John Wiley & Sons,. Inc., 1995. [48] D.J. Shaw. Introduction to ...... [1] O.O. Omatete, M.A. Janney, and R.A. Strehlow. Gelcasting-a new ceramic forming process.
DISS ETH No. 13388
Wet Processing of Nanosized Ceramic Particles A dissertation submitted to the SWISS FEDERAL INSTITUTE OF TECHNOLOGY ZURICH
for the degree of Doctor of Technical Science
presented by
Jens Helbig Dipl. Ing. univ. born on July, 9, 1969 Germany
ACCEPTED ON THE RECOMMENDATION OF Prof. L.J. Gauckler, examiner Prof. H. Hofmann, co-examiner
Zürich, 2000
Acknowledgments I like to thank Prof. Gauckler for the opportunity to do this thesis at his laboratory, for the freedom to realize my ideas and his help and support throughout the work. He showed me that scientific research not only consists of working in the lab, but also in communicating that work to other people. I want to thank Prof. Hofmann from the EPF Lausanne that he accepted to be my co-examiner. The final version of this work was substantially improved by the help of many people, especially Prof. Bayer, Dr. Markus Hütter and Urs Schönholzer. Their patience and their advice is gratefully acknowledged. Furthermore I like to thank all the colleagues and guests at the chair of nonmetallic materials for the good times, especially Markus Hütter, Urs Schönholzer, Elena Tervoort and Julia Will, who had to show utmost patience with my ’stupid’ questions. To work together with them always was fun. They also helped me a lot by simply being there. Thank you. Prof. Margarida Almeida for the long time which we were working together and without whom some of this work would not have been possible. all the other members of the colloid chemistry/ceramics processing group, who always were ready to discuss even the most bizarre ideas: Beate Balzer, Frank Filser, Daniela Hesselbarth, Martin Hruschka, Lorenz Meier und Hans Wyss. Peter Kocher for his support in all questions concerning machines and other problems and Irene Urbanek who helped me many times with the administration. my students Rene Hummel, Reto Joray, Reto Kessler, Gabor Markus, Steve Merillat, Cyrill Rüegg and Michael Zäch. I wish to thank my friends outside the institute for their support and many ’recreative’ night meetings, especially Conny, Dirk, Erik, Lars, Peter, Ronny and Walter. Last but not least, I like to thank my parents and my brother for always believing in and supporting me.
Financial support by KTI is gratefully acknowledged.
Table of Contents Synopsis/Zusammenfassung ....................................................... 1 1 Examples of Structural Ceramic Nanomaterials .................. 5 1.1 Examples of Processing of Nanosized Ceramic ............................ 6 1.1.1 Single Phase Nanoceramics ...........................................................................6 1.1.2 Nano-SiC Composites ....................................................................................7 1.1.3 Other Nanocomposite Systems ......................................................................9
1.2 Mechanical Properties of Nanomaterials .................................... 11 1.2.1 Strength ........................................................................................................11 1.2.2 Hardness .......................................................................................................19 1.2.3 Toughness ....................................................................................................20 1.2.4 Plasticity and Creep .....................................................................................23
1.3 Concluding Comments .................................................................. 25 1.4 References ...................................................................................... 27
2 Aim of the Study .................................................................... 37 3 Wet Processing of Ceramic Suspensions ............................. 39 3.1 Introduction ................................................................................... 39 3.1.1 Wet Processing of Submicron Alumina Powders ........................................39 3.1.2 Influence of Particle Shape on Wet Processing ...........................................43 3.1.3 Processing Method .......................................................................................43
3.2 Experimental .................................................................................. 46 3.2.1 Powders Used and Particle Characterization ...............................................46 3.2.2 Slurry Preparation ........................................................................................47 3.2.3 Viscosimetry ................................................................................................48 3.2.4 Electrophoretic Mobility and Surface Charge .............................................48 3.2.5 Electric Conductivity ...................................................................................48
3.3 Results and Discussion .................................................................. 49 3.3.1 Particle Size and Particle Size Distribution .................................................49 3.3.2 Viscosity of urea containing suspensions ....................................................51 3.3.3 Liquefying Effect of Urea Additions to Particle Suspensions .....................55 3.3.4 Liquefying Effect of Urea in TM-DAR and SiC Slurries ............................65 3.3.5 Amount of Urea Dissolved .......................................................................... 67 3.3.6 Estimation of the Double Layer Thickness with the DLVO-Approach .......68
3.4 Summary and Conclusions ........................................................... 70 3.5 References ...................................................................................... 71
4 Strength of Coagulated α-alumina-Boehmite Wet Green Bodies .................................................................. 77 4.1 Introduction ................................................................................... 77 4.2 Experimental Procedure ............................................................... 78
4.2.1 Slurry Preparation ........................................................................................78 4.2.2 Slurry Evaluation .........................................................................................80 4.2.3 Casting .........................................................................................................82 4.2.4 Measurement of Wet Green Body Strength .................................................82 4.2.5 Drying, Sintering and Microstructure ..........................................................83
4.3 Results and Discussion .................................................................. 86 4.3.1 Electrophoretic Mobility ..............................................................................86 4.3.2 Viscosity .......................................................................................................86 4.3.3 Compressive Wet Green Body Strength ......................................................88 4.3.4 Sintering and Microstructure .......................................................................94
4.4 Conclusions and Summary ......................................................... 100 4.5 References .................................................................................... 101
5 Centrifugal Casting with Submicrometer Alumina Powders ................................................................. 105 5.1 Introduction ................................................................................. 105 5.2 Experimental Procedure ............................................................. 108 5.2.1 Slurry Preparation ......................................................................................108 5.2.2 Mold preparation ........................................................................................109 5.2.3 Casting .......................................................................................................109 5.2.4 Drying ........................................................................................................110 5.2.5 Sintering .....................................................................................................111 5.2.6 Microstructural Evaluation ........................................................................112 5.2.7 Hardness and KIc .......................................................................................112
5.3 Results and Discussion ................................................................ 115 5.3.1 Casting .......................................................................................................115 5.3.2 Drying ........................................................................................................116 5.3.3 Sintering .....................................................................................................118 5.3.4 Microstructure ............................................................................................120 5.3.5 Hardness and Toughness ............................................................................126
5.4 Summary and Conclusions ......................................................... 129 5.5 References ................................................................................... 130
6 Outlook ................................................................................. 133 6.1 Wet Processing of Ceramic Supensions ..................................... 133 6.2 Strength of Coagulated α-alumina-Boehmite Wet Green Bodies ......................................................................... 134 6.3 Centrifugal Casting With Submicrometer Alumina Powder .. 135 6.4 References .................................................................................... 137
A.1 Fitparameters of the Viscosity Measurements .............. 138 A.1.1 Parameters from Chapter 3 .................................................... 138 A.1.1.1 Viscosity of Slurries .....................................................................138 A.1.1.2 Viscosity and Conductivity of Urea Solutions .............................139
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Synopsis The research of nanoscaled ceramics is still of great interest. This lies in the promise of improved material properties of these materials. In the recent years, the processing and the properties of single phase nanoceramics and nanocomposites was greatly advanced. Unfortunately, the processing of single phase nanoceramics mostly utilizes elaborate processes, such as hot pressing. Even with these advanced forming techniques, no dense nanoscaled monolithic material could be produced, apart from few notable exceptions. For further research and for large scale applications a more simplistic approach is desired. In this work, the possibilities of processing nanoscaled alumina particles with the goal to achieve dense, nanostructured single phase α-alumina are tested. Wet processing of ceramic particles is chosen as a simple way to produce homogenous green bodies with high green densities. To show the possibilities of wet processing of nanoparticles, the three stages from dispersing the powder, to the forming of green bodies and finally to the sintering of the consolidated nanoparticles are addressed in this work. In the first part, the influence of the particle size of the particles on the viscosity of the suspension in dependence of the ionic strength is investigated. The smaller the particles, the lower the true solids loading is becoming, as the double layer takes up more volume. With variations in the ionic strength this influence in high solids loading suspensions could be quantified. Furthermore, a model for the influence of the double layer on the effective volume based on the Debye-Hückel equation of the double layer thickness is developed. The second part of this thesis deals with the solidification of nanoscaled particles with the DCC process. It is shown that boehmite particles can be solidified with the same mechanisms as used for larger α-alumina. The wet green body strength of coagulated bodies of
α-alumina and boehmite, as well of bodies made of mixtures of these two powders is also investigated. It is shown that the addition of boehmite particles (which are ten times smaller than the α-alumina particles) to high solids loading α-alumina bodies doubles the strength to over 20 kPa. The addition of α-alumina to boehmite rich bodies does not result in a significant increase of the strength. The strength of the wet particulate bodies is explained with the development of force chains under compressive stress. The strengthening of the α-alumina rich bodies by the addition of the smaller boehmite particles is attributed
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to an increase of particles at the contact points of two α-alumina particles, thus preventing the collapse of the force chains at lower stresses. The third part of this study looks into the use of centrifugal casting as a means of solidification of nanoscaled particles. High solids loading suspensions could be densified to very high green densities of approximately 74 % theoretical density with centrifugal accelerations of up to 4700 g. After sintering the bodies are all at almost full density. The microstructures exhibits grain sizes of between 0.36 µm and 0.7 µm, depending on the solids loading of the starting slurry and the centrifugal acceleration, the higher both, the smaller the grain size. Hardness and toughness are evaluated using Vickers indentations.
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Zusammenfassung Die Festigkeit von Keramiken bleibt durch inhärente Fehler in den Bauteilen immer weit unterhalb der theoretisch möglichen Festigkeit. Deswegen ist die Erforschung von keramischen Nanomaterialien mit weniger und kleineren Fehlern und der damit verbundenen Aussicht auf verbesserte Materialeigenschaften immer noch von grossem Interesse. In den vergangenen Jahren wurden die Herstellungsmethoden und die Eigenschaften von einphasigen Nanokeramiken und Nanokompositen stark verbessert. Trotz der fortgeschrittenen, aufwendigen Formgebung, wie etwa Heisspressen, konnten jedoch bis auf ein paar wenige bemerkenswerte Ausnahmen keine dichten einphasigen Nanokeramiken hergestellt werden. Um eine weitere Erforschung und die Anwendung im grossen Massstab zu ermöglichen sind einfachere Herstellungsmethode erwünscht. In dieser Arbeit wurden die Möglichkeiten der Herstellung von dichten, einphasigen
α-Aluminiumoxid Keramiken aus Nanometer grossen Aluminiumoxid-Partikeln untersucht. Um homogene Grünkörper mit hohen Gründichten zu erzeugen, was eine Vorraussetzung für möglichst hohe Endfestigkeiten ist, wurden Nassformgebungsverfahren gewählt. Die Verbesserungsmöglichkeiten der Nassformgebung von Nanopartikeln wurden in dieser Arbeit bei der Aufbereitung der Pulver, beim Direct Coagulation Verfahren (DCC) und beim Zentrifugalguss untersucht. Im ersten Teil wird der Einfluss der Partikelgrösse auf die Viskosität der Suspensionen in Abhängigkeit der Ionenstärke untersucht. Je kleiner die Partikel sind, desto niedriger wird der realisierbare Feststoffgehalt da die Doppelschicht relativ zum Teilchenvolumen immer mehr Volumen einnimmt. Dieser Einfluss wurde in hochfestoffhaltigen Suspensionen von 15 bis 28 Vol % durch Veränderung der Ionenstärke quantifiziert. Basierend auf der Debye-Hückel Gleichung wurde ein Modell abgeleitet, das den Einfluss der Doppelschicht des kolloiden Teilchens einbezieht und dem Ganzen ein neues scheinbares Volumen zuordnet, welches für die rheologischen Eigenschaften wirksam wird. Der zweite Teil dieser Arbeit handelt von der Koagulation von Nanopartikel mittels dem DCC Prozess. Es wird gezeigt, dass Böhmit Suspensionen mit dem gleichen Mechanismus destabilisiert werden können wie grössere α-Aluminiumoxidpartikel. Die Festigkeit der nassen Grünkörper aus α-Al2O3 und Böhmit, als auch von Körpern die aus einer Mischung der beiden Pulver bestehen, wurde untersucht. Die Zugabe von Böhmitpartikeln
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(die zehnmal kleiner sind als das α-Al2O3) zu den hochfeststoffhaltigen α-Al2O3 Suspensionen erhöht die Grünfestigkeit der daraus geformten Körper von 10 kPa auf 22 kPa. Die Zugabe von α-Al2O3 zu Böhmitsuspensionen führt zu keiner signifikanten Zunahme der Grünfestigkeit. Die Festigkeit der granularen, nassen Körper wird mit der Entstehung von ’force chains’ (Kraftlinien) unter Druck erklärt. Diese Festigkeitssteigerung bei den
α-Al2O3 reichen Körpern durch die Zugabe der kleineren Böhmit Partikeln lässt sich durch die Anlagerung von mehr Partikeln an den Kontaktpunkten von zwei Al2O3 Partikeln erklären, was die Forcechains verstärkt und deren Zusammenbruch bei normalen Belastungen verhindert. Im dritten Teil wird Zentrifugalguss als Formgebungsverfahren von Nanopartikeln untersucht. Hochfestoffhaltige Suspensionen konnten zu sehr hohen Gründichten von etwa 74 % der theoretischen Dichte durch Beschleunigungen von bis zu 4700 g verdichtet werden. Nach dem Sintern waren alle Körper mit einer Dichte von mehr als 98 % der theoretischen Dichte fast vollständig dicht. Die Korngrössen in den gesinterten Mikrostrukturen lagen zwischen 0.36 µm und 0.7 µm. Diese Korngrössen sind die niedrigsten die bisher bei dichten, drucklos gesinterten Aluminiumoxid erreicht wurden. Je höher die Beschleunigung und der Feststoffgehalt, desto kleiner wurde die Korngrösse.
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1. Examples of Structural Ceramic Nanomaterials Over the last 10 years the development and the properties of nanomaterials got increasing attention. These materials promise improved mechanical behavior as well as special properties for electric, magnetic and catalytic applications. Theoretically, nanomaterials can be defined in two ways. Verweij [1] defined nanomaterials generally as materials with structures or special features in a size range below 100 nm. Roy et. al. [2] characterized nanocomposites as materials containing more than one phase with at least one of the phases in the nanometer range. These definitions of nanomaterials do not distinguish wether one or all of these phases either exist in amorphous, semicrystalline or crystalline states. For nanoceramics and ceramic nanocomposites however at least the dominant phase, i.e. the matrix phase, is a (poly-)crystalline ceramic material. Generally, in the ceramics community, all materials whose special features can be expressed in nanometers are considered to be nanomaterials. When nanoceramics are defined this way, the emphasis of the research lies on a change of properties instead of purely evaluating the microstructure. The research on nanoceramics concentrated on two main areas: 1) the properties of the materials and 2) the processing of these nanoceramics. Thereby it is not recommended to separate these two from each other, as the processing influences the final structure as well as the resulting properties. Nanoceramics can also be divided into dense and porous materials which are either singlephase ceramics with nanoscaled structures or nanocomposite ceramics. In the case of the composite materials pores are not regarded as an additional phase.
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1.1. Examples of Processing of Nanosized Ceramic 1.1.1. Single Phase Nanoceramics The most prominent representatives of single phased ceramics are Al2O3, ZrO2, SiC and Si3N4. Nanoscaled microstructures of these materials are processed using high external pressures and low sintering temperatures to avoid grain coarsening in the final stage of sintering [3-20]. These processing techniques limit their usefulness for widespread applications. Gallas et al. and Gonzales et al. used a diamond anvil cell with pressures of up to 5.6 GPa at room temperature to compact γ-alumina [13,14,15]. Depending on the compaction pressure the densities of these samples varied between 50 and 90 % of the theoretical density. Gonzales heat treated his samples after compaction at 1600oC, but a porosity of about 5 % still remained. Mishra et al. [9,21] and Nordahl et al. [8,10] used pressures of 1 GPa respectively 100 MPa at temperatures of around 1000oC for the compaction of nanocrystalline alumina. Again the resulting densities depended on the pressure applied and reached about 95% TD for high pressures. All of this materials were not really nanosized according to the definition given by Verweij since the observed grain sizes were between 200 and 500 nm in the sintered state. Comparable compaction techniques were used to produce nanoscaled zirconia. Chaim et al. [5,11], Winnubst et al. [6] and Chokshi [4] applied hot isostatic pressing at temperatures from 1000 to 1600oC and pressures between 400 MPa and 1 GPa. They achieved dense zirconia ceramics with grain sizes between 100 and 250 nm. Radonjic [7] and Theunissen et al. [3] used precompacted specimens by uniaxial pressing and sintered them at temperatures from 1100oC to 1600oC. They achieved nearly full densities for samples with grain sizes exceeding 150 nm. Radonjic showed that samples which were sintered at lower temperature retained the nanometer scale but their density never exceeded 80 % TD. Single phased nanostructured silicon carbide ceramics were produced by Mitomo et al. [19] and Upadhya [18] by hot isostatic pressing with 100 MPa pressure at 1750oC. They reached almost full density with grain sizes between 100 to 500 nm and of 100 to 250 nm respectively. Although the four structural ceramics mentioned could not be sintered to full density re-
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taining grain sizes smaller than 100 nm, they still are considered as nanoceramics. The only monolithic ceramic which could be compacted up to 95 % TD with grain sizes smaller than 100 nm was TiO2, as shown by Hahn and Siegel et al. [22,23,24,25] who used hot pressing at 700oC and 100 MPa to 1 GPa pressure. Karch and Birringer [26] pressed samples of titania in vacuum with 1-2 GPa at 500-1000oC and achieved grain sizes of about 150 nm at 99 % TD. Mayo et al. [27] showed that pressureless sintered TiO2 at temperatures of 700-900oC retained the small initial grain size of 12 to 30 nm. However they found that samples prepared this way have densities less than 80 % TD.
1.1.2. Nano-SiC Composites Because single phased nanoceramics are very difficult to achieve and the results mostly are not satisfying with regard to density and grain size, further efforts concentrated on the preparation and evaluation of ceramic composites. One of the first researchers who adopted the concept of nanocomposites given by Roy [2] were Niihara et al. with the investigation of Al2O3-SiC, Al2O3-Si3N4 and MgO-SiC nanocomposite systems [28-36]. He showed that the inclusion of nanometer sized SiC into an alumina matrix increases fracture toughness and strength. These materials were produced by hot pressing at 1500-1900oC with 300 MPa pressure in N2 atmosphere. The resulting alumina matrix showed grain sizes of 2 to 10 µm with nanometer sized (10 to 250 nm) inclusions of silicon carbide (see Figure 1.1). Later on Niihara and his co-workers achieved alumina/SiC nanocomposites also by pressureless sintering [33]. Other matrix materials, such as zirconia, were used as well [32]. Other groups repeated the experiments of Niihara [37-52], and managed to achieve nano-SiC/alumina matrix composites. Usually the SiC nanopowder and the alumina powder were mixed by ball milling or in an attritor as either aqueous or organic suspensions. Special care had to be taken that no agglomerates remain after drying of the powder mixtures. Consolidation usually was carried out by hot pressing in Ar or N2 atmospheres at 1550-1800oC and between 20 and 40 MPa pressure. Pressureless sintering under inert gas atmosphere was carried out with injection molded and slip cast samples [44,48], and here the achieved densities were only slightly lower than for hot pressed samples with full densities. Gao et al. [53] used spark plasma sintering to densify a mixture of nano-SiC and alumina to almost full density at a temperature between 1350 and 1550oC
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Figure 1.1: Microstructures of a single phase alumina (A) and an alumina/5 vol % SiC nanocomposite [39] and a heating rate of 600 K/min. Bamba et al. [32], Miao et al. [54] formed ZrO2/SiC composites by hotpressing a mixture of the powders. With temperatures above 1600oC and a pressure of 20 MPa a fully dense body could be achieved. Depending on the amount of added SiC a refinement of the zirconia matrix was observed: the more SiC particles were present, the narrower was the grain size distribution of the matrix. SiC/Si3N4 composites with SiC nanoparticles in a Si3N4 matrix were prepared by Niihara and co-workers [55,56]. The preparation of these samples was comparable to that of the alumina/SiC nanocomposites. The powders were dispersed in ethanol together with sintering aids such as alumina or yttria and then dried. The powder mixture was then hotpressed at 1850oC under 30 MPa pressure in N2. This achieved a fully dense material. Other groups tried to reproduce the results of Niihara with other techniques. Akimune et al. [57] tried cold isostatic pressing the samples which then were pressureless sintered in N2 at 1700oC followed by HIPing at 1850oC with 100 MPa. A similar approach was later used by the group of Niihara [58]. Polymer pyrolysis was chosen by Riedel and his group [59,60,61,62] to produce SiC nanoparticles in-situ in a Si3N4 matrix with alumina and yttria as sintering aids. After pyrolysis at 1000oC under Argon and pressureless sintering
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between 1750 and 1850oC, these cold isostatically pressed samples had a density of 97 % TD, the size of the SiC particles in the Si3N4 matrix was about 200 nm. Zhang et al. found a different way to prepare SiC/Si3N4 nanocomposites [63] by crystallization of glass-coated SiC particles and subsequent hotpressing. They achieved dense SiC ceramics containing 20-30 wt % SiAlON particles with a grain size of 50 nm which were dispersed along the SiC grain boundaries.
1.1.3. Other Nanocomposite Systems Apart from the best investigated systems SiC in alumina and SiC in Si3N4, also SiC in MgO respecteively TiN in alumina should be mentioned. MgO/SiC composites are produced similar to the alumina/SiC nanocomposites, only differing in the higher hotpressing temperatures ranging from 1700 to 1900 o C [31,64,65,66,67]. The achieved material was fully dense, only varying in the grain sizes of MgO between 3.0 µm for 5 vol% SiC addition to 1.6 µm for 30 vol % SiC. ZrO2 ceramics toughened by SiC nanoparticles were described by Bamba et al. [32]. They produced a fully dense material by hotpressing between 1600 and 1900oC at 30 MPa in Ar-atmosphere. Depending on the SiC content they achieved a matrix grain size of 25 µm for no and of about 1 to 2 µm for 20 vol % SiC. The SiC particle sizes remained in the range of 150 to 500 nm. Nanoscaled TiN in an alumina matrix was achieved by attritor milling of the initial powders, drying and subsequent hotpressing at 1700oC. The TiN showed a rather large grain size of 1µm after sintering at full density [68].
As can be seen, it is possible to produce nanoscaled ceramics with a wide variety of processing methods. Most often, hotpressing or hot-isostatic pressing are applied to reduce the grain growth of the material. Single phase oxides with grain sizes in the range of few hundred nanometers were not produced with more conventional processing techniques. Although applying high pressure many sintered nanocomposites were not dense. The remaining porosity is in some cases desired, where the achieved specimens are used as catalysts or filters. For structural, load bearing ceramics the remaining porosity hinders the application of the material substantially. The production of nanocomposite ceramics is comparable to the processing of single
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phase nanoceramics. Here additionally the preparation of the powder mixtures has to be done carefully to avoid the formation of inhomgeneities and agglomerates. In almost all cases the pressing was carried out in inert atmospheres to retain the non-oxide phases. Most nanocomposites didn’t reach full density, which is basic for a comparison of the mechanical properties with conventional dense ceramics. As long as the processing can not be developped into easier to use and cheaper methods, the production of nanocomposites will be restricted to the laboratory and the potential of nanoceramics will not be exploited industrially.
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1.2. Mechanical Properties of Nanomaterials In the following chapter the mechanical properties of nanomaterials will be addressed. It will be distinguished between monolithic and composite materials wherever it is meaningful.
1.2.1. Strength a) Strength of Single Phase Nanoceramics Among the mechanical properties which change drastically with a reduction in grain size G is the strength σ. This relation between strength and grain size was already well known for metallic materials. One of the first researchers to consider this relation for ceramics was Carniglia [69], who showed that MgO, Al2O3 and BeO exhibited two distinctive different regimes for the strength-grain size relationship. Schematically, this is shown in Figure 1.2 (after Rice [70]). Three different regions can be identified for the strength-grain size dependence of polycrystalline materials (see Figure 1.2). The first region refers to the dependence of the strength on the coarser grains. In this region, the size of existing flaws are on the same scale as the grains or smaller. A crack will grow stably until the ratio
γ(c)/c of the fracture-surface energy and the crack length reaches a maximum. After that, failure will occur. This ratio will increase with decreasing grain size until the crack will be controlled only by the polycrystalline fracture-surface energy, i.e. the KIc [71]. This is true for ceramic bodies where the initial flaw sizes are identical. Different initial flaw sizes lead to deviations in the σ/G-slopes and to varying initial strengths. Alford et al. proposed the largest grain as initiator of the fracture [72], thus leading to higher strengths for smaller grains. In earlier reports the strength for extremely coarse grained ceramics was extrapolated to zero [69,73,74]. This evidently is not reasonable, as single crystals also have a strength. Later Rice corrected this extrapolation considering the strength of the single crystals as well [70,75,76]. As the extrapolated strength data would hold for an interception at zero strength, an anomaly in the strength-grain size relation shows up (see Figure 1.2 for large grains). Rice explained this behavior with stresses which develop at grain boundaries, with initiation of fracture at grain boundaries with a KIc smaller than in
σ
σ
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Strength of the single crystal
Figure 1.2: σ-G-1/2-plot schematically for polycrystalline, monolithic ceramics [88]. single crystals and with elastic anisotropies [75]. For smaller grain sizes the slope of the strength-grain size levels off to smaller values or even to zero. In this region the initial flaws which lead to fracture are larger than the grains. Variations in the grain size do not influence the strength limited by the Griffith flaws anymore [34,70,71,72,75]. When the initial flaw size is kept always the same for decreasing grain sizes, the resulting strength is not expected to vary anymore. It was proposed that the special processing required for small grain sizes decreases the flaw sizes in finer grained ceramics, thus leading to a slight strength increase [40,75]. The introduction of larger flaws after sintering by grinding the samples, showed that the strength for fine grained ceramics indeed becomes independent of the grain size and only determined by
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the flaws introduced by the grinding [77,78].
For large grain sizes a Knudsen type dependence was found: 1 σ L = k 1 ⋅ -------G
(1.1)
(with G as the median grain size and k1 as a materials dependent constant) and for small grain sizes a Hall-Petch like dependence with
σ = σ2 + k 2 G
(1.2)
(σ2 and k2 constants) was derived. These equations were basically confirmed by several other authors who investigated the strength-grain size relationship of single oxide ceramics [34,70-90]. Some of these authors used the largest grain size present in their specimen instead of the median grain size as proposed by Carniglia [74,77,78,80,81,83]. This dependence was based on the assumption that the largest grains acted as initiating flaws under the Griffith criterium. Rice stated later that there are two different sorts of ceramics whose strength-grain size relationship could either be described by the median grain size or by the largest grain [70]. Reviewing his own data and those of other researchers Rice proposed that the largest grain size should be used for ceramics which exhibit nonuniform microstructures (e.g. Al2O3, β-Al2O3, B4C, SiC, Si3N4 and TiO2) in contrary to ceramics with more uniform grain sizes (e.g. MgO, Y2O3, ZrO2 and MgAl2O4), where the median grain size should be preferred. Hohjo and co-workers [90] took this one step further. They tried to establish parameters for the uniformity of the microstructure and for the grain size, exemplary for Si3N4. They then related the strength to these parameters and found again a Knudsen-type relation (see equation 1.1).
One of the most extensively researched ceramic materials for the strength-grain size dependence is alumina. Figure 1.3 shows a plot of the strength-grain size relationship for hot pressed alumina [74,78,79,82]. In this graph, circles indicate failure where homogeneous microstructures are involved, squares where exaggerated grains were present. All grrin
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Figure 1.3: σ-G-1/2 data for hot-pressed alumina measured at 22oC [82]. sizes are average grain sizes, arrows point to the grain sizes of exaggerated grains [74,79,82]. Symbols with crosses represent specimen ground after hotpressing (triangle 220 grit, diamonds 400 grit, squares 600 grit, circles 1400 grit) [78]. It can be seen, that the strength is increasing with decreasing grain size. For the ground samples, one can see that the newly introduced flaws have a bigger influence on the fracture strength than the grain size in the fine grained region. In the coarse grained region, the surface flaws do not show such big inflence. The crosshatched region summarises the range of values found by Rice in earlier reviews [79] With zirconia shows the same dependence of the strength on the grain size is observed, which is shown in Figure 1.4 for various zirconias with different stabilizers (data taken from [75]). The numbers next to the symbols show the amount of stabilizer used. Again it can be seen that the strength increases with decreasing grain size and levels of for finer
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Figure 1.4: σ-G-1/2 data for zirconia with varying stabilizers (graph taken from [103]).
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grains. The strength-grain size dependence was also confirmed for SiC ceramics. Prochazka and co-workers [80] found a Hall-Petch like relation for grain sizes between 1700 and 150 µm. The strength increased from 130 MPa for the large grains to 358 MPa for the smaller grains. Cranmer et al. [77] then derived a Knudsen-type relation for SiC with grain sizes between 103 and 6.2 µm. The strength increased with decreasing grain size from about 375 to 450 MPa for these samples. For Si3N4 only few data are available. Hohjo et al. [90] showed some strength-grain size data and proved that there is an increase of strength with decreasing grain size. However they also stated that the grain shape has a very strong influence on the strength.
All of these measurements were carried out using materials with grain sizes larger than 0.8 µm in case of alumina, 0.3 µm for zirconia and 6 µm for SiC and Si3N4. Strength data for smaller grain sizes, even for nanoceramics such as TiO2, are not available in the literature. The reason for this is the difficulty to produce fully dense specimens with such small grain sizes. Even if possible, the special processing techniques allow small samples in limited quantities. Therefore, no reliable strength data are available. More data on hardness and toughness are available for small grained (i.e. nanoscaled) materials.
b) Strength of Non-Textured Nanocomposite Ceramics The immense interest in nanocomposite materials derives from the fact that these materials exhibit superior strength and toughness compared to the monolithic matrix materials. Niihara et al. [28,32,36,55,56,59,58,64,67,91-100] found that the addition of nano sized particles of different compositions to coarser grained matrix materials increased the strength and the toughness. E.g., the addition of nano-SiC to alumina increased the strength from 350 MPa for pure alumina to 1100 MPa for alumina containing about 5 vol % nano-SiC. Further addition of SiC decreased the strength down to 800 MPa. The combination of SiC and alumina then was examined by many other authors, but nobody reached the high strength values of Niihara [46,101,102]. The strength values of those researchers steadily increased with the amount of SiC nanoparticles to the same strength which Niihara found for high SiC contents (800 MPa) (see Figure 1.5) [46,66,101,102].
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Figure 1.5: Strength and toughness values of alumina/SiC nanocomposites compared by different authors[39]. Niihara et al. (full circles) , Borsa et al. (squares) , Davidge et al. (down triangles) and Zhao et al. (up triangles) For Si3N4/SiC nanocomposites, Niihara reported a modest increase of the strength from about 900 MPa to 1100 MPa. SiC nanoparticles also increased the strength of MgO to 550 MPa when more than 10 vol % SiC were added [67]. Bamba et al. [32] examined 8YTZP/SiC nanocomposites and found an increase for the strength from about 200 MPa for pure zirconia up to 700 MPa for Y-TZP with 20 vol % SiC. The increase of the strength for nanocomposite materials is attributed to various factors. First, the addition of the nanoparticles inhibits the grain growth of the matrix grains, thus leading to the above mentioned strengthening due to the Hall-Petch relation by reducing the initial flaw size [46]. This refinement of the microstructure is accomplished by Zener grain boundary pinning [40,44,46] (see Figure 1.6), whereby the grain boundaries of the matrix grains are pinned by the smaller SiC inclusions. The quantity of this effect depends on the amount and the size of the smaller particles in the matrix. The higher the amount of the SiC, the less the matrix grains grow during sintering. In Figure 1.6 the measured grain sizes of the matrix for SiC/a-alumina nanocomposites is compared with an estimation with the Zener modell [40,103]. Monolithic alumina e.g., sintered under similar con-
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Figure 1.6: Comparison of grain refinement in an alumina matrix by the addition of SiC with the prediction made by the Zener model [39]. ditions as the alumina/SiC nanocomposite, tends to develop grain sizes of 20 µm and larger whereas in composites the average grain size usually is less than 2 µm. Another mechanism for the grain refinement and reduction of flaw sizes is the formation of dislocation networks due to the SiC inclusion, as proposed by Niihara [64,65] (see Figure 1.7). During cooling down the mismatch of the thermal expansion coefficients leads
Figure 1.7: Dislocation networks around SiC inclusions in alumina [39]
to stresses and to the formation of dislocation networks which could increase the strength. On the other hand, Sternitzke pointed out [40] that the Griffith flaws which would yield
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the strength values of Niiharas materials was 18 µm. He therefore concluded that the refinement of the microstructure is sufficient to explain the strength increase. A third explanation for the increase in strength was given by Carrol et al. [104] and Sternitzke [40]. They stated that due to careful processing necessary for nanoparticles, the size of processing derived flaws decreased drastically, thus increasing the highest possible strength. Sternitzke furthermore argued that the inclusion of SiC nanoparticles increases the R-curve effect in the material, thus leading to higher toughness and therefore to higher strength [40] (this is discussed in more detail in the following chapter on toughness). He proposed that the inclusion of the particles might lead to crack deflection and crack bowing. The latter effect was taken up and confirmed by Pezotti et al. [105] by a theoretical derivation. In addition to the already mentioned effects, several authors pointed out that the grain boundaries in the nanocomposites are strengthened and therefore grain pullout is reduced [68,106]. This is accomplished by the deflection of cracks which run along grain boundaries and by the introduction of internal stresses at the grain boundaries.
1.2.2. Hardness There is a lot of data available on hardness and toughness of single phased small grained ceramics [3,9,15,21,22,27,34,85,86,87,90,107-120]. All authors show that the hardness increases linearly with decreasing grain size. For α-alumina, Skrovanek and Bradt measured hardness values between 1700 and 2000 KHN400 for grain sizes ranging from 8 down to 1 µm [107]. Mishra et al. reported hardness values for γ-alumina between 16 and 25 GPa in the grain size range from 1.5 µm down to 0.5 µm [9]. In TZP ceramics, Ruiz et al. measured 11.3 GPa for 5.1 µm to 12.7 GPa for 0.3 µm. For the same material, Singh et al. [114] found hardness values between 12.6 and 13.8 GPa for the grain size region from 1.6 to 0.3 µm. For TiO2, Mayo and co-workers [27] determined hardness values from 9 GPa for 0.25 µm grain size up to 15 GPa for 0.01 µm by nano indentation. In ZnO they measured values between 1.5 GPa and 5 GPa for grain sizes from 0.14 to 0.02 µm. So far no limit has been found for the increase of the hardness. As one exception, some authors measured a decrease of the hardness in case of true nanoceramics based on TiO2
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and ZnO, but this can be attributed to the increased porosity in all of these materials [27,115,121]. Another exception to the general rule of increasing hardness was found by Cook et al. [122]. They showed that the hardness for MgO ceramics is independent from the grain size due to a pronounced plastic deformation zone in the material at the indentation tip. The increase in hardness with decreasing grain size is explained by Skrovanek and Bradt [107] and Siegel [119] by the increased amount of grain boundaries which is hindering the movement of dislocations, thus impeding the plastic deformation of the material.
The hardness of alumina/SiC nanocomposites was investigated by Nakihara and Niihara [123]. They found that the hardness scales linearly with the amount of SiC mixed into the Al2O3. The hardness increased from 17 GPa for pure alumina to 17.5 GPa for a 5 vol % mixture of alumina and SiC. These values were measured by the indentation method. The hardness for SiC-AlN composites was examined by Chen et al. [124], who found an increase in hardness by the AlN addition of 28.6 GPa. The combination of Al2O3/ZrO2 nano/nanocomposites was examined by Bhaduri et al. [125]. For this material they found a decrease in hardness from about 20 GPa for pure alumina to about 4.5 GPa for compositions with zirconia additions.
1.2.3. Toughness The grain size dependence of the toughness of nanoscaled materials is difficult to interpret, as the fine grained materials exhibit a pronounced R-curve behavior [86,87,108,110,111,126,127]. Considering the intrinsic crack length to estimate the toughness-grain size dependence [9,86,87,109,110,116], no dependence of the toughness on the grain size could be observed. The R-curve behavior of these materials often leads to a strong scattering in the measured toughness values [34,128]. The R-curve behavior of the monolithic, non-phase transforming ceramics is attributed to bridging grain pullout [86,108,111]. Myahara and co-workers [86] found toughness values for α-alumina at around 3-4 MPa⋅m1/2 (see Figure 1.8). They observed increased scattering of these values with decreasing grain size. Seidel and Rödel reported identical toughness for all grain sizes in alumina at the crack tip in the order of 2-2.5 and 3.4-4 MPa⋅m1/2. These values de-
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Figure 1.8: Toughness of α-alumina in dependence on the grain size [102]. pend on whether the toughness was determined by the single-edge notch beam method or by indentation methods respectively. For alumina ceramics, Koyama et al. [34] and Tuan et al. [87] found a slightly different behavior. Koyama observed a steady increase of the toughness with the grain size for alumina having nonuniform grain sizes. Tuan measured an increase of the toughness with increasing grain sizes for grains smaller than 2 µm. For coarser grains the toughness remained constant. For phase transforming ceramics, like ZrO2, another reason has to be considered. The transformation toughening of zirconia relies on the tetragonal to monoclinic phase transformation of grains due to applied stress. This stress induced transformation depends on various factors. First, the larger the transformation zone around a crack tip is, the more effective is the strengthening of the material [3]. This transformation zone increases with decreasing grain size [129]. On the other hand, if the grain size becomes too small, no transformation toughening is observed. This critical grain size was found to be around ~0.3 µm for 3 mol Y-TZP [3]. It is assumed, that if grain sizes are in the range of the crystal domains of tetragonal zirconia, the toughening occurs by transformation due to crack deflection [3]. Wang and co-workers [127], Singh et al. [114] and Ruiz et al. [113] showed that phase transforming zirconia exhibits a maximum of the toughness at a certain grain size before it drops again. Wang [127] attributed the decrease of the toughness in coarser grained to a competition of higher transformability of larger grains and to the introduction of larger Griffith flaws which decreases toughness and strength. On the other hand, he proposes that the decrease in the transformation zone and the suppressed transformation
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toughening is responsible for the lower toughness of the zirconia for very fine grains. Theunissen et al. [3] observed no change in toughness with decreasing grain size. But he found that the toughness of fine grained zirconia was comparable to that of coarse grained zirconia, despite lacking transformation toughening.
The toughness of nanocomposites is reported to increase with the addition of the nanoparticles [28,46,101,102]. Niihara et al. found that the toughness of alumina/SiC composites shows a maximum in dependence on the SiC addition [28]. It increased from 3 MPa⋅m1/ 2
for pure alumina to 4.6 MPa⋅m1/2 for Al2O3 with 5 vol % SiC. The toughness reaches a
plateau with a value of 4.1 MPa⋅m1/2 with further SiC addition. Again, other researchers could not reproduce this finding and these high values [46,101,102]. They also found an increase of the toughness with the addition of SiC nanoparticles to alumina, but the value increased moderately from 3 MPa⋅m1/2 in alumina without SiC addition to a plateau of around 3.5 MPa⋅m1/2 in composites when 25 vol % SiC were added to alumina. The toughness of Si3N4/SiC composites is strongly influenced by the morphology and the density of the Si3N4 matrix grains [55,56,130,131]. The addition of 15 vol % SiC increases the toughness from 4-5 MPa⋅m1/2 to 6.6 MPa⋅m1/2. After further addition of SiC, the toughness drops to a plateau between 5 and 6 MPa⋅m1/2. The decrease in toughness was attributed to a homogenization of the Si3N4 grains. Without SiC addition the Si3N4 exhibits elongated grains which lead to high toughness. The addition of SiC prevents the formation of these elongated grains, eventually leading to a lower toughness [55,131]. The toughness of 8Y-TZP/SiC nanocomposites increases slightly with the addition of SiC from 2 to 2.5 MPa⋅m1/2 [32]. The toughness in alumina/ZrO2 nano-nanocomposites was determined by Bhaduri et al. with 8.4 MPa⋅m1/2 [125], although no transformation toughening could be detected.
The increase of the toughness in the nanocomposite materials was already discussed in some detail in the chapter on strengthening of nanoceramics, as the increase of toughness generally results in an increase of the strength. But it still remains difficult to give an detailed explanation for the toughening mechanisms in nanocomposites, as the toughness increase generally is rather small. One reason is that these changes also strongly depend on the properties of the matrix material. In the case of Si3N4 the refinement of the matrix
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grains by addition of SiC even leads to a decrease in toughness. In all other cases two competing effects are discussed to explain the slight increase in toughness. One is the refinement of the matrix structure which might lead to internal stresses which causes a general decrease of the toughness [40]. The second reason is the increase of the toughness by the addition of a second phase which leads to crack deflection and grain boundary strengthening which both leads to an increase in toughness.
1.2.4. Plasticity and Creep One of the main reasons why the nanosized materials evoked so much interest is the extremely high strain and high creep rates which these materials can exhibit. Materials with high creep rates and high strains are called superplastic. In 1934, Pearson observed elongations of up to 2000 % in lead-tin and bismuth-tin alloys [132]. He postulated that intergranular displacement, or grain boundary sliding, is the major factor for superplasticity. This discovery started a search for more and more fine grained materials to enhance the creep rate for manufacturing complex shaped parts. First, this behavior was researched using other metallic materials such as palladium, copper and other metals as well as on intermetallic compounds, like TiAl, Nb3Sn and FeMoSiB [120]. All these metallic m aterials s ho wed s up erp las ticit y at room temperature or s light ly above [119,120,132,133,134,135]. Later, the search for superplastic deformable ceramics led many researchers to the development of fine grained ceramics. Unfortunately, so far true nanoscaled ceramics could only be produced with TiO2, ZnO and ZrO2 [4,6,11,22,23,26,121]. Nanoscaled titania showed deformation which qualified for plasticity, but could not meet the requirements for superplasticity which means a strain rate of 0.4 (0.04 for the titania) and an elongation of at least 100% [27]. Similar results were obtained for nanosized ZnO, which exhibited a strain rate sensivity of less than 0.03 [121]. All these measurements for the superplastic deformation had to be conducted at temperature higher than 1300oC. At lower temperatures all ceramic materials showed brittle behavior. The measurements on titania and ZnO also proved that the grain size plays also a crucial role for the plastic deformation of ceramics. The smaller the grains, the higher the strain rates and deformations [121]. Until now, superplasticity in ceramic materials could be found in a wide range of ceramic
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materials, but especially in zirconia and zirconia composites [4,6,120,133,136-142]. Superplastic deformation occurs only at high temperatures (1300oC and more) [4]. At 1400oC, Kajihara et al. were able to show that a 3Y-TZP with 5 vol % silica could be plastically deformed to an elongation of 1038 % [136]. This high elongation was possible by the distribution of amorphous silica along the grain boundaries. In pure zirconia, plastic deformation showed rarely more than 50 % strain. Chen et al. showed superplasticity in TZP, alumina, Si3N4 and TZP/alumina composites [137,138]. Carry et al. investigated superplastic deformation in BaTiO3, zirconia, alumina and TiO2 [139,140]. The mechanism for superplasticity in ceramics were identified by Siegel as increased grain boundary sliding in fine grained materials [120]. Chen et al. proposed grain boundary diffusion for the fine grained ceramics [137]. They reported that the presence of a second phase can help in the superplastic deformation (e.g. a liquid phase at the grain boundaries) or can help to reduce damage to the microstructure during deformation. In the case of nanocomposites, the addition of nanosized particles to ceramic matrices seems to inhibit creep as long as they do not form glassy phases [29,40,143]. Ohji et al. could showrd by TEM measurments, that the inclusion of SiC in alumina prevent grain boundary sliding and possible dislocation movement.
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1.3. Concluding Comments The intensive investigations on true nanoscaled, dense single phase ceramics with grain sizes smaller than 100 nm were not rewarded with success so far. When nanosized microstructures were accomplished, the densities achieved of the ceramics usually were below 90 % TD. As the mechanical properties of strucural ceramics strongly depend on the density of the samples, it is important to find ways to produce dense nanosized ceramics. Especially for common structural ceramics as alumina, zirconia and SiC this is of interest for further studies on the grain size dependence of the mechanical properties. Zirconia ’nanoceramics’ (grain sizes
