-Teller method (BET). Pore volumes and size distributions were determined by the Barrett-Joyner-Halenda method (BJH). Materials and methods: Characterization methods 87 3.2.7.2 Mercury porosimetry Porosity was also evaluated by mercury porosimetry in a Micromeritics AutoPore II 9220 porosimeter. Prior to analysis, the samples were degassed at 0.001 psia to remove physisorbed gases. The penetrometer was filled with mercury and the low pressure analysis started at 1.5 psia and terminated at 15 psia. The penetrometer was then transferred to the high pressure chamber and the analysis with mercury intrusion (15-60000 psia) followed by mercury extrusion (60000-15 psia) was conducted. 3.2.8. pH measurements The pH measurements were carried out using a Mettler Toledo InLab®413 SG pH-meter with combination polymer electrode. Prior to each measurement, the pH-meter was calibrated with the standard buffer solutions pH = 6.86 and pH = 4.01 (Wako). For the CPCs samples, 1 g of the CPC powder was mixed with 20 mL of distilled water under vigorous stirring for 5 seconds, and then the pH of the suspension was monitored every 30 seconds with continuous stirring. 3.2.9. Final setting time determination Final setting time (ST) of the as-prepared cements was determined according to the American Dental Association specification No. 9. For this purpose, equimolar amounts of ball-milled DCPD and TTCP powders (total 0.6 g) were thoroughly mixed in an agate mortar and pestle and placed inside a latex finger cot. After this, an appropriate amount of distilled water (0.3 mL) was added to the mixed powders using a micropipette. The used solid to liquid (S/L) ratio was 2 g/mL. This mixture was hand kneaded for 1 min and the resultant putty-like material was placed inside a polytetraflouroethylene (PTFE) mold of 6 mm diameter and 3 mm deep. The Vicat needle technique (Fig. 3.1) was used to determinate the final setting time (Section 2.5.2). ST was considered to be reached when the tested CPC specimen bore the weight of the Vicat needle (weight = 453.6 ± 0.5 g, tip diameter = 1.06 ± 0.05 mm) without appreciable indentation on the surface. All values are the average of at least six replicates. Materials and methods: Characterization methods 88 Fig. 3.1. Vicat needle tool used for the estimation of the ST of the CPCs. 3.2.10. In vitro test In vitro test was performed in the optimal CPC in order to investigate its bioactivity. Bioactivity studies were carried out using an acellular simulated body fluid (SBF) solution reported by Kokubo et al.154 SBF is almost identical to human blood plasma in terms of its ion concentrations, but it does not contain organic substances such as proteins (Table 3.1). Table 3.1. Ions concentration at human blood plasma and employed SBF. Ions concentration, mM Na+ K+ Mg2+ Ca2+ Cl− HCO32− PO42− Human blood plasma157 142.0 5.0 1.5 2.5 103.0 27.0 1.0 SBF 142.0 5.0 1.5 2.5 148.8 4.2 1.0 Nevertheless, the soaking of bioactive materials in SBF can reproduce the apatite formation seen on such materials in the living body. Therefore, it is a useful tool for evaluating the bone-bonding bioactivity of materials by simply examining the ability of apatite to form on the surface of these materials (Section 2.3.3.3).158 The specimens, in the form of pellets, were prepared, as for the ST determination, by mixing equimolar amounts of the ball-milled DCPD and TTCP powders (total 0.2 g) with 0.067 mL of distilled water (S/L ratio: 3 g/mL). The mixture was kneaded by hand in a latex finger cot for 1 minute and the resulting putty-like material was placed into a cylindrical mold assembly of 6.9 mm diameter. The sample was pressed under 5.3 MPa (20 Kg/cm2) for 1 minute using a pressure-loading device. The specimen was removed from the mold and kept Materials and methods: Characterization methods 89 for 1 hour at ambient conditions to set. The test samples were then placed into a polystyrene bottle (V = 5.7 mL) previously washed with 1M HCl and water, filled with 3 mL of SBF and incubated at 37°C. The polystyrene bottle and SBF solution were refreshed every 3 days. After soaking for 3, 7 and 10 days, the CPC specimens were collected, their surfaces were washed with acetone and finally dried in ambient conditions before further analysis. The in vitro tested samples were studied by XRD, IR, TGA/DTA and SEM characterization techniques in order to investigate the formation of hydroxyapatite on their surfaces. 3.2.11. Mechanical properties measurements 3.2.11.1 Compressive strength testing For the compressive strength (CS) measurements, the CPC specimens were prepared in a similar manner as in the case of in vitro test, the difference being that three times more TTCP and DCPC precursors were used (S/L ratio: 3 g/mL), resulting in the specimen with an aspect ratio of 2. It has to be noted here that the applied pressure may not be applicable for clinical practice, but in the current study it was selected with the aim of obtaining reproducible results in a well-defined and systematic way for comparison purposes. Biocomposite pellets were prepared in a similar way. In a typical procedure, equimolar amounts of ball-milled DCPD and TTCP powders were thoroughly ground together using an agate mortar and pestle. Next, the desired amount of HA crystals with various morphology were well-dispersed using a hand spatula in the CPC precursors mixture, followed by gentle mixing in an agate mortar. The resulting fine powder (total 0.6 g) was placed inside a latex finger cot and an appropriate amount of distilled water was added to the composite CPC precursor using a micropipette. This amount of liquid phase was calculated to maintain a solid to liquid (S/L) ratio of 2, on the basis of the DCPD and TTCP cement precursors, with the reinforced phase being not considered. The mixture was hand kneaded for 1 minute and then placed into a cylindrical mold assembly of 6.9 mm diameter. The batch was pressed under 5.3 MPa (20 Kg/cm2) for 1 min using a pressure-loading device. After pressing, the biocomposite specimen was removed from the mold and kept for 1 h at ambient conditions to set before being incubated in acellular simulated body fluid solution at human body temperature of 37°C for 23 hours. Furthermore, the samples were placed into a polystyrene bottle (V = 5.7 mL) filled with 3 mL of simulated body fluid solution (SBF) and incubated at 37°C for 23 hours. Finally, the wet specimens were subjected to compressive strength measurements on a Universal Testing Machine (Shimadzu Autograph AG-I) with a crosshead speed of 0.5 mm/min until Materials and methods: Characterization methods 90 fracture. The compressive strength value was determined as the critical force divided by the loaded area. All values are the average of at least six replicates. 3.2.11.2. Nano-indentation measurements The hardness (H) and elastic modulus (Es) of the optimal CPC specimen were probed at the sub-micro/micro level by means of the nano-indentation technique. Es is a measure of the inherent material stiffness –material resistance against suffer permanent deformations (elastic) by an applied force, while H is related with the material resistance to various kinds of shape change when a force is applied, and thus, related to plastic deformation. Experiments were performed on a Hysitron TriboScan UBI-1 nanohardness tester in conjunction with a Digital Instruments Nanoscope III atomic force microscope (AFM). A Berkovich diamond indenter tip with regular triangle pyramid geometry and nominal radius of 300 nm was used. The system calibration and data evaluation were performed using the method of Oliver and Pharr.159 Thermal drift is corrected in the Hysitron TriboScan software by automatically measuring the drift for up to 20 seconds before starting each indent. The reduced indentation modulus (Er) and the hardness were calculated from the load-displacement curve according to the Eq. 3.2 and 3.3, respectively: )( 2C rhA S E π = (3.2) )( max C hA P H= (3.3) where S is the slope of the fitted unloading curve at the point of initial unloading (20%-95% of the unloading curve was used for fitting); Pmax is the peak load and A(hC) is the projected contact area at the peak load –i.e. area of the indentation at the contact depth of the residual indentation hC (Fig. 3.2). Materials and methods: Characterization methods 91 Fig. 3.2. Schematic illustration of an indentation load vs. displacement plot showing quantities used to determine the reduced indentation modulus and hardness from the obtained curve. The reduced modulus (Er) is related to the elastic modulus (Es) of the test specimen through the Eq. 3.4: s s i i rEEE )1()1( 122 νν − + − = (3.4) where νi is the Poisson’s ratio for the diamond tip (0.07),160 νs is the Poisson’s ratio for synthetic HA (0.28)161 and Ei is the diamond elastic modulus (1141 GPa).160 Two loading functions were employed: (i) Peak load of 1 mN, loading rate of 0.2 mN/s, holding time 60 seconds, unloading to 0.05 mN with rate of 0.2 mN/s, holding 20 seconds at 0.05 mN and unloading back to 0 mN with rate of 0.5 mN/s; (ii) Peak load of 5 mN, loading rate of 1 mN/s, holding time 60 seconds, unloading to 1 mN with rate of 0.4 mN/s, holding at 1 mN for 20 seconds and unloading back to 0 mN with rate of 0.5 mN/s. The first loading function was used to probe Es and H in HA single crystals with hexagonal prism-like particle morphology (Chapter V). The second one was employed for nano-indentation testing of CPCs (Chapter IV) and biocomposites (Chapter VI) systems. One hundred fifty indents were randomly performed in order to obtain average values for CPC’ and composites’ Es and H. For the nano-indentation test on the optimal CPC and composites (optimal CPC reinforced by 10 wt.-% and 15 wt.-% of fine-plate and hexagonal HA crystals respectively), a pellet was prepared in a similar manner as for the CS investigation. Then, the specimen was Materials and methods: Characterization methods 92 embedded in a solution formed by 77.0 wt.-% of methylmethacrylate and the polymerization initiators –1.4 wt.-% of benzoylperoxid and 21.6 wt.-% of nonylphenylpolyethylenglycol acetate. In order to fill up the pores in the CPC pellet, a slow polymerization of methylmethacrylate to polymethylmethacrylate at a moderated low temperature profiles [(39°C, 12 h) → (48°C, 12 h) → (55°C, 12 h)] was performed. Then, the section for nano-indentation was cut in parallel to the pellet base, polished by a set of abrasive papers of decreasing size of grit followed by polishing with 1 µm diamond powder suspended in water and subjected to nano-indentation test. IV PREPARATION OF CALCIUM PHOSPHATE BONE-CEMENTS 4.1.- EXPERIMENTAL PART...........................................................................................95 4.1.1. Synthesis of dicalcium phosphate dihydrate......................................... 95 4.1.2. Synthesis of tetracalcium phosphate.....................................................95 4.1.3. Mechanochemical grinding................................................................... 96 4.1.4. Preparation of the bone cements ...........................................................98 4.2.- RESULTS ................................................................................................................... 99 4.2.1. DCPD and TTCP products analysis...................................................... 99 4.2.2. Bone cements synthesis and CPCs comparison: selection of the optimal cement.............................................................................................. 104 4.2.3. Phase composition evolution during setting reaction.......................... 106 4.2.4. Spectroscopic CPC product analysis...................................................108 4.2.5. Thermal stability of CPC product .......................................................111 4.2.6. Textural properties of CPC .................................................................112 4.2.7. CPC product morphology and microstructure .................................... 113 4.2.8. In vitro bioactivity test of CPC product .............................................. 116 4.2.9. Mechanical properties of the optimal CPC .........................................118 4.3.- DISCUSSION...........................................................................................................119 4.4.- CONCLUDING REMARKS....................................................................................127 Preparation of calcium phosphate bone-cements: Results 100 Fig. 4.3. XRD patterns for DCPD powder prepared by precipitation reaction (A) following by mechanochemical grinding using balls with diameter of 3.6 mm (B). Tick marks below the patterns correspond to the positions of the Bragg reflections expected for the monoclinic DCPD (CaHPO4·2H2O, JCPDS No.9-77). Fig. 4.4. XRD patterns for TTCP powder prepared by solid-state reaction (A) following by mechanochemical grinding using balls with diameter of 4.8 mm (B). Tick marks below the patterns correspond to the positions of the Bragg reflections expected for the monoclinic TTCP (Ca4(PO4)2O, JCPDS No. 70-1379). In order to decrease the crystal sizes of as-prepared D-raw and T-raw, the powders were subjected to a wet rolling ball-mill process (Tables 4.1, 4.2). IR analysis of the ball-milled DCPD and TTCP products (see Annexe II, Fig. AII.1) confirm the absence of residual acetone and isopropanol from the ball-milling procedures. According to the powder XRD, mechanochemical grinding does not lead to changes in the phase composition and all as-prepared samples were found to be phase-pure TTCP and DCPD (Table 4.2). As an example, we show phase analysis results for D-36 and T-48 products (Fig. 4.3B, 4.4B). XRD reveals that the milling of D-raw and T-raw using balls with diameters of 3.6 and 4.8 mm results in the pure dicalcium phosphate dihydrate and tetracalcium phosphate (samples D-36 and T-48), respectively. It should be noted that for the ball-milled TTCP products, a broadening of the XRD peaks is observed. This is likely due to the comminution of agglomerates, reduction of crystallites sizes and crystal structure deformation from the ball-milling procedure, overall resulting in partial amorphization of the ball-milled material,86 in good agreement with reported data for TTCP.50 Preparation of calcium phosphate bone-cements: Results 101 Fig. 4.5. SEM and PSD characterizations of CPC precursors illustrating the considerable reduction of the primary particle’s and aggregate’s sizes after ball-milling. (A) – the overall morphology of the DCPD powder (left panel) synthesized by a precipitation reaction together with the PSD, determined by sedimentation analysis (right panel); (B) – morphology of the same sample (left panel, inset: high magnification SEM) after mechanochemical grinding using balls with diameter of 3.6 mm in conjunction with the PSD, determined by DLS technique (right panel); (C) – the overall morphology of the TTCP powder (left panel) prepared by solid-state reaction together with the PSD, determined by sedimentation analysis (right panel); (D) – morphology of the same sample (left panel, inset: high magnification SEM) after mechanochemical grinding using balls with diameter of 4.8 mm in conjunction with the PSD, determined by DLS technique (right panel). Preparation of calcium phosphate bone-cements: Results 102 The SEM morphologies of polycrystalline raw DCPD and TTCP are shown in Fig. 4.5A and 4.5C (left panels), respectively. The sample D-raw consists of relatively large aggregates formed by faceted particles with plate-like morphology, while the T-raw product exhibits a less well-defined microstructure. The morphologies of the raw samples, as determined by SEM, are micrometer-sized particles with typical sizes in the range of several micrometers to several tens of micrometers. Particle size distributions (PSD) of the D-raw and T-raw samples were estimated using sedimentation analysis (Fig. 4.5A, 4.5C (right panels)) and obtained values are in good agreement with the SEM observations. According to the PSD analysis, the as-prepared DCPD exhibits a non-uniform size distribution with particles sizes ranging from ~0.5 μm to ~100 μm. The average particle size for D-raw was established to be ~6 μm (Table 4.2). The latter value most likely corresponds to the overall sizes of the particles in D-raw powder, while the presence of grains with size up to ~100 μm appears to be attributed to the overall sizes of the particle agglomerates. According to the PSD analysis, sample T-raw is also characterized by a non-uniform size, with particles sizes ranging from ~0.6 μm to ~100 μm, which is quite typical for the powder prepared by solid-state reaction followed by thorough grinding in an agate mortar and pestle. The average particle size for T-raw was determined to be ~7 μm (Table 4.2). A wet rolling mechanochemical grinding of the raw DCPD and TTCP powders with three different types of balls (Table 4.1) leads to significant changes of the microstructure, particles size and PSD (Table 4.2). As an example, we show SEM results in conjunction with the PSD data, determined by dynamic light-scattering technique, for D-36 and T-48 samples (Fig. 4.5B, 4.5D, respectively). In sharp contrast to the morphologies of the raw materials, even at low magnification, the ball-milled D-36 and T-48 products show drastic reduction in particle sizes and narrowing of the PSD. Higher magnifications of these powders reveal almost non-aggregated sub-micrometer-sized particles with typical sizes in the range of several hundred nanometers to several micrometers (Fig. 4.5B, 4.5D (left panels, insets)). It can be expected that the use of three different ball diameters (3.6, 4.8 and 9.6 mm) in the course of a wet rolling mechanochemical grinding would provide a useful variable to tailor the grain sizes. To investigate the PSD evolution with respect to the variation of the grinding media, DLS analysis was applied and a summary of the results is presented in Fig. 4.6 and Table 4.2. Preparation of calcium phosphate bone-cements: Results 103 Fig. 4.6. Comparison of the particle size distributions data, determined by DLS technique, for ball-milled DCPC (A) and TTCP (B) cement precursors. According to the PSD analysis, samples D-96 and T-96 posses quite broad peaks centred approximately at 3.66 and 2.96 μm with full width at half maximum (FWHM) of 1.72 and 1.11 μm, respectively; samples D-48 and T-48 have a rather narrow peaks centred approximately at 1.69 and 1.52 μm with FWHM of 0.52 and 0.61 μm, respectively; while the samples D-36 and T-36 exhibit quite narrow peaks at about 1.13 and 1.33 μm with FWHM of 0.41 and 0.51 μm, respectively (Fig. 4.6, Table 4.2). It should be noted that ball-milled D-36 and T-48 particles do not posses any shape anisotropy (Fig. 4.7) and, consequently, the spherical approximation employed by DLS technique to estimate the particles’ diameters, can be considered to be a reliable estimation. Fig. 4.7. Low-magnification TEM images of the DCPD (A) and TTCP (B) powders after mechanochemical grinding using balls with diameters of 3.6 mm and 4.8 mm, respectively. Preparation of calcium phosphate bone-cements: Results 104 4.2.2. Bone cements synthesis and CPCs comparison: selection of the optimal cement Different CPCs were synthesized from ball-milled dicalcium phosphate dihydrate and tetracalcium phosphate via a dissolution-precipitation reaction (Eq. 4.3). Table 4.3 compares the final setting time, phase composition and conversion data for the CPCs produced by a combinatorial method utilizing DCPD and TTCP of various crystal sizes as precursors. Table 4.3. Final setting time, phase composition and conversion date for CPCs produced by combinatorial method utilizing DCPD and TTCP precursors of various particle sizes. The used solid to liquid ratio was 2 g/mL. CPC precursors combination Sample DCPD TTCP Final setting time (min) a XRD phase composition b CPC conversion to HA (%) c C-D96/T96 D-96 T-96 40 ± 4 HA ⎯ C-D96/T48 D-96 T-48 53 ± 4 HA ⎯ C-D96/T36 D-96 T-36 25 ± 2 HA 16.4 ± 1.6 C-D48/T96 D-48 T-96 47 ± 3 HA ⎯ C-D48/T48 D-48 T-48 35 ± 3 HA 11.8 ± 1.1 C-D48/T36 D-48 T-36 31 ± 2 HA 10.9 ± 1.2 C-D36/T96 D-36 T-96 42 ± 3 HA ⎯ C-D36/T48 D-36 T-48 22 ± 2 HA 33.3 ± 3.1 C-D36/T36 D-36 T-36 37 ± 3 HA 12.9 ± 1.4 a– Final setting time data estimated by Vicat needle technique (Fig. 3.1). b– XRD analysis results after 24 hours of setting reaction. Reference: JCPDS card number 72-1243. c– Conversion data after 3 hours of setting reaction calculated by Eq. 3.1. The powder XRD patterns of nine derived cements after 24 hours of setting reaction (1 hour setting at ambient conditions and 23 hours of incubation in distilled water at 37°C) exhibit low intensities and very broad peaks (Fig. 4.8). This observation would suggest that these samples are essentially amorphous, however, as established by TEM investigation, these samples are crystalline (vide infra). Hence, aforementioned features of the powder X-ray diffraction patterns are most likely caused by the nanocrystalline nature and defective HA structure of the as-synthesized cements. The XRD patterns of all samples are very similar to each other and most of the relatively sharp peaks closely resemble hexagonal hydroxyapatite (JCPDS No. 72-1243). According to the powder XRD, the peaks related to the presence of separate DCPD and/or TTCP phases are not detected, indicating that all nine as-prepared cements are totally converted to phase-pure HA already after 24 hours of setting reaction (Fig. 4.8, Table 4.3). Preparation of calcium phosphate bone-cements: Results 105 Fig. 4.8. Comparison of the nine CPC products after 24 hours of setting reaction by powder XRD. The samples were combinatorially prepared from stoichiometric amounts of DCPD and TTCP precursors of various crystal sizes (Table 4.2). The as-produced CPCs were allowed to set for 1 hour at ambient conditions and then incubated in distilled water for 23 hours at 37°C. Tick marks below the patterns correspond to the positions of the Bragg reflections expected for the hexagonal HA phase (JCPDS No. 72-1243). Note the similarity between powder XRD patterns of all CPCs as well as their entire conversion to HA after 24 hours of setting reaction. Fig. 4.9. Comparison of the conversion data after 3 hours of setting reaction (columns) and setting time (dots) for the selected cements. Solid line connects setting time data for visual clarity. The conversion percentages of CPCs to HA phase were calculated according to Eq. 3.1. Note the higher conversion rate and shortest setting time for the C-D36/T48 CPC. Owing to the lack of structural differences amongst all combinatorially-derived CPC samples after 24 hours of setting reaction, setting time measurements were performed. The ST data are compared in Table 4.3. The CPCs show quite varied values of ST, ranging from ~22 Preparation of calcium phosphate bone-cements: Results 106 to 53 minutes for the samples C-D36/T48 and C-D96/T48, respectively. But only several of the cements, namely C-D96/T36, C-D48/T48, C-D48/T36, C-D36/T48 and C-D36/T36, exhibit setting time values which are suitable with respect to surgeons’ requirements (Section 2.5.2).70 Additionally, only these CPCs result in the formation of the required easy to handle putty-like materials when mixed with water (S/L ratio: 2 g/mL) (Fig. 4.2). For further comparison of these CPCs, their conversion rate to HA at a moderated reaction time of 3 hours (1 hour setting at ambient conditions and 2 hours incubation in distilled water at 37°C) were investigated. Fig. 4.9 shows the conversion percentages of the selected CPCs to HA. The calculated values are also enlisted in Table 4.3. Analysis of the obtained data reveals that the cement C-D36/T48 possesses the highest percentage conversion to HA after 3 hours (~33 %). Taking into account the quickest achievement of setting and fastest conversion rate, C-D36/T48 cement has been chosen as the optimal one amongst all the combinatorially prepared CPCs. Consequently, in the research described below, C-D36/T48 cement is used to realize further investigations by different characterization methods. 4.2.3. Phase composition evolution during setting reaction C-D36/T48 cement phase composition evolution during the setting reaction at the selected times of 0, 0.5, 1, 2, 3, 4, 5, 6, 7 and 8 hours is shown in Fig. 4.10A. According to XRD, this cement at t = 0 (unreacted D-36 and T-48 powders mixture) consists of, as expected, DCPD and TTCP (JCPDS No. 9-77 and 70-1379, respectively). Further phase analysis reveals total consumption of D-36 already at t = 4 hours, while T-48 is fully exhausted at t = 6 hours. At this time of reaction (t = 6 hours), C-D36/T48 cement product consists of only phase-pure HA (JCPDS No. 72-1243). The powder XRD patterns of this cement at t = 6, 7, 8 and 24 hours as well as after 1 and 3 weeks of incubation are very much similar and no distinction in phase composition, patterns shapes and positions/intensities of Bragg reflections of an apatite phase are detected. This indicates that the conversion rate of C-D36/T48 CPC to HA appears to be very high and the dissolution-precipitation reaction is fully complete already after 6 hours of cement setting. Preparation of calcium phosphate bone-cements: Results 107 Fig. 4.10. Comparison of the X-ray powder diffraction patterns at different times of setting reaction (A) together with estimated conversion data (B) for the optimal C-D36/T48 cement sample. The dotted lines in (A) correspond to the diffraction peaks of DCPD, TTCP and HA used for the calculation of the conversion percentages of this CPC to HA phase by Eq. 3.1, presented in (B). Note the disappearance of DCPD and TTCP diffraction peaks already after 4 and 6 hours of setting reaction, respectively. Fig. 4.10B shows the C-D36/T48 to HA conversion data determined according to Eq. 3.1, considering the disappearance of the strongest DCPD and TTCP diffraction peaks – (020), (021) and (013), (040), respectively; and the appearance of the HA (002) diffraction peak. This curve clearly shows that the CPC setting reaction does not follow a linear progression and it exhibits four distinct stages during the dissolution-precipitation reaction. The first (I) stage (up to half hour) is most likely attributed to the wetting and initial dissolution of the D-36 and T-48 cement precursors.91 The following stages resemble a sigmoidal curve, which is typically assumed to be a result of the rate-limiting process of the formation and growth of nuclei.168 Hence, the second (II) stage between approximately a half Preparation of calcium phosphate bone-cements: Results 108 and three hours is probably owing to the induction period of the reaction. At this time, no change in the extent of reaction is occurs. As the reaction proceeds, it shifts towards a proliferation period, caused by the rapid increasing of the reaction extent per unit of time (stage III, ~3–6 hours). Finally, the fourth (IV) stage (tapering off period, ~6–8 hours) indicates that the rate of transformation decreases appreciably until it finally reaches the total conversion.168 Overall, it is apparent that after the wetting and initial dissolution of the D-36 and T-48 precursor powders (stage I), the processes involving nucleation and nucleus growth takes place, resulting in a sigmoid-like plot of conversion rate versus time (stages II to IV, Fig. 4.10B). 4.2.4. Spectroscopic CPC product analysis 4.2.4.1. Energy-dispersive X-ray spectroscopy The semiquantitative Ca and P content of the 24 hour-derived optimal cement C-D36/T48 was determined by EDX. The average atomic percent of the calcium and phosphorous elements are 34.3% and 65.7%, providing the average molar ratio of the calcium to phosphorous elements (Ca/P) ∼ 1.49. The molar ratio of calcium and phosphorous elements in stoichiometric HA is 1.67, indicating that C-D36/T48 cement is Ca-deficient HA. Ca-deficient HA can be expressed as Ca10-x(HPO4)x(PO4)6-x(OH)2-x, with 0 ≤ x ≤ 2, where the Ca/P molar ratio varies between 1.67 (x = 0, stoichiometric HA) and 1.33 (x = 2, Ca-deficient HA).169 Hence, the EDX data reveals a rough composition Ca8.94(HPO4)1.06(PO4)4.94(OH)0.94 for the 24 hour-derived C-D36/T48 cement product. 4.2.4.2. Diffuse reflectance infrared Fourier-transformed spectroscopy The chemical and structural composition of C-D36/T48 CPC was studied by IR, a useful technique for determining the incorporation of anions like CO32− and/or HPO42− into the structure of the as-prepared cement, since HA possesses a strong ability to incorporate different ions into its crystal lattice.9 Fig. 4.11 displays a representative spectrum for the 24 hour-derived C-D36/T48 cement product. Peaks centred near 3568 and 634 cm−1 are assigned to the stretching (νs) and librational modes (νL) of the hydroxyl groups respectively. Their harmonic overtones and/or combination bands are also detected135 at ~2085 and 1996 cm−1. Peaks observed at ~2358 and 2339 cm−1 correspond170 to the atmospheric CO2(g). The set of characteristic bands representing apatitic PO43− groups appear at ~1095, 1070 (shoulder) and 1043 cm−1 (triply Preparation of calcium phosphate bone-cements: Results 109 degenerated asymmetric stretching mode of the P-O bond, ν3); ~962 cm−1 (symmetric stretching mode, ν1); ~604, 577 (shoulder) and 567 cm−1 (triply degenerated bending mode of the O-P-O bond, ν4); and ~472 cm−1 (double degenerated symmetric bending mode of O-P-O, ν2).135 The broad peak at ~3406 cm−1 reflects the physisorbed water, while the appearance of the H−O−H deformation band170 at ~643 cm−1 suggests the existence of free water molecules trapped in the crystal lattice of the C-D36/T48 HA cement product (Fig. 4.11). The latter observation is consistent with the non-stoichiometry of this cement seen by EDX, since water molecules can be incorporated only into the crystal structure of nonstoichiometric HA.171 There are also bands that confirm the incorporation of CO32− anions into the structure of hydroxyapatite, which indicate that the obtained HA cement is not only Ca-deficient, but carbonated HA. This fact is proved well by the representative combustion bulk elemental analysis, which shows that C-D36/T48 cement contained approximately 0.55 wt.-% of carbon after 24 hours of setting reaction. These carbonate anions are believe to come from the synthesis procedures, namely, from the atmospheric/water CO2 and/or from tiny impurity of the CaCO3 phase (if any) utilized for the DCPD and TTCP synthesises. According to the assignation of the bands owing to CO32− groups (1462, 1419 and 875 cm−1), CD36/T48 can be referred as B-type carbonated HA, i.e., with the CO32− ions occupying PO43− positions in the HA structure.172 Fig. 4.11. IR spectra collected from the 24 hour-derived C-D36/T48 cement product. It should be emphasized that the band at ~875 cm−1 can also be assigned to the symmetric stretching mode of P-O(H) bond in the HPO42− groups present in Ca-deficient HA.135 This characteristic band of HPO42− anions normally appears near 870 cm−1, where it is Preparation of calcium phosphate bone-cements: Results 116 direct measurement of the d100 spacing on the HRTEM image (Fig.4.16B), gives a slightly smaller distance (~7.4 Å) than a similar measurement for the ideal hexagonal HA structure (~8.2 Å). The calculated ED pattern for hexagonal HA does not fit perfectly with the FT pattern (given as inset in Fig.4.16B). These experimental data suggest that C-D36/T48 cement does not crystallize in the ideal HA structure; the observed discrepancy can be explained by the non-stoichiometry of the as-synthesized C-D36/T48 CPC Ca8.94(HPO4)1.06(PO4)4.94(OH)0.94, as determined by EDX analysis. 4.2.8. In vitro bioactivity test of CPC product To the first approximation, the bioactivity of the optimal C-D36/T48 cement was evaluated by means of the in vitro test using soaking in an acellular simulated body fluid solution at the human body temperature of 37°C. Fig. 4.17A shows low and high magnification SEM images of the C-D36/T48 specimen’s surface after soaking in water at 37°C for 10 days. The sample exhibits solid microstructure with porous surface. According to SEM analysis, for the C-D36/T48 sample soaked in SBF at 37°C the first isolated morphologies of precipitated apatite crystals occur already after three days of incubation, whereas after ten days an extensive formation of sub-micrometer-sized scaly-like apatite crystals is observed (Fig. 4.17B). Closer examination of the surface of the present specimen by cross section SEM confirms the apatite crystal growth, as well as shows the presence of knoll-like aggregates, which together generate a continuous layer on the sample surface (Fig. 4.17C). The phase purity of the C-D36/T48 cement immersed in SBF at 37 °C for ten days was confirmed by powder X-ray diffraction (Fig. 4.18). IR spectroscopy and TGA/DTA analysis of the 10 days SBF-derived C-D36/T48 (see Annexe II Fig. AII.3, AII.4) do not reveal any significant differences cf. IR and TGA/DTA data for the 24 hours water-derived C-D36/T48 cement product (Fig. 4.11 and 4.13, respectively), thus additionally indicating phase purity of the C-D36/T48 after soaking in a SBF solution. Preparation of calcium phosphate bone-cements: Results 117 Fig. 4.17. Scanning electron microscopy results for the in vitro test of the C-D36/T48 cement, illustrating the extensive formation of HA phase on the surfaces of SBF incubated specimen. Comparison of the low magnification (main panel) and high magnification (right panel) SEM images from the surface of the C-D36/T48 sample soaked in distilled water (A) and in SBF (B) at 37°C for 10 days. The SEM images (B) clearly confirm the growth of hydroxyapatite crystals forming a continuous layer on the surface of the cement. A low magnification (left panel) and high magnification from selected regions (right top and right bottom panels) cross section SEM images (C) are provided for clarification of the HA crystal growth on the surface of the C-D36/T48 CPC immersed in SBF. Fig. 4.18. Powder X-ray diffraction pattern of the C-D36/T48 specimen incubated in a SBF solution at 37°C for ten days. Tick marks below the pattern correspond to the positions of the Bragg reflections expected for the hexagonal HA (Ca10(PO4)6(OH)2, JCPDS No. 72-1243). Note the absence of the XRD peaks related to the other phases, confirming the phase purity of the in vitro tested CPC. Preparation of calcium phosphate bone-cements: Results 118 4.2.9. Mechanical properties of the optimal CPC Compressive strength measurements were applied to investigate the mechanical properties of the optimal C-D36/T48 specimen, obtaining a value for CS of 25 ± 3 MPa. The local material properties elastic modulus (Es, related to elastic deformations) and hardness (H, related to plastic deformations) of the C-D36/T48 specimen were estimated by means of the nano-indentation technique. A representative nano-indentation load vs. displacement profile is shown in Fig. 4.19A. This load-displacement curve smoothly follows the loading function without any discontinuities or pop-in marks, confirming that no cracks arise during nano-indentation.180 According to the displacement profile, the peak load of 5 mN results in residual indentation depths ranging from approximately 250 to 300 nm. The average elastic modulus and hardness of the C-D36/T48 cement were calculated to be 23.1 ± 2.6 and 0.73 ± 0.2 GPa, respectively. The inset in Fig. 4.19A shows a typical FE-ESEM image of the C-D36/T48 specimen surface after nano-indentation, which, as a whole, exhibits clearly visible macropores, in good agreement with the above-mentioned mercury porosimetry and SEM analyses. The nano-indentation probing expectedly generates equilateral triangle-shaped microindent impressions with typical sides’ lengths of approximately 4.5 μm. No cracks are observed by FE-ESEM (Fig. 4.19A, inset), consistent with the load-displacement curve (Fig. 4.19A). Fig. 4.19. A representative load versus depth curve acquired by nano-indentation probing (peak load: 5 mN) of the C-D36/T48 cement (A). The inset is a typical FE-ESEM image of a specimen surface after nano-indentation, showing the microindent impression. AFM three-dimensional topography view (B) along with the corresponding height profile (C) from the microindent impression on a cement surface (lateral size: 10 μm by 10 μm; height: from 76.36 nm to 266.50 nm). Preparation of calcium phosphate bone-cements: Discussion 119 A view of the microindents can be observed more closely from the AFM topography imaging, owing to the higher spatial resolution of this technique, and a representative 3D AFM image of the regular triangle pyramid faceted impression is shown in Fig. 4.19B. The AFM thickness profile of the microindent (Fig. 4.19C) reveals that the surface exhibits a residual indentation depth consistent with the one observed in the load-displacement curve (Fig. 4.19A). According to the AFM observations, the appearance of pile-up around the microindents is not detected. 4.3.- DISCUSSION In the current study, the dissolution-precipitation reaction between DCPD (basic character) and TTCP (acid character) has been chosen to generate hydroxyapatite calcium phosphate bone-cements (Eq. 4.3). This setting reaction occurs most readily when both precursors exhibit high solubility, since the reaction initially takes place at the surface before becoming limited by diffusion.70 Consequently, the dissolution rate of the CPC precursors becomes a key factor that accelerates cement setting and conversion rates. It is reasonable to expect that the solubility of the DCPD and TTCP will increase with decreasing of their crystal sizes that, accordingly, will facilitate the formation of respective cements. On the basis of this concept, we pursued a strategy that includes several stages to explore advanced HA cement biomaterial. The initial stage involves the convenient synthesis of the phase-pure DCPD and TTCP crystalline compounds –samples D-raw and T-raw respectively. Next, these raw materials are mechanochemically treated using three different grinding media to reduce and to tailor the size of the particles. Further, a set of calcium phosphate bone-cements are generated, via a combinatorial approach, from the ball-milled DCPD and TTCP powders. At the latter stage, the optimal cement is selected with respect to the quickest setting time and the fastest conversion rate after 3 hours of setting reaction. Finally, this optimal CPC is subjected to detailed characterization studies employing various techniques. In our investigation, sample D-raw is a product of the precipitation reaction (Eq. 4.1), while sample T-raw is synthesized by high temperature reaction (Eq. 4.2). As a result, phase-pure compounds dicalcium phosphate dihydrate and tetracalcium phosphate were produced (Fig. 4.3A, 4.4A, Table 4.2). The SEM and PSD examination of DCPD and TTCP clearly show that these powders are composed of micrometer-sized crystals and aggregates (Fig. 4.5A, 4.5C, respectively), which are expected to exhibit relatively low dissolution rates. In order to improve the precursors contact and to enhance their reactivity, the primary crystal Preparation of calcium phosphate bone-cements: Discussion 120 sizes of the D-raw and T-raw were reduced by mechanochemical grinding. XRD data reveals that all six CPC precursors prepared by a wet rolling ball-mill process are phase-pure powders and no by-products/admixtures were found in the samples (Fig. 4.3B, 4.4B, Table 4.2). SEM and DLS analyses clearly confirm that ball-milling serves to effectively break up the large crystals and aggregates of TTCP and DCPD raw materials (Fig. 4.5B, 4.5D), while PSD data obviously indicates that employing smaller grinding media promotes the formation of finer particles with narrowing particle size distributions (Table 4.2, Fig. 4.6). Many research groups have previously investigated the preparation of the cement biomaterials based either on dicalcium phosphate dihydrate or dicalcium phosphate anhydrous (DCPA) and tetracalcium phosphate precursor systems.88, 94, 156, 165, 166, 181-185 The starting TTCP materials were normally approximately 10 times greater in crystal sizes than the DCPD or DCPA ones,88, 94, 156, 165, 181, 182, 185 since originally Chow and Takagi stated that CPC can only be solidified with this particular size ratio of starting compounds.181 In contrast, in other reports, HA CPCs were successfully prepared by varying the crystal sizes of both precursors183, 184 or exclusively the dicalcium phosphate precursor while the size of tetracalcium phosphate crystals remained constant.185 In these efforts, the common trend of a progressive decrease in the setting times with decreasing orthophosphates crystal sizes was observed, showing that the TTCP size effect is predominant.183 Additionally, Otsuka et al. reported that the crystal sizes of both precursors affected the CPCs’ compressive strength performance, wherein the optimal results were obtained when both precursors have the crystal sizes of about 1 μm.184 It should be stressed that upon examination of the literature related to HA CPCs on the basis of the dicalcium and tetracalcium phosphates (particle size of TTCP larger than DCPD ~10 times), it is apparent that the TTCP precursors were not entirely converted to the apatitic cements after 24 hours of setting reaction and in all cases the as-produced CPCs still contained unreacted TTCP.88, 94, 156, 165, 182, 185 The biocompatible disadvantage of such cements with a non-homogeneous composition is clear. It is believed that unconsumed TTCP is related to the greater size of TTCP crystals cf. DCPD ones and also can be connected to the formation of HA layer on the surface of the reactants.186 The results of all reviewed studies demonstrate that only Greish and Brown have reported before a cement formulation which is totally converted to HA after approximately 24 hours of setting reaction.166 This was achieved by using a liquid phase containing phosphate ions, which significantly increased the saturation of the liquid phase with PO43− ions compared to traditional water mixing liquid.88 Preparation of calcium phosphate bone-cements: Discussion 121 However, as-prepared CPC exhibits a high pH for a long period of time (~7 days), prone to generating cell damage when employed in vivo.166 In the current investigation of CPCs synthesis, a combinatorial approach, utilizing equimolar amounts of ball-milled DCPD and TTCP powders with different crystal sizes, was used (Table 4.3). In sharp contrast to the published data, after 24 hours of setting reaction, all nine as-prepared CPCs have very similar powder X-ray diffraction patterns and show only evidence for nanocrystalline hydroxyapatite (Fig. 4.8), i.e. the conversion reaction is complete after one day. One potential reason for this is that the mechanochemically-derived powders are sub-micrometer sized and usually characterized by storing a certain amount of the received mechanical energy, therefore, exhibiting enhanced reactivity owing to their capability to consume this extra energy in different ways.86 The setting time and conversion rate studies of the nine combinatorially-derived cements indicate that it is C-D36/T48 CPC which exhibits the quickest achievement of setting (~22 minutes, Table 4.3) and the highest percentage conversion to HA (~33 %) at a moderated setting reaction time of 3 hours (Fig. 4.9). It is noteworthy that this setting time fulfils the medical47 requirements’ (Section 2.5.2) and that it was achieved by using exclusively water as the liquid mixing phase, i.e. without any of the setting accelerators widely used in cement formulations88 that may potentially generate biocompatibility problems.71, 72 The insert data conclusively support that C-D36/T48 cement, prepared from D-36 and T-48 powders with average crystal sizes of ~1.13 and ~1.52 μm, respectively, is the optimal one among the nine combinatorially-derived CPCs (Table 4.3). Furthermore, it was established that C-D36/T48 is entirely converted to HA already after 6 hours of setting reaction (Fig. 4.10). To our knowledge, this is the highest reaction rate observed for CPCs based on dicalcium and tetracalcium phosphates, an important step toward enhanced biological and physiological characteristics. According to the published data, complete hydrolysis of dicalcium phosphate to HA can only be achieved in very dilute suspensions (S/L ratio ~0.01),187 whereas tetracalcium phosphate with particles size of about 1.4 μm (similar to our T-48) is only partially hydrolysed to HA after 24 hours of incubation at 37°C (S/L ratio : 2 g/mL).50 Therefore, it is believed that the observed rapid formation of HA during C-D36/T48 setting is a result of a straightforward reaction between ball-milled D-36 and T-48 (Eq. 4.3) and not product of their hydrolysis. Although we have not attempted to investigate the kinetics of the optimal cement formation, a qualitative understanding of the reaction pathway can be deduced from the pH Preparation of calcium phosphate bone-cements: Discussion 122 evolution as a function of the reaction’s duration (Fig. 4.20). It should be emphasized that the solubility of TTCP is higher than that of DCPD in acidic and neutral pH71 and, as expected, at the beginning of the dissolution-precipitation reaction, the suspension pH rises up very quickly to a value of 11.4, presumably due to the release of OH− ions during TTCP dissolution evolving water (Eq. 4.4). However, at pH higher than ~8.5, tetracalcium phosphate becomes progressively less soluble than dicalcium phosphate dehydrate (Fig. 2.18).53 Hence, at highly alkaline pH, DCPD starts to dissolve at a greater rate over TTCP (Eq. 4.5), leading to the reduction of the pH owing to the release of H+ ions (Eq. 4. 6). −−+ ++⎯⎯→⎯+ )( )( 3 4 2 )()(2)(244 224)( 2 aq aq aq OH aqs OHPOCaOHOPOCa (4.4) )(2 )( 2 4 2 )()(24 22 2 aq aq aq OH sOHHPOCaOHCaHPO ++⎯⎯→⎯⋅ −+ (4.5) )( 3 4 )( )( 2 4aq aq aq POHHPO −+− +→ (4.6) Fig. 4.20. pH versus time curve for the C-D36/T48 cement (S/L ratio : 0.05 g/mL). As a result of DCPD (predominant) and TTCP co-dissolution, after approximately one and a half hours, the suspension pH reaches the value of ~8.45 (Fig. 4.20), which remains constant until the end of the C-D36/T48 CPC setting reaction (t = 6 hours). The observed steady-state pH value is consistent with the pH value of ~8.5 representing DCPD-TTCP invariant (singular) point (Fig. 2.18),53 suggesting that a pseudo-metastable equilibrium is established in the suspension for the remaining period of time. This equilibrium is characterized by a continuous saturation of the liquid phase with respect to dicalcium dihydrate and tetracalcium phosphates and supersaturation with regards to hydroxyapatite, thus providing the driving force toward the nucleation and growth of a Preparation of calcium phosphate bone-cements: Discussion 123 more thermodynamically stable (less-soluble) calcium phosphate compound (HA) under these conditions (Fig. 2.18).53 As long as DCPD and TTCP are present in excess and their rates of dissolution are greater than the rate of HA precipitation, the liquid phase of the cement will remain near the DCPD-TTCP singular point. Therefore, the solution will stay at a quasi-constant pH and composition, allowing the HA formation to progress at a steady rate. However, with the exception of stage I (Fig. 4.10B), the conversion of C-D36/T48 to HA does not occur at a constant rate (zero-order reaction), indicating that the kinetic effects have profound influence on C-D36/T48 setting. The observed sigmoid-like plot for the conversion (Fig. 4.10B) suggests that the formation and growth of HA nuclei is the rate-limiting process of the setting reaction.168 Furthermore, according to our conversion data, D-36 is consumed faster than T-48 (Fig. 4.10A), which is consistent with most previous reports on the dicalcium-tetracalcium phosphates CPC systems.88, 94, 156, 182, 185, 188 Although more data are needed to be certain, it is believed that, in our case, the enhanced dissolution rate of D-36 is associated with an intrinsic capacity of CaHPO4·2H2O for in situ releasing of H2O during dissolution, thus, taking advantage over T-48. Specifically, in situ released water facilitates the dissolution of dicalcium phosphate dihydrate, while the dissolution rate of tetracalcium phosphate is likely to be limited by the diffusion of water to the TTCP crystals; therefore D-36 and T-48 are consumed at different rates. One important feature of this data is that the C-D36/T48 setting should result in the formation of Ca-deficient HA, since it is expected that stoichiometric HA would form directly only when the reactants are consumed at similar rate. With regards to the application of CPC in the human body for bone repair and substitution, whereby it should be similar to biological apatites, it is also beneficial to investigate the structural, textural and mechanical properties, as well as bone-bonding capacity of the selected optimal cement. Therefore, detailed characterization studies were carry out on the C-D36/T48 CPC. According to XRD, the 24 hour-derived C-D36/T48 sample is phase-pure nanocrystalline hydroxyapatite (Fig. 4.8). EDX study clearly confirms that this sample is Ca-deficient HA, having an average Ca/P molar ratio of ~1.49. This result is consistent with the different CPC precursors’ dissolution rate (Fig. 4.10A). Also, the obtained CPC present a partial substitution of the PO43− anions by HPO42− and CO32− anions in the HA crystal structure, as established by IR spectroscopy (Fig. 4.11), combustion bulk elemental analysis, Raman scattering (Fig. 4.12) and TGA (Fig. 4.13). IR spectroscopy also confirms water incorporation into the crystalline CPC product. Based on these data, the complex formulation Preparation of calcium phosphate bone-cements: Discussion 124 Ca10-x-y(HPO4)x(CO3)y(PO4)6-x-y(OH)2-x-y(H2O)z is a possible representation of 24 hour-derived C-D36/T48 CPC, however the exact formula was not established in the current study. It is worth noting that the Ca-deficient HA is more similar to bone mineral than stoichiometric HA and, consequently, exhibits better osteogenic and osteoconductive properties.189 Moreover, biological apatite contains 3-5 wt.-% carbonate groups,9 thus prepared carbonated Cadeficient HA CPC is the requirement rather than a problem for its application in medical practice. BET/BJH and mercury porosimetry studies of the 24 hour-derived C-D36/T48 sample reveal some remarkable features related to the textural properties of this cement. The as-prepared sample possesses a relatively high specific surface area (~169 m2/g), high total pore volume (~0.41 cm3/g) and average mesopore size ~55 Å, as estimated by BET/BJH methods (Fig. 4.14). It should be noted that CPCs are normally characterized by developed macroporous structure owing to the water diffusion during the setting reaction.190 Clear evidence for macroporosity in C-D36/T48 cement can be found from SEM and FE-ESEM images (Fig. 4.15B and Fig. 4.19A, inset). According to the mercury porosimetry analysis, the 24 hour-derived C-D36/T48 CPC possesses ~35 % porosity, however only the presence of mesopores with an average diameter of ~50 Å was determined for this cement. From these data, it can be deduced that macroporous structure observed by SEM and FE-ESEM is likely not mechanically stable at the relatively high pressure of mercury porosimetry experiment (up to 60000 psia), leading to the shrinkage of the CPC structure. Therefore, the presence of macropores was not detected by this technique. According to the SEM data, C-D36/T48 CPC setting reaction results in a closely associated microstructure (Fig. 4.15B). Detailed TEM, ED and HRTEM studies on the 24 hour-derived C-D36/T48 reveal that the sample has the HA crystal structure, devoid of other impurity phases (Fig. 4.16). TEM clearly confirms that the close-packed microstructure of this CPC is generated by strongly interlocked nanocrystallites (Fig. 4.16A) and it appears that namely such spontaneous assembly via interlocking of the precipitated nanocrystals results in the structural solidity of the as-prepared CPC biomaterial. The materials exhibit mesoporosity due to voids in between the interlocked nanocrystals. HRTEM results also demonstrate that those nanocrystallites do not tend to orient with their axis along aand b-axes, indicating that the precipitated Ca-deficient HA nanocrystals exhibit a much higher tendency to grow along the c-axis, the natural tendency of HA. The potential bioactivity of the optimal C-D36/T48 cement –ability to form a direct chemical bond with surrounding bone tissue– was explored via soaking in SBF at the human Preparation of calcium phosphate bone-cements: Discussion 125 body temperature of 37°C. This bone-bonding capacity was evaluated by analyzing the formation of an apatite layer on the surface of the test sample, revealing the C-D36/T48 apatite-inducing ability. The results of the in vitro tests on the C-D36/T48 CPC clearly demonstrate a widespread formation of apatite crystals already after soaking in a SBF solution for ten days at 37°C. Furthermore, the crystals tend to generate a relatively thick and persistent apatitic layer on the specimen surface, as it can be seen in Fig. 4.17, thus proving the potentially high bone-bonding ability of the Ca-deficient HA C-D36/T48 CPC.191 Compressive strength performance of the CPCs on the basis of dicalcium and tetracalcium phosphates is described in the literature.91, 94, 95, 156, 167, 183, 190, 192-194 In preparative chemistry of CPCs, there are several key factors that have essential influences on the CS values: (i) synthesis with lower solid to liquid ratio (S/L) leads to an increasing in the porosity, resulting in low compressive strength and vice versa;94, 167, 192 (ii) products derived from precursors with reduced crystal sizes are normally characterized by higher CS values;183 (iii) the use of HA crystal seeding in the preparation course of the calcium phosphate bone-cements results in the rise of the compressive strength;91, 183 (iv) the DCPD or DCPA to TTCP ratio,192 the stoichiometry of tetracalcium phosphate94 and the variety of the incubating solutions75 also might alter the CS characteristics of the final products. Moreover, some aspects of the specimens’ fabrication for the compressive strength measurements, such as applying a compaction pressure,75, 94, 167, 192 thermal treatment of the CPC precursors95, 193 or the specimen’s aspect ratio,195 are shown to have a great impact on the final CS value. Hence, the reported CS values for the cements obtained from dicalcium and tetracalcium phosphates vary essentially from ~10 MPa156 to 174 MPa,167 depending upon the CPCs composition and experimental aspects. As an example, a value of ~174 MPa was achieved using a S/L ratio of 9 g/mL and a compaction pressure of 70 MPa.167 However, based on our experience, such a S/L ratio does not result in CPCs with putty-like behaviour, as it required in medical practice. The ultimate compressive strength value of the optimal C-D36/T48 specimen was determined to be 25 ± 3 MPa, indicating that as-prepared cement is able to bear ~255 kg/cm2. It should be noted that larger CS values can be achieved for C-D36/T48 CPC by increasing the S/L ratio, pressure and/or compaction time in the course of the specimens’ fabrication for the CS measurements. However, in the current study, the objective was to obtain a CPC under experimental conditions that closely mimic those in clinical practice, ensuring the possibility of CPC in situ shaping at low pressure for a short period of time. In addition to the compressive strength, elastic modulus and hardness of C-D36/T48 sample were estimated using nano-indentation. According to nano-indentation experiments, it