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Rapid Prototyping of Ceramic/Metal Composites
(Rapid Prototyping von Keramik/Metall-Verbundwerkstoffen)
Der Technischen Fakultät der
Universität Erlangen-Nürnberg
Zur Erlangung des Grades
DOKTOR-INGENIEUR
vorgelegt von
Wei Zhang
Erlangen - 2010
Als Dissertation genehmigt von
der Technischen Fakultät der
Universität Erlangen-Nürnberg
Tag der Einreichung: 12. Januar 2010
Tag der Promotion: 02. Juli 2010
Dekan: Prof. Dr. Reinhard German
Berichterstatter: Prof. Dr. Peter Greil
Prof. Dr. Erdmann Spiecker
- III -
ACKNOWLEDGEMENTS This Ph.D. work has been carried out from April 2006 to September 2009 in the Institute of Glass
and Ceramics, Department of Materials Science, University of Erlangen-Nuremberg. Financial
support of the German Research Foundation (Deutsche Forschungsgemeinschaft, DFG) is gratefully
acknowledged.
First of all I would like to thank my supervisor, Prof. Dr. Peter Greil, for giving me this opportunity
to work in his institute, for the very interesting project, for his useful discussions and strong
support, for his encouragement and great suggestions and teachings.
I also would like to thank my group leader Dr. Nahum Travitzky for his guidance and support, for
his good ideas and many fruitful discussions, for his time and help, for his extensive knowledge and
experience, and for his practical advices.
I would like to thank all organizers, professors and colleagues of the DFG-Gradate School 1229
“Stabile und metastabile Mehrphasensysteme bei hohen Anwendungstemperaturen” for many
useful discussions and suggestions.
I also thank all my colleagues of the Rapid-Prototyping group and other groups, and all the
technical staff in the Institute of Glass and Ceramics for the pleasant working climate and their
support.
I would like to thank Dr. Pavel Leiva-Ronda for Density functional theory calculations.
Special thanks to my parents, my parents in law, and my sister, for their everlasting support,
comprehension and love.
Last, I want to thank my wife, Chun, and my daughter, Yanwen, who have helped me through all
the hard moments.
- Contents -
- IV -
Contents
Contents
0.1 List of figures VII
0.2 List of tables XI
Chapter 1 Introduction 1
Chapter 2 Basic principles 3
2.1 Ceramic/metal composites 3
2.2 MAX phases 7
2.3 Rapid prototyping: Three-dimensional printing (3DP) 11
2.4 Reactive melt infiltration 14
Chapter 3 Experimental procedure 17
3.1 Raw materials and powder processing 17
3.2 3DP 19
3.3 Pyrolysis and sintering 21
3.4 Reactive melt infiltration 22
3.5 Hot pressing 23
3.6 Microstructure analysis 24
3.7 Property measurements 24
3.8 Thermodynamic calculations 28
3.9 Density functional theory (DFT) calculations 29
- Contents -
- V -
Chapter 4 Results 31
4.1 Nb-Al-O system 31
4.1.1 Microstructure of preforms 31
4.1.2 Wetting of Al melt on Nb2O5 and NbO2 36
4.1.3 Microstructure of reaction composites 40
4.2 Nb-Al-C system 41
4.2.1 Printed, CIP-ed and sintered Nb2AlC 41
4.2.2 Hot-pressed Nb2AlC 47
4.2.3 Thermal properties 50
4.2.4 Mechanical properties 50
4.3 Ti-Al-O-C system 56
4.3.1 Microstructure 56
4.3.2 Fracture behavior 58
Chapter 5 Discussion 64
5.1 3DP multistep processing of composites 64
5.1.1 3DP 64
5.1.2 Reaction and microstructure control 65
5.1.3 Wetting and infiltration 71
5.1.4 Surface finish and accuracy of 3DP 75
5.1.5 Comparison and application 77
5.2 Mechanical behavior of MAX phase composites 80
5.2.1 Deformation and damage mechanisms 80
- Contents -
- VI -
5.2.2 Quasi-plasticity 81
5.2.3 Crack propagation and structure modeling 83
Chapter 6 Summary and Conclusions 89
References 93
List of publications 113
- Inhaltsverzeichnis -
- IV -
Inhaltsverzeichnis
Inhaltsverzeichnis
0.1 Abbildungsverzeichnis VII
0.2 Tabellenverzeichnis XI
Kapitel 1 Einleitung 1
Kapitel 2 Grundlagen 3
2.1 Keramik-Metall-Verbundwerkstoffe 3
2.2 MAX Phasen 7
2.3 Rapid Prototyping: Dreidimensionales Drucken (3D-Drucken) 11
2.4 Reaktive Schmelzinfiltration 14
Kapitel 3 Experimentelle Durchführung 17
3.1 Rohstoffe und Pulveraufbereitung 17
3.2 3D-Drucken 19
3.3 Pyrolyse und Sintern 21
3.4 Reaktive Schmelzinfiltration 22
3.5 Heißpressen 23
3.6 Mikrostrukturanalyse 24
3.7 Eigenschaftsmessungen 24
3.8 Thermodynamische Berechnungen 28
3.9 Dichtefunktionaltheorie (DFT) - Berechnungen 29
- Inhaltsverzeichnis -
- V -
Kapitel 4 Ergebnisse 31
4.1 Nb-Al-O System 31
4.1.1 Mikrostruktur des Vorkörpers 31
4.1.2 Benetzung von Al auf Nb2O5/NbO2 36
4.1.3 Mikrostruktur der Verbundwerkstoffe 40
4.2 Nb-Al-C System 41
4.2.1 3D-gedrucktes und gesintertes Nb2AlC 41
4.2.2 Heißgepresstes Nb2AlC 47
4.2.3 Thermische Eigenschaften 50
4.2.4 Mechanische Eigenschaften 50
4.3 Ti-Al-O-C System 56
4.3.1 Mikrostruktur 56
4.3.2 Bruchverhalten 58
Kapitel 5 Diskussion 64
5.1 3D-Drucken-basiertes Multi-Step-Verfahren von Verbundwerkstoffen 64
5.1.1 3D-Drucken 64
5.1.2 Reaktion- und Mikrostrukturkontrolle 65
5.1.3 Benetzung und Infiltration 71
5.1.4 Oberflächenrauigkeit und Genauigkeit von 3D-Drucken 75
5.1.5 Vergleich und Anwendungen 77
5.2 Mechanisches Verhalten von MAX-Phasen verstärkten Verbundwerkstoffen 80
5.2.1 Verformung und Schadensmechanismen 80
- Inhaltsverzeichnis -
- VI -
5.2.2 Quasi-Plastizität 81
5.2.3 Rissausbreitung und Strukturmodellierung 83
Kapitel 6 Zusammenfassung 89
Literatur 93
Veröffentlichungen 113
- List of figures -
- VII -
0.1 List of figures
2.1 Schematic showing the three main types of metal matrix composite [Cly93, Ncn].
2.2 Schema of Ceramic/Metal composites.
2.3 Schema of IPCs (left) [Ven06]; Microstructure of NiAl(Si)/Al2O3 composite prepared by a
metal infiltration of molten Al and Ni into silica preform (right) [Man08].
2.4 Elements of MAX phases summarized in periodic table [Bar01].
2.5 Unit cells of 211, 312 and 413 phases [Bar01].
2.6 Schematic of 3DP process.
2.7 Schematic of build bay: Machine directions and sample orientations [Zha09].
2.8 Effect of layer thickness on porosity and Young’s modulus of 3D-printed and sintered alumina
samples [Zha09].
2.9 Effect of Gibbs free energy per unit area by the reaction between liquid metal and ceramic
substrate on the contact angle [Aks74, Eus98].
3.1 Schema of 3DP [Yin06].
3.2 STL data designed by a CAD-program (Solid Edge Version 20, Siemens Product Lifecycle
Management (PLM) Software GmbH, Köln, Germany).
3.3 Turbine blade printed from powder blend TiO2/TiC/Dextrin (top) and metal blade (bottom).
3.4 Experimental set-up for investigation of wetting behavior between Al-foil and ceramic
substrate.
3.5 Schema of set-up for Young’s modulus measurement [Bos05].
3.6 Four-Point bending test set-up for for the in-situ investigation of crack propagation.
4.1 Particle size distribution of fabricated powders of CN1, CN2 and CNA.
4.2 Coral-like microstructure of CN1 preform uniaxially pressed (5 MPa) and sintered at 1400 °C
for 1 h 3.
4.3 Pore size distribution of pressed and sintered preforms.
4.4 CN1 preform printed and sintered at 1400 °C for 1 h.
4.5 Pore size distribution of printed, pyrolyzed and sintered preforms CN1.
4.6 Preforms printed and sintered at 1400 °C for 1 h of (a) CN2 and (b) CNA.
4.7 Pore size distribution of printed, pyrolyzed and sintered preforms CN2.
4.8 Pore size distribution of printed, pyrolyzed and sintered preforms CNA.
- List of figures -
- VIII -
4.9 Photographs of Al/Nb2O5 samples during the wetting experiments (wetting angle) at different
temperatures.
4.10 Photographs of Al wetting experiment on NbO2 at different temperatures.
4.11 Variations in the wetting angle of the molten Al on Nb2O5 and NbO2 between 700 °C and
1300 °C.
4.12 Variation in the wetting angle of the molten Al on Nb2O5 and NbO2 with time at temperature
of 1200 °C.
4.13 SEM micrograph of the polished cross section of pressed and infiltrated CN1 composite (light
gray: NbAl3, dark gray: Al2O3) 3.
4.14 SEM micrograph of the polished cross section of printed and infiltrated CAN (Light gray:
NbAl3, dark gray: Al2O3 and residual Al.
4.15 Particle size distribution of powder mixture NNA1 for 3DP.
4.16 Position of printed samples in build bay of 3D-printer.
4.17 Effect of position on the density of the printed Nb-Al-C samples.
4.18 Facture surface of prepared Nb-Al-C samples: (a) printed; (b) printed and CIP-ed at a pressure
of 200 MPa.
4.19 Pore size distribution of printed and CIP-ed Nb-Al-C samples.
4.20 Effect of applied CIP pressure on the linear shrinkage of Nb-Al-C samples.
4.21 Effect of applied CIP pressure on the volume shrinkage of Nb-Al-C samples.
4.22 Effect of applied CIP pressure on the density of Nb-Al-C samples.
4.23 XRD patterns for the surface and the center of the sintered sample.
4.24 Pore size distribution of reactive sintered Nb-Al-C samples prepared at different CIP pressure.
4.25 SEM micrograph of Nb2AlC sample etched cross section after hot-pressing at 1650 °C under a
pressure of 30 MPa for 90 min 2.
4.26 Temperature dependence of thermal expansions of Nb2AlC.
4.27 Dependence of Vickers hardness of Nb2AlC on the applied load 2.
4.28 Indent morphology of Nb2AlC after indentation with load of 100 N.
4.29 Indentation load dependence of residual bending strength of Nb2AlC 2.
4.30 Fracture surface of Nb2AlC tested in four-point bending after Vickers indentation with 300 N
load 2.
4.31 Fracture surface showing delamination and laminate kinking of the Nb2AlC grains 2.
4.32 Effect of quenching temperature on the bending strength of Nb2AlC 2.
- List of figures -
- IX -
4.33 Typical microstructure of Ti3AlC2/Al2O3/TiAl3 composite. 1: Ti3AlC2, 2: Al2O3, 3: TiAl3.
4.34 Stress-strain curve of Ti3AlC2/Al2O3/TiAl3 composite in four-point bending test.
4.35 SEM micrograph of crack propagation of the Ti3AlC2/Al2O3/TiAl3 specimens (The arrow
indicates the direction of crack propagation).
4.36 In-situ fracture series for SEVNB specimen of Ti3AlC2/Al2O3/TiAl3 composite in four-point
bending test.
4.37 Layered crystal structure of Ti3AlC2 (left); bonding charge density of Al/Ti terminated (10ī0)
surfaces in Ti3AlC2 (right) (Red, green and blue color means high, middle and low electron
density, respectively. Bonding charge calculation provided by Dr. Pavel Leiva-Ronda).
5.1 (a) Partial pressure of CO for Eq. (1); and (b) phase stability diagram in the system
Nb2O5 - NbO2 - NbO - NbC - C associated with the reactions upon reduction 3 (calculated by
means of equiTherm Version 5.04i [Bar97-2]).
5.2 Microstructure of sintered CN1 preforms prepared under different processing: (a) printed; (b)
pressed.
5.3 Decrease in the packing density of printed powder bed as a result of binder-powder interaction.
Left portion of the micrograph shows the unprinted region while the other half shows lower
packing due to rearrangement of granules [Yoo96].
5.4 Interfacial microstructure for the sample of molten on Nb2O5 at 1200 °C for 1 h in vacuum (<
10 Pa).
5.5 Ternary Nb-Al-O phase diagram at 1100 °C [Zha94, Sch98-2, Sch00].
5.6 Differential thermal analysis results of a sintered TiO2/TiC preform infiltrated with Al melt
[Yin07-2].
5.7 Testing part for surface finish measurement [Mel09]: (A) data model; (B) sintered Al2O3; (C)
Cu-O-infiltrated.
5.8 Surface finish of testing parts in green, sintered and infiltrated state depending on different
planes (0°/45°/90°) [Mel09].
5.9 Complex geometry parts by 3DP: a Al2O3-based moulding dies [Rep04, Mel06]; b glass-
infiltrated half skull [Zha09]; c infiltrated turbine wheel [Zha09]; d glass-infiltrated jaw; e
macrocellular SiSiC [Sch10].
5.10 Schematic of deformation and damage mechanisms of Ti3AC2 (A: Al and Si) [Zha04].
5.11 Fracture surface of Ti3AlC2/Al2O3/TiAl3 composite after four-point bending test showing
fracture mechanism of buckling, delamination and cleavage fracture.
- List of figures -
- X -
5.12 Model of the formation of kink bands [Hes49, Bar99-1, Bar04].
5.13 Variation of fracture toughness as a function of crack extension (R curves) for Ti3AlC2
reinforced TiAl3/Al2O3 prepared at 1300 °C and 1400 °C [Yin07-1].
- List of tables -
- XI -
0.2 List of tables
2.1 Selection of reaction-based processing techniques to fabricate ceramic/metal composites
[Fah06].
2.2 Summary of 211, 312 and 413 MAX phase materials.
2.3 Physical and mechanical properties of Ti2AlC [Bar00-1, Bar00-2, Wan02-1, Hu08-1], Ti3SiC2
[Bar96, Rag99-2, Fin00], Ti3AlC2 [Tze00, Wan02-2, Bao04] and Ti4AlN3 [Bar00-1, Bar00-3,
Raw00, Pro00, Bar00-4].
3.1 Raw materials used to fabricate Nb-Al-O composites.
3.2 Compositions of prepared powders for Nb-Al-O system.
3.3 Compositions of prepared powders for Nb-Al-C system.
3.4 Thermodynamic data of system Nb2O5-NbO2-NbO-NbC-C at T= 1523 K used for calculations
with a software package 3 (equiTherm Version 5.04i [Bar97-2]).
3.5 Characterizing of MAX phase Ti3AlC2 in TiAl3/Al2O3 composites.
4.1 Summary of phase analysis of the Nb-Al-C samples hot-pressed at the temperature range of
600 −1650 °C 2.
4.2 Properties of Nb2AlC phase materials compared to isomorphous Cr, Ti and Ta MAX phases 2.
4.3 EDS taken from 1, 2 and 3 regions of Fig. 4.33.
4.4 DFT-calculated cleavage energies for Ti3AlC2.
5.1 Vapor pressure of aluminum [Hat84].
5.2 EDS results taken from 1, 2, 3 and 4 region of Fig. 5.4 (b).
5.3 Comparison of fracture toughness and bending strength of ceramic materials fabricated by
3DP and other technique.
5.4 Properties E, CTE and ν of TiAl3, Ti3AlC2 and Al2O3.
5.5 DFT-calculated cleavage energy of Ti3AC2 (A = Al, Si).
- 1 Introduction -
- 1 -
1 Introduction
Advanced ceramic materials and ceramic composites are considered to play a key role in future
lightweight components to be used in various fields such as automobile and aerospace industries.
Material researchers and manufacturers try to develop the ceramic materials which combine the best
properties and can be used for many different applications. Intermetallic/ceramic composites were
developed which combine the properties of metal such as high ductility and toughness with ceramic
properties such as high modulus and wear resistance, low density, good corrosion and oxidation
resistance. The material group of ternary carbides/nitrides, which are characterized by a unique
nano-layer microstructure and a general formula Mn+1AXn (or MAX), where n is 1, 2, or 3, M is an
early transition metal, A is an A-group element (mostly A and VA), and X is either C or N, offers
a high potential to make accessible engineering applications which require improved mechanical
performance. Fabrication of MAX based materials, however, is mainly based on hot-pressing
technique. An external pressure is required to accelerate solid state reaction between the powder
components. Thus, only simple shape components can be produced which limits the freedom of
shaping as well as the manufacture of products.
Rapid prototyping offers a wide range of forming technologies which can produce complex
shaped parts directly from Computer Aided Design (CAD) data. Three dimensional printing
(3DPTM) is a rapid prototyping technique that constructs parts by spreading powders in thin layers
and then subsequently binding it with appropriate additives. 3DP can be used to produce the objects
with complex geometry. A subsequent post-processing such as reactive melt infiltration can be
performed to fabricate dense composite materials. The advantages of combining 3DP and reactive
infiltration processing include the realization of complex shaped parts, the use of cheap raw
materials/precursors, low-temperature processing, the optimization of microstructure and properties
of precursors and corresponding dense materials.
The aim of this work is to explore and develop a novel processing chain for MAX-based
composites which involves 3DP of a porous preform and subsequent metal melt infiltration reaction.
Based on preliminary thermodynamic work and reaction studies the following systems were
selected: Nb-Al-O, Nb-Al-C and Ti-Al-O-C. The working plan focused on preprocessing, shaping
by 3DP, and post-processing to convert the shaped component into MAX phase composite.
- 1 Introduction -
- 2 -
Fundamental scientific questions addressed in the work include the wetting and infiltration process
and the liquid-solid reaction process. 3D-printed MAX-phase composites were evaluated with
special emphasis on the damage tolerance behavior triggered by local deformation mechanisms in
the nano-laminate structure.
- 2 Basic principles -
- 3 -
2 Basic principles
2.1 Ceramic/metal composites
Ceramic Matrix Composites (CMCs) and Metal Matrix Composites (MMCs), are being developed
for a number of high-temperature and high-performance applications in industrial, aerospace, and
energy conservation sectors [Tua92, Gau95, Cla96, Gar97, Sch97].
CMCs consist of reinforcing phases and a ceramic matrix to create composite materials with
new and enhanced properties [Fre98]. Reinforcing phases for CMCs include discontinuous phases
(particles, whiskers, or short fibers) and continuous fibers. CMCs offer improved mechanical
properties such as strength and toughness compared to the unreinforced ceramics. In addition,
electrical and thermal properties can be optimized using adequate reinforcing phases [Fre98].
Therefore, CMCs have a unique combination of properties such as low density, high temperature
strength and fracture toughness, high corrosion resistance, good damage tolerance and thermal
shock resistance [Kre08]. Due to these enhanced properties the applications of CMCs include:
cutting tools, wear components, space engines, thermal protection systems, industrial and nuclear
applications [Kre08]. Superior properties of CMCs can be attractive alternatives to traditional
structural materials such as monolithic ceramics, intermetallic compounds, titanium-aluminum
alloys, steels and nickel-based superalloys [Kre08]. The disadvantages such as the high
manufacturing costs, the lack of commercial processing methods, the high material costs and
difficult-to-repair, however, have limited the use of CMCs [Fre98].
MMCs combine reinforcing phases with a continuous metal matrix. MMCs can be classified
according to the type and the geometry of reinforcement: continuous reinforced composites
(filaments) and discontinuous reinforced composites (particles, whiskers, or short fibers), Fig. 2.1
[Cly93, Ncn]. The common ceramic reinforcements are alumina, silicon carbide, titanium boride,
boron and graphite [Mor-Cly]. Similar to CMCs, MMCs combine the properties of metal matrix
such as light weight, high thermal conductivity, ductility and toughness with ceramic properties
such as high modulus, strength and wear resistance [Mor-Cly]. Contrary to conventional metals,
steels or alloys MMCs exhibit the following advantages [Ncn]:
- 2 Basic principles -
- 4 -
Low density
Optimized thermal expansion coefficients / thermal conductivity
Good specific mechanical properties
Improved wear, fatigue, creep, corrosion and oxidation resistance
Dimensional stability
Fig. 2.1 Schematic showing the three main types of metal matrix composite [Cly93, Ncn].
Fahrenholtz [Fah06] hat summarized reactive-based processes to fabricate ceramic/metal
composites, Table 2.1. According to these reactive fabrication processes, the types of ceramic/metal
composites are summarized in Fig. 2.2. Fahrenholtz pointed out that Gibbs free energy changes and
thermodynamic compatibility of the reaction phases are the criteria to produce dense composite
materials using reactive-based processes [Fah06].
- 2 Basic principles -
- 5 -
Fig. 2.2 Scheme of Ceramic/Metal composites.
Table 2.1 Selection of reaction-based processing techniques to fabricate ceramic/metal composites
[Fah06].
Process Material systems Ref.
Reaction bonding (RB)
Directed metal oxidation (DIMOX)
Displacive compensation of porosity (DCP)
Alumina-aluminide alloys (AAA)
Co-continuous ceramic composites (C4)
Reactive metal penetration (RMP)
Reactive hot pressing (RHP)
Al2O3, Si3N4
Al2O3-Al
MgAl2O4-Fe/Ni/Al, ZrC-WC-W
Al2O3-Ni3Al, -NbAl, or –TiAl
Al2O3-Al
Al2O3-Al
Al2O3-Nb, Al2O3-MoSi2, ZrB2-SiC
[Hol94, Gau99, Ril89]
[New86]
[Rog99]
[Sch98-1]
[Bre94]
[Loe96]
[Fah00, Fah02, Zha00]
Ceramic/Metal Composites
Non-reaction composites
Reaction-based composites
Ceramic reinforced metal matrix composites
Metal infiltrated ceramic matrix
composites
Reaction bonding composites
Directed metal oxidation composites
Displacive compensation of porosity composites
Alumina-aluminide alloys composites
Co-continuous ceramic composites
Reactive metal penetration composites
Reactive hot pressing composites
- 2 Basic principles -
- 6 -
More recently, much interest has arisen in research on interpenetrating composites (IPCs)
[Cla92, Mat04, Dob07]. Contrary to conventional ceramic or metal matrix composites, both phases
in the IPCs form homogenous microstructure with continuous interpenetrating three-dimensional
networks [Kla98], Fig. 2.3. Similar to CMCs and MMCs, IPCs combine the properties of metal
with ceramic properties. In addition, their continuous interpenetrating networks lead to significantly
enhanced mechanical properties [Tra03]. In the last few years, alumina/aluminide alloy (3A)
interpenetrating composites have been developed and studied extensively, which offer an excellent
combination of properties for high-temperature structural and functional applications [Kla98,
Hor02]. Focusing on the fabrication process of 3A composites, one approach is the in-situ reactive
powder processing technique where metal oxides (e.g. Fe2O3, Nb2O5, TiO2, ZrO2, etc.), and
elemental metals (e.g. Al, Fe, Nb, Ti, Cr, etc.), are milled, pressed and sintered, resulting in the
formation of oxide/intermetallic composites [Gar97, Sch97, Sub98, Sch98-1, Sch00, Hor02, Tra03].
)()(6)()3()( 322 sMeAlsOAlnlAlxnsOMe xnn (2.1)
where Me is Fe, Nb, Ti etc., n is 2, 3, 4, 6 and x is 1/3, 1 and 3. Another method is the infiltration
technique, where a porous metal oxide is infiltrated by a liquid metal [Röd95, Sch98-2, Avr06,
Yin06].
Fig. 2.3 Scheme of IPCs microstructure (left) [Ven06]; Microstructure of NiAl(Si)/Al2O3 composite
prepared by a metal infiltration of molten Al and Ni into silica preform (right) [Man08].
10 µm
- 2 Basic principles -
- 7 -
2.2 MAX phases
Ceramic usually are characterized by a combination of ionic and covalent bonding, which leads to
typical “ceramic” properties such as high elastic modulus and hardness, high melting point, low
thermal expansion, and good chemical resistance. On the other hand, ceramics are also brittle and
not easily machinable. Due to a wide scatter of strength and a low Weibull modulus, application of
ceramics as engineering materials is limited. One class of ternary ceramics, however, provides a
combination of ionic, covalent and metallic bonds, which give a combination of unique properties
from metals and ceramics. The nano-layered ternary carbides and nitrides with the general formula
Mn+1AXn (abbr. MAX), where n = 1, 2, or 3, M is an early transition metal, A is an A-group element
(mostly IIIA and IVA), and X is either C or N, represent a new class of solids [Sch80, Bar97-1,
Bar00-1, Bar01, Bar04], Fig. 2.4. MAX phases are layered hexagonal with space group of P63/mmc.
Fig. 2.5 [Bar01] shows the unit cells of the 211, 312 and 413 phases. The unit cell is characterized
by near close-packed M layers interleaved with layers of A-group element, with the X-atoms filling
the octahedral sites between the former [Bar00-1].
To date, more than 50 M2AX phases [Now71, Bar00-1], 5 M3AX2 phases [Jei67, Wol67, Pie94,
Dub07, Etz07], and 5 M4AX3 [Raw00, Etz07, Pal04, Hög05-1, Hu07-1, Hu07-2] were reported in
literature, Table 2.2.
Fig. 2.4 Elements of MAX phases summarized in periodic table [Bar01].
- 2 Basic principles -
- 8 -
Fig. 2.5 Unit cells of 211, 312 and 413 phases [Bar01].
Table 2.2 Summary of 211, 312 and 413 MAX phase materials.
Unit cells of MAX phases MAX phase materials
211
312
413
Ti2AlC, Nb2AlC, Ti2GeC, Zr2SnC, Hf2SnC, Ti2SnC
Nb2SnC, Zr2PbC, Ti2AlN, (Nb, Ti)2AlC, Cr2AlC, Ta2AlC
V2AlC, V2PC, Nb2PC, Ti2PbC, Hf2PbC, Ti2AlN0.5C0.5
Zr2SC, Ti2SC, Nb2SC, Hf2SC, Ti2GaC, V2GaC
Cr2GaC, Nb2GaC, Mo2GaC, Ta2GaC, Ti2GaN, Cr2GaN
V2GaN, V2GeC, V2AsC, Nb2AsC, Ti2CdC, Sc2InC
Ti2InC, Zr2InC, Nb2InC, Hf2InC, Ti2InN, Zr2InN
Hf2InN, Hf2SnN, Ti2TlC, Zr2TlC, Hf2TlC, Zr2TlN
Ti3SiC2, Ti3GeC2, Ti3AlC2, Ti3SnC2, Ta3AlC2
Ti4AlN3, Ta4AlC3, Ti4SiC3, Ti4GeC3, Nb4AlC3
- 2 Basic principles -
- 9 -
MAX phases exhibit an unusual combination of properties [Bar96, Rag99-1, Rag99-2, Bar99-1,
Bar99-2, Rad00, Rad02, Gil00, Rag00-1, Rag00-2, Bar00-1]. Like ceramics, they are stiff (high
elastic modulus ~ 300 GPa), have coefficients of thermal expansion in the range of 8 − 10 × 10-6 K-1,
and are resistant to chemical attack, oxidation and corrosion; like metals, they are highly damage
tolerant, thermal shock resistant, machinable, and relatively soft with Vickers hardness values of 2 –
5 GPa [Bar00-1, Bar04, Gup06-1]. They go through a ductile-brittle transition at temperatures >
1000 °C and retain mechanical properties at high temperature [Gup06-2, Bar04, Rad02, Rag99-2].
Physical and mechanical properties of typical MAX phase materials are summarized in Table 2.3.
In the last years many researchers have tried to produce MAX phase powders and bulk materials
by different methods such as chemical vapor deposition [Nic72, Got87, Rac94-1, Rac94-2, Fak06]
mechanically activated sintering [Li07-1], solid-state synthesis [Rac94-3], self-propagating
high-temperature synthesis (SHS) [Lis95, Kho02, Lis08], arc-melting and annealing [Aru95], solid-
liquid reaction process [Zho98, Sun99, Don01], hot isostatic pressing (HIP) [Gao02, Sal02, Gan04],
hot pressing (HP) [Luo02, Zhu04], pulse discharge sintering (PDS) [Zha02-1, Zha02-2, Zha01],
spark plasma sintering (SPS) [Zha07-1, Zho03], pressureless sintering [Tan02, Sun05, Has08]. In
addition, the deposition of MAX phase thin films has been achieved by magnetron sputtering from
corresponding ternary compound targets or individual elemental targets [Hög05-1, Hög05-2, Hög06,
Dub07, Pal02, Wal06].
- 2 Basic principles -
- 10 -
Table 2.3 Physical and mechanical properties of Ti2AlC [Bar00-1, Bar00-2, Wan02-1, Hu08-1],
Ti3SiC2 [Bar96, Rag99-2, Fin00], Ti3AlC2 [Tze00, Wan02-2, Bao04] and Ti4AlN3 [Bar00-1,
Bar00-3, Raw00, Pro00, Bar00-4].
Properties Ti2AlC Ti3AlC2 Ti3SiC2 Ti4AlN3
Density (g/cm3)
Coefficient of thermal expansion (× 10-6 K-1)
Thermal conductivity at 25 °C (W m-1 K-1)
Electrical conductivity (Ω-1 m-1)
Vickers hardness (GPa)
4-point Bending strength (MPa)
3-point Bending strength (MPa)
Compressive strength (MPa)
Fracture toughness (MPa m1/2)
Young’s modulus (GPa)
Shear modulus (GPa)
Brit. to duct. Trans. T (°C)
4.1
8.2
46
4.42 × 106
2.8
-
275
763
6.5
305
127
-
4.2
9.0
-
3.48 × 106
3.5
-
357±15
560±20
7.2
297
124
1050
4.5
9.2
43
4.5 × 106
4.0
465
-
885
6.9
333
139
1050
4.6
9.7
12
0.5 × 106
2.5
-
350±15
475±15
-
310±2
127±2
-
- 2 Basic principles -
- 11 -
2.3 Rapid prototyping: Three-dimensional printing (3DP)
Solid Freeform Fabrication (SFF) or Rapid Prototyping (RP) techniques can be defined as the
constructing of freeform solid objects directly from Computer-Aided Design (CAD) data without
the use of tooling, dies, or molds [Spr92, Ash94, Mar93, Pha98]. The working principle of major
RP techniques can be summarized as follows:
A CAD model is constructed using a CAD software package such as Solid Edge (Solid Edge
Version 20, Siemens Product Lifecycle Management (PLM) Software GmbH, Köln, Germany);
the CAD model is converted to Standard Triangular Language (STL) format;
RP device processes the STL file by creating sliced layers;
RP device constructs the model layer by layer;
Clean and finish the model.
Over the last two decades, several rapid prototyping techniques were developed. Laminated
Object Manufacturing (LOM) [Gri94] can be used to fabricate objects out of paper, plastic, metal
sheet stock, or ceramic tape. Selective Laser Sintering (SLS) uses a laser. Powder is spread in thin
layers and the laser energy is directed toward the surface of the layer to initiate localized sintering
[Dec87, Vai93]. Stereolithography also uses lasers to selectively cure resins in a laminated fashion
to build complex shapes [Ben89, Jac92]. For Fused deposition modeling (FDM), starting material in
a form of thermoplastic filament is fed to a heated dispenser. Thermoplastic is then molten and
extruded through the head. Three dimensional objects are made by controlling the movement of the
extrusion head to control the placement of the thermoplastic melt [Wal91].
Three-dimensional printing (3DPTM), an advanced RP technique, was first developed at the
Massachusetts Institute of Technology (MIT) [Sac92-1, Sac92-2, Sac93]. Fig. 2.6 shows the set-up
of 3DP process:
A powder feed roller spreads a layer of powder from the feed bay to cover the surface of the
build platform;
The print head then prints binder solution onto the powder causing the powder particles to bind
together, forming the first layer of the object.
- 2 Basic principles -
- 12 -
When the first layer is printed, the build platform is lowered slightly, and the printer spreads a
new layer. The process is repeated until the whole object is completed.
Fig. 2.6 Schematic of 3DP process.
The processing parameters of 3DP strongly influence the microstructure and mechanical
properties of the printed bodies. Moon et al. [Moo01] pointed out that the binder saturation, and the
interaction between binder drop and powder can have a significant influence on the surface finish
and microstructure of the printed preforms. In addition, the preforms printed with slower printing
speed show better surface structure compared to those with faster printing speeds [Moo01, Sun02].
Zhang et al. [Zha09] reported on the influence of layer thickness and sample orientation within the
build bay (Fig. 2.7) of the 3D-printer on microstructure, porosity and mechanical properties of the
printed objects. The increase of the layer thickness results in an increase of the total porosity of the
3D-printed and sintered alumina samples and thus, in a decrease of the mechanical properties of the
sintered preforms such as Young’s modulus, Fig. 2.8 [Zha09].
Build bay
Overfall
Feed bay
Powder bed
Print head
“Ink”
Powder feed roller
- 2 Basic principles -
- 13 -
Fig. 2.7 Schematic of build bay: Machine directions and sample orientations [Zha09].
Fig. 2.8 Effect of layer thickness on porosity and Young’s modulus of 3D-printed and sintered
alumina samples [Zha09].
Z-axis
Y-axis
X-axis
Z-orientated sample
X-orientated sample
Y-orientated sample
Piston
Gantry
Print head
- 2 Basic principles -
- 14 -
3DP was used to produce complex components with a wide range of material systems including
ceramics [Yoo93, Yoo98], glasses [Cim95], metals [Mic92], and polymers [Cim94]. In order to
fabricate dense composite materials, a combination processing from 3DP with reactive melt
infiltration has been developed and studied. It has been reported that complex-shaped TiC/Ti-Cu
[Ram05], Al2O3/Cu-O [Mel06], Si/SiC [Tra06], TiAl3/Al2O3 [Yin06], and Al2O3/glass [Zha09]
composites were fabricated by the combination processing (3DP with reactive melt infiltration).
Focusing on fabrication of MAX phase materials by 3DP, Sun et al. [Sun02, Dco02] have pointed
out a development using three-stage fabrication process, i.e., 3DP, cold isostatic pressing (CIP), and
sintering processing to freeform fabrication of three-dimensional Ti3SiC2 structures with complex
geometry and high density (> 99%). In addition, Yin et al. [Yin07-1, Yin07-2] reported on the
fabrication of Ti3AlC2/Al2O3/TiAl3 composites by 3DP and pressureless melt infiltration.
2.4 Reactive melt infiltration
The infiltration of molten metals or alloys into a porous ceramic preform is a processing route to
fabricate ceramic/metal composites [Hil87, Hil88, Tra98-1]. Pressureless melt infiltration is
versatile and offers near net shape capability [Toy90, Rit93, Tra97, Tra98-2, Gre99, Gre02]. A
critical wetting angle much smaller than 90° is required to achieve pressureless infiltration driven
by capillary force [Yan95]. The characteristics of ceramic preforms such as pore size and pore
shape can have a strong influence on the wettability of ceramic preform by metal melt [Hil88]. The
infiltrant composition and infiltration parameters such as temperature, time and atmosphere can also
affect the wettability and the infiltration height [She06]. When the molten infiltrant reacts with the
ceramic preform, the reaction kinetics at the liquid-solid interface can have a strong effect on their
wettability [Sin95, Ast00].
In the absence of a chemical reaction between liquid metal and ceramic substrate, the contact
angle can be described on the basis of Young’s equation,
LV
SLSV
cos (2.2)
- 2 Basic principles -
- 15 -
where σSV, σSL, and σLV are the interfacial tensions at the boundaries between solid (S), liquid (L),
and vapor (V). When a reaction occurs at the interface, the decreased contact angle θmin according to
Laurent Lau88 is given by
LV
r
LV
r tGt
coscos min (2.3)
in which the last two terms represent the time (t) dependent of the chemical reaction on wettability
[Kal95, Eus98]. The terms Δσr and ΔGr are the change in interfacial energy and in Gibbs free
energy per unit area by the reaction, respectively. Aksay et al. and Naidich claimed that ΔGr is the
predominant factor for reactive wetting Aks74, Nai81, Nai83. Furthermore, during the early stage
of interfacial reaction, initial contact angle θinit reaches a minimum θm duo to maximum change in
Gibbs energy after the time of tm; thereafter, the reaction kinetics slow down, the contact angle
increases again and gradually approaches an equilibrium value θe in the time of te , Fig. 2.9 [Aks74].
Fig. 2.9 Effect of Gibbs free energy per unit area by the reaction between liquid metal and ceramic
substrate on the contact angle [Aks74, Eus98].
- 2 Basic principles -
- 16 -
The infiltration height of the metal melt into the porous ceramic preform, h, can be described
according to Darcy’s law:
21
2
t
PKh
p (2.4)
where K is the permeability of the preform; P is the pressure drop in the metal melt; t is the
infiltration time; is the viscosity of the metal melt; p is the pore volume fraction [Mol05, Piñ08,
Yin06]. According to Kozeny-Carman equation the permeability of the preform, K, can be
described [Yin06]:
223
15.37 p
p rK
(2.5)
where r is the particle radius.
For pressureless melt infiltration, the pressure drop, P, can be described [Wan05, Yin06]:
cPPP (2.6)
cos13
p
pc r
P (2.7)
where P is external pressure; Pc is capillary pressure.
Combining Eqs. (2.4), (2.5), (2.6) and (2.7), the infiltration height is given as [Yin06]:
21
cos25.61
r
thp
p (2.8)
where λ is the porosity shape factor, and γ is the surface tension. Thus, accelerated infiltration could
be achieved with the increase of the time, surface tension and porosity shape factor and with the
decrease of the wetting angle and viscosity of the melt.
- 3 Experimental procedure -
- 17 -
3 Experimental procedure
3.1 Raw materials and powder processing
Nb-Al-O system
The raw materials, which were used to fabricate Nb-Al-O composites, are shown in Table 3.1.
Table 3.1 Raw materials used to fabricate Nb-Al-O composites.
Materials Mean particle size
D50 (µm) Supplier
Nb2O5 CERAMIC GRADE
Flammruß 101
α-Al2O3, CT 3000 SG
Dextrin, Gelb mittel F
Ammonium polymethacrylate, Darvan C
Al-foil
0.6
0.095
0.8
150
-
0.3 mm thick
H.C. Starck , Goslar, Germany
Degussa, Hanau-Wolfgang, Germany
Almatis, Ludwigshafen, Germany
Südstärke GmbH, Schrobenhausen, Germany
R.T. Vanderbilt Company, USA
Merck, Darmstadt, Germany
Powder blends were prepared by mixing Nb2O5, α-Al2O3 and C powder with Dextrin (C6H10O5)n
(n = 10 – 200) powder. Dextrin was used as a binder in order to enhance the green strength of 3D-
printed objects. A polymethacrylate-based dispersant agent (Darvan C) can be used for the mixing
and suspension. Dextrin (6 wt. %) decomposed to amorphous carbon (1.57 wt. %) upon pyrolysis.
Three powder blends with different compositions labeled as CN1, CN2 and CNA were prepared,
Table 3.2. Each mixture was tumbled in a polyethylene bottle with Al2O3 grinding balls for 48 h
(Reax 20, Heidolph, Schwabach, Germany). Each slurry was freeze-dried at 50 °C / 37 Pa (Delta 2-
24, Christ, Osterode/Harz, Germany). Each dried batch was jar-milled for 72 h and sieved through
150 μm mesh.
Table 3.2 Compositions of prepared powders for Nb-Al-O system.
Samples Compositions of powder (wt. %) Molar ration
(C/Nb2O5) Nb2O5 α-Al2O3 Carbon black Dextrin
CN1
CN2
CNA
91.44
87.66
69
0
0
23
2.56
6.34
3
6
6
5
1
2
1.38
- 3 Experimental procedure -
- 18 -
Nb-Al-C system
Two types of powder blends were prepared by mixing NbC (d50 ~ 0.9 µm, NIOBIUM CARBIDE
HGS, H.C.Starck, Goslar, Germany), Nb (5 − 45 µm, AMPERIT® 161.3, 99.9 % purity,
H.C. Starck, Goslar, Germany), Al ( 45 µm, 99.5 % purity, Eckart-Werke, Fürth, Germany) and
Dextrin, Table 3.3. Each powder mixture was milled (Reax 20, Heidolph, Schwabach, Germany)
with Al2O3 grinding balls for 24 h. After evaporation of the acetone, the milled mixture was passed
through a 200 µm sieve.
Table 3.3 Compositions of prepared powders for Nb-Al-C system.
Samples Compositions of powder (wt. %) Molar ratio
(NbC/Nb/Al) NbC Nb Al Dextrin
NNA1
NNA2
36.5
46.7
46.1
41.3
13.4
12
4
0
0.7/1/1
1/1/1
Ti-Al-O-C system
The powder blend was prepared by mixing TiC powder (d50 ~ 1.2 µm, H.C. Starck, grade HV
120, Germany) and nanometer TiO2 powder (d50 ~ 30 nm, Degussa P 25, Hanau, Germany) with
Dextrin (C6H10O5) n (n = 10 – 200) powder (d50 ~ 115 µm, Superior Gelb mittel F, Suedstaerke
GmbH, Schrobenhausen, Germany). The weight ratios of TiC, TiO2 and dextrin in the powder blend
were 56.4 wt. %, 37.6 wt. % and 6 wt. %, respectively. Slurry mixing was carried out in aqueous
suspension containing a polymethacrylate-based dispersant agent (Darvan C, R.T. Vanderbilt,
Norwalk/CT, USA). The slurry was tumbled in a polyethylene bottle with Al2O3 grinding balls for
48 h (Reax 20, Heidolph, Schwabach, Germany) and then freeze-dried at 50 °C / 37 Pa (Delta 2-24,
Christ, Osterode/Harz, Germany). The dried batch was jar-milled for 72 h and sieved through 200
μm mesh.
- 3 Experimental procedure -
- 19 -
3.2 3DP
The green parts were designed using a conventional CAD-program (Solid Edge Version 20,
Siemens Product Lifecycle Management (PLM) Software GmbH, Köln, Germany). 3DP was
carried out in a 3D printer (ZPrinter 310, Zcorporation, Burlington, MA, USA). This 3D printer is
allowed to print layer thickness in the range of 88 to 225 µm. A building up speed of ~ 20 mm
thickness per hour could be achieved, Fig. 3.1 [Yin06]. Water-based printing solution was passed
through a bubble jet print head (nozzle diameter ~ 60 µm, number of nozzles ~ 304). Dextrin could
be dissolved by the printing solution and the powder particles became bonded together, providing
mechanical integrity during the 3DP process. The amount and distribution of injected printing
solution depend on the mass flow rate (3.7 cm3/h, when a cylinder with a diameter of 20 mm and a
thickness of 18 mm as reference sample was printed) and the printer head velocity (15 cm/s). In the
present work, the thickness of an individual layer was set to 90 µm and the binder saturation was
kept constant at 0.35 g/cm3. For Nb-Al-O system, rectangular plates of 50 × 50 × 6 mm3 were
printed. For reference, samples were uniaxially pressed applying a pressure of 5 MPa and 10 MPa,
respectively. For Nb-Al-C system, rectangular plates of 9 × 8 × 7 mm3 were printed using the
powder blend NNA1 (Table 3.3). In order to fabricate Ti-Al-C-O composites, the prepared powder
blend TiO2/TiC/Dextrin was used to print rectangular plates of 49 × 49 × 7 mm3.
In order to prove the capability to fabricate components with complex geometry a turbine blade
was designed and printed as a demonstration component, Fig. 3.2 and 3.3. After drying in air at
room temperature for 48 h, the printed parts were removed and cleaned from the unbound powder
bed. Because 3DP is a powder-based process where particles are glued together by a binder fluid,
the printed parts have high open porosity and low strength. After sintering the printed parts are
porous with an open porosity in the range of ~ 20 % – 40 %, Fig. 2.7 [Zha09]. Duo to the porous
structure and low strength of printed samples, further post-processing such as cold isostatic pressing
(CIP), pyrolysis, sintering and melt Infiltration can be performed to achieve dense materials with
enhanced properties.
- 3 Experimental procedure -
- 20 -
Fig. 3.1 Scheme of 3DP [Yin06].
Fig. 3.2 STL data designed by a CAD-program (Solid Edge Version 20, Siemens Product Lifecycle
Management (PLM) Software GmbH, Köln, Germany).
X
Y Z
Print head
Single layer thickness, 90 m
Build-up direction
2 cm
- 3 Experimental procedure -
- 21 -
Fig. 3.3 Turbine blade printed from powder blend TiO2/TiC/Dextrin (top) and metal blade (bottom).
3.3 Pyrolysis and sintering
Nb-Al-O system
The printed as well as pressed preforms were pyrolyzed at 800 ºC for 2 h in N2 atmosphere to
decompose the dextrin binder into carbon (C). The pyrolyzed preforms were sintered at 1400 °C for
1 h in Ar atmosphere. Carbon serves as a reduction agent to reduce Nb2O5 at least on the surface
above 1150 °C:
Nb2O5 (s) + C (s) 2 NbO2 (s) + CO (g) (3.1)
The reduced NbO2 preforms with interconnected porosity provide a lower wetting angle for Al-
melt and hence facilitates complete infiltration. After sintering the mechanical stability of the
sintered samples increased when compared with the 3D-printed samples.
- 3 Experimental procedure -
- 22 -
Nb-Al-C system
3D-printed samples of the Nb-Al-C system were additionally cold isostatically pressed (Loomis
Products Kahlefeld GmbH, Kaiserslautern, Germany) with pressures of 50 MPa, 100 MPa, 150
MPa and 200 MPa for 60 s to reduce porosity. Debinding and sintering were accomplished in one
step in vacuum (~ 1 − 5 Pa). A multistep heating program was applied with hold steps at 250 °C (2
h), 600 °C (2 h) and 1450 °C (30 min). The heating rate was increased from 1 °C/min (25 – 600 °C)
via 2 °C/min (600 – 800 °C) to 5 °C/min (800 – 1450 °C). A cooling rate down to room temperature
of 5 °C/min was applied. The sintering temperature and holding time were selected according to
differential thermal/thermogravimetric analysis (DT/TGA) to identity corrective temperature region
of maximum decomposition rate.
Ti-Al-O-C system
Pyrolysis of the printed preforms was carried out at 800 ºC for 2 h in N2 atmosphere, followed by
sintered in flowing Ar at 1400 °C for 0.5 h.
3.4 Reactive melt infiltration
Al foil (see Table 3.1) was used for infiltration, which was placed on the top and bottom of the
sintered samples. The impurity composition of the Al foil given by supplier was: N ~ 0.005 wt. %,
As ~ 0.0002 wt. %, Cu ~ 0.005 wt. %, Fe ~ 0.006 wt. %, Mn ~ 0.002 wt. %, Si ~ 0.02 wt. %, Zn ~
0.005 wt. %. The specimen (Al foil and sintered sample) was placed in an alumina crucible. In
order to present the crucible sticking to the molten Al, the bottom of crucible was covered with a
coarse alumina powder (d50 ~ 14.7 µm). The reactive melt infiltration was carried out in two steps
under Ar: at 1200 °C for 1h and at 1400 °C for 1 h, in order to ensure complete infiltration and
reaction between the Al-melt and sintered preform.
Wetting of Al melt on porous Nb2O5 and NbO2 was investigated. 5 g of Nb2O5 and NbO2
powder was poured into a tool steel die, 15 mm in diameter, and uniaxially pressed by applying a
pressure of approximately 10 MPa, respectively. The specimen cylinders were encapsulated in an
elasto mold and cold isostatically pressed (CIP) with a high pressure of 200 MPa (Loomis Products
Kahlefeld GmbH, Kaiserslautern, Germany). Finally, the samples were annealed up to 1400 °C for
1 h in Ar. After polishing using diamond pastes down to about 1 μm the specimens were
- 3 Experimental procedure -
- 23 -
ultrasonically cleaned in acetone. The set-up for investigation of wetting behavior between Al-foil
and ceramic substrate is shown in Fig. 3.4. The Al-covered sample was heated (Astra Model 1100-
4080-MI, Thermal Technology LLC, Santa Rosa, USA) up to 1200 °C and 1400 °C for 1 h under
vacuum (< 10 Pa). Melting and interfacial behavior were recorded by a video camera. The wetting
angle between Al and niobium oxide was evaluated from the photographs by a graphics program
package (CorelDRAW Graphics Suite 12, Version 12.0.0.458, (C) 2003 Corel Corporation).
Fig. 3.4 Experimental set-up for investigation of wetting behavior between Al-foil and ceramic
substrate.
3.5 Hot pressing
For reference, the powder blend NNA2 was uniaxially pressed at 5 MPa in a BN-coated graphite
die. In order to study the reaction mechanisms of the powder mixture NNA2, one set of the samples
was heated to 600, 700, 900, 1100, 1300, 1500 and 1650 °C in Ar atmosphere, respectively. The
heating rate was 15 °C/min and the holding time was 30 min under a pressure of 1 MPa. The
cooling rate down to room temperature was 15 °C/min. Another set was heated to 1650 °C for 90
min under 30 MPa in Ar atmosphere. Discs, 50 mm in diameter and 8 mm in hight, can be achieved
via hot pressing.
Al foil
Ceramic substratVideo camera Light source
Furnace
- 3 Experimental procedure -
- 24 -
3.6 Microstructure analysis
Phase composition of the fabricated composites was analyzed by X-ray diffraction (XRD,
Kristalloflex D 500, Siemens, Karlsruhe, Germany) using monochromatic Cu Kα radiation ( =
1.54178 Å) at a scan rate 0.75° min-1 over a 2 theta range of 5 − 70°. The polished and the fractured
surfaces of the composites were analyzed by a scanning electron microscope (Quanta 200, FEI,
Praha, Czechia) equipped with an energy-dispersive X-ray spectroscope (EDS, Inca x-sight, Oxford
Instr., Oxford, UK). The chemical composition in composites was determined by an inductively
coupled plasma optical emission spectrometer (ICP-OES, Spectro Flame Modula, Spectro
Analytical Instruments, Kleve, Germany) using powdered samples. The crystallographic orientation
of grains was investigated by electron backscattering diffraction (EBSD) in a SEM (Carl Zeiss 1540
crossbeam system, Germany).
3.7 Property measurements
Particle size distribution
A laser diffractometer (Mastersizer 2000, Malvern Instruments Ltd., Malvern, UK) was used for the
particle size distribution measurements of powders. The instrument was equipped with a Scirocco
2000 dry powder dispenser for prepared granulate powders. Raw materials was measured being
suspended in distilled water (Hydro 2000S). The diagram of volume percent versus the particle size
was determined together with the particle diameters of d10, d50 and d90.
Density, porosity and pore size distribution
The geometrical density ρg of porous preforms was determined by measuring the dimensions and
the weight of the samples. The skeleton density ρs was determined by He-pycnometry (Accu Pyk
1330, Micromeritics Inc., USA). The relative density ρr of the sample was calculated from the ρg
and ρs:
s
gr
(3.2)
- 3 Experimental procedure -
- 25 -
The total open porosity and pore size distribution of porous preforms was measured by Hg-
porosimetry (Pascal 140, Thermo Electron, Rodano/Milan, Italy). The density of infiltrated
composites was measured by Archimedes method.
Thermal properties
The coefficient of thermal expansion (CTE) was measured in a dilatometer (Dil-402E, Netzsch,
Selb, Germany) under flowing Ar atmosphere from room temperature to 1050 °C at a heating rate
of 5 °C/min. The specimen size is ~ 4.9 × 3.9 × 2.7 mm3. CTE was calculated as:
dTl
dlCTE
1 (3.3)
where l denotes the length of the measured sample, dl the length difference, and dT the temperature
range of the measurement.
The thermal diffusivity α was measured on carbon-coated cylindrical samples with a diameter of
15 mm and a thickness of 0.3 mm (Thermal Pulse System XP20, CompoTherm Messtechnik, Syke,
Germany). The thermal conductivity, λ (W (m K)-1), was calculated from the thermal diffusivity
[Tia06]:
pc (3.4)
where ρ is the density and cp is the specific heat capacity (382.6 J (kg K)-1 [Bar02-1]).
Mechanical properties
Bending strength of prepared samples with dimensions of 3 × 4 × 36 mm3 was measured by four-
point bending method using a universal testing device (Instron 4204, Instron Corp., Canton, MA,
USA) according to [DIN95]. Specimens were machined by diamond cutting as well as electrical
discharge machining (EDM). The tensile surfaces of the sample bars were polished to 1 µm
diamond finish prior to bending. For reference, another set of bars was unpolished. The crosshead
speed and corresponding strain rate for bending strength tests was 0.5 mm/min and ~ 4.5 × 10-3 s-1,
respectively. The inner and outer span was 10 mm (ls) and 20 mm (Ls), respectively. In order to
- 3 Experimental procedure -
- 26 -
compare with literature data, a three-point bending method with a span length of 30 mm was
performed to measure the bending strength. Four-point bending strength σB was calculated as
22
3
hb
lLF ssB
(3.5)
where F is the load, b is the specimen width, and h is the specimen thickness. A mean value of
bending strength was calculated by averaging over ten measurements. The fracture toughness was
measured by four-point bending method with spans of 10 mm and 20 mm (Instron 4204, Instron
Corp., Canton, MA, USA). The crosshead speed and corresponding strain rate for fracture
toughness tests was 0.05 mm/min and ~ 4.5 × 10-4 s-1, respectively. The fracture toughness was
determined by means of ‘single-edge-v-notch-beams’ (SEVNB) [DIN91, Küb02]: bar specimens
with a v-notch of finite width were introduced by a saw cut and sharpened by a razor-blade; the
fracture toughness value was derived from the maximum load and the dimensions of the specimen
and the notch; the depth of the notch a was optically measured with a CCD camera (Leica DC200,
Leica, Heerbrugg, Schweiz); ten specimens were tested for one data point; fracture toughness, KIC,
was calculated from the fracture load Fmax by:
Yhb
FKIC
max (3.6)
where b is specimen width, and h is specimen height; Y was calculated by the following formula:
2
2
1
135.168.049.3326.19887.1
12
32
3
e
eeeee
e
e
h
lLY ss (3.7)
where Ls is outer span, ls is inner span, and e is:
h
ae (3.8)
- 3 Experimental procedure -
- 27 -
Young’s modulus was measured using the impulse excitation technique (Buzz-o-sonic,
BuzzMac Software LLC, Glendale, WI, USA) [Rad04]. The bar sample was placed on polymer
foams; a small impulse tool (3−4 mm ball bearing cemented to a flexible plastic strip) struck the bar
specimen and created a standing wave; the frequency of vibration was measured using a
microphone; together with the dimensions and mass of the specimen the elastic constants can be
calculated, Fig. 3.5 [Bos05].
Fig. 3.5 Scheme of set-up for Young’s modulus measurement [Bos05].
The Vickers hardness was measured at indentation loads of 1, 3, 5, 10, 30, 50, 100, 200 and 300
N with a dwell time of 10 s (Zwick 3212, Zwick, Ulm, Germany). The average hardness values
were determined from twenty indentation measurements for each load. In order to analyze damage
tolerance of prepared samples the residual bending strength after indentation was measured by four-
point bending method [Pro00]. Thermal shock resistance was determined by water quenching
method: the prepared samples were heated to 600, 800, and 1000 °C for 10 min in Ar atmosphere,
respectively, and then immediately quenched into a room temperature water bath. The retained
bending strength after water quenching was measured by four-point bending method.
SEVNB specimens with dimensions of ~ 2 × 4 × 18 mm3 were used for the in-situ investigation
of crack propagation by four-point bending tests using a tensile/compression apparatus (maximum
loading 5000 N, Kammrath & Weiss GmbH, Dortmund, Germany) with a crosshead speed of 0.5
µm/s, Fig. 3.6. The load-displacement curves were monitored and the crack propagation path in the
specimen during the loading was examined in a scanning electron microscope (SEM, JSM-6400,
Jeol, Japan).
- 3 Experimental procedure -
- 28 -
Fig. 3.6 Four-point bending test set-up for the in-situ investigation of crack propagation.
3.8 Thermodynamic calculations
A thermodynamic calculation software package (equiTherm Version 5.04i) [Bar97-2] was used to
calculate partial pressure of CO and phase stability diagrams in the system Nb2O5-NbO2-NbO-NbC-
C. This software can be used to calculate equilibrium compositions by minimizing the Gibbs energy
at constant pressure (or volume) and constant temperature; the data including the standard enthalpy
change, the absolute entropy, the heat capacity of formation of a compound or reaction can be used
for calculations [Bar97-2]. Gaseous CO and five condensed phases in the system Nb-O-C were
considered in the calculations, and the basic thermodynamic data are summarized in Table 3.4.
2 cm
- 3 Experimental procedure -
- 29 -
Table 3.4 Thermodynamic data of system Nb2O5-NbO2-NbO-NbC-C at T= 1523 K used for
calculations with a software package 3 (equiTherm Version 5.04i [Bar97-2]).
Species ΔG (kJ/mol)
Nb2O5 -2305
NbO2 -968
NbO -548
NbC -250
C -28
3.9 Density functional theory (DFT) calculations
In order to characterize the fracture and crack propagation behavior on the level grains the cleavage
energy was calculated in layered crystal structure of Ti3AlC2 by means of ab initio density
functional theory (DFT) [Hoh64, Koh65]. In this work, density functional theory calculations were
carried out within the cooperation project of simulation of MAX phases with the institute of general
materials properties, University Erlangen. The cleavage energy Gc is defined as the energy
difference between the sum of the two fractured surfaces energies (EFS-1 and EFS-2) and the total
energy of the corresponding bulk supercell (EBS), normalized by the total fractured surface area A
[Zha07-2]:
A
EEEG BSFSFS
c
21 (3.9)
Total energies EBS were calculated using exchange-correlation effects described by the Perdew-
Burke-Ernzerhof generalized-gradient-approximation (PBE-GGA) functional [Per96], as it is
implemented in the ABINIT open source program [Gon02]. The electron wave function was
expanded in a plane wave basis set (Energy cut-off of 35 Hartree) and the core-valence interaction
was modeled by Goedecker, Teter and Hutter (GTH) norm-conserving pseudopotentials [Kra05].
Brillouin-zone integrations were performed using Monkhorst-Pack [Mon76] k-point meshes with a
- 3 Experimental procedure -
- 30 -
density corresponding to 6 × 6 × 2 k-points for the bulk unit cell and accordingly less in larger
supercells. Furthermore, the bonding charge was calculated as the charge density difference
between Ti3AlC2 and the superposition of neutral Ti, Al and C atomic densities at the corresponding
lattices sites. This bonding charge distribution was also displayed using a software package
XCrySDen [Kok03]. The basic data of Ti3AlC2, TiAl3 and Al2O3 are shown in Table 3.5.
Table 3.5 Characterizing of MAX phase Ti3AlC2 in TiAl3/Al2O3 composites.
.
Properties TiAl3
[Nak91, Mil01]
Al2O3
[Zhu98]
Ti3AlC2
[Bar00-1, Wan02-2]
Crystallography Tetragonal Hexagonal Hexagonal
Density (g/cm3) 3.3 3.98 4.5
Lattice parameter (Å) a=3.863
b=8.587
a=4.760
b= 13.000
a=3.075
c=18.578
Young’s modulus (GPa) 216 [Nak91] 430 [Zhu98] 297
Vickers Hardness (GPa) 5 15 3.5
Bending strength (MPa) 162 380 340
Fracture toughness (MPa m1/2) 2 3.5 7.2
CTE (× 10-6 K-1) 13 8.3 9
- 4 Results -
- 31 -
4 Results
4.1 Nb-Al-O system
4.1.1 Microstructure of preforms
Fig. 4.1 shows the particle size distribution of fabricated Nb2O5 powders. While Al2O3 free powder
mixtures CN1 and CN2 display a broad particle size distribution, addition of Al2O3 in CNA resulted
in a narrow particle size distribution. After annealing to 1150 °C the carbon caused reduction of
Nb2O5:
CONbOCONb 252 2 (4.1)
Pressed and sintered sample CN1 exhibits a porous “coral-like” microstructure consisting of
interconnected pores, Fig. 4.2: the open porosity and average pore diameter of sintered CN1 is 63 %
and 1.8 µm, respectively; the mean grain size and bulk density of sintered CN1 is 2 µm and 1.7
g/cm3, respectively. In the pressed and sintered sample CN2 the carbothermal reduction can be
described according to Eq. (4.2):
Nb2O5 (s) + 2 C (s) (28/17) NbO2 (s) + (1/17) Nb6C5 (s) + (29/17) CO (g) (4.2)
The pressed and sintered samples exhibit monomodal pore distribution: the pore diameters are
estimated to be between 1 and 2 µm with a narrow distribution as derived from Hg-porosimetry
measurements, Fig. 4.3. The open porosity of sintered CN2 preforms as-pressed at 5 MPa and 10
MPa are 63 % and 64 %, respectively. The increase from 5 MPa to 10 MPa in pressure resulted in
an increase in bulk density of CN2 preforms from 1.32 g/cm3 to 1.45 g/cm3.
- 4 Results -
- 32 -
Fig. 4.1 Particle size distribution of fabricated powders of CN1, CN2 and CNA.
Fig. 4.2 Coral-like microstructure of CN1 preform uniaxially pressed (5 MPa) and sintered at
1400 °C for 1 h 3.
10 µm
- 4 Results -
- 33 -
Fig. 4.3 Pore size distribution of pressed and sintered preforms.
Fig. 4.4 CN1 preform printed and sintered at 1400 °C for 1 h.
20 µm
- 4 Results -
- 34 -
Fig. 4.4 shows a typical microstructure of printed and sintered CN1 preform. Compared to the
pressed reference material connectivity of open pore channels is substantially reduced due to the
granulated structure of the print powder. While the pressed and sintered samples exhibit
monomodal pore distribution, the printed, pyrolyzed and sintered preforms CN1 offer a bimodal
pore size distribution with two maxima: first peak at 0.15 µm for printed samples, 0.18 for
pyrolyzed samples and 0.6 µm for sintered samples, respectively; second peak at 40 µm for all
samples, Fig. 4.5. After sintering, the volume fractions of the intra-agglomerate pores decreased and
that of the inter-agglomerate pores increased, Fig. 4.5. Fig. 4.6 shows the “coral-like”
microstructure of printed and sintered CN2 and CNA preforms. These results are similar to those
obtained in uniaxially pressed samples (Fig. 4.2). The printed CN2 and CNA samples also exhibit a
bimodal pore size distribution consisting of similar pore volume with respect to the pore size, Fig.
4.7 and Fig. 4.8. After pre-sintering, the volume fractions of the intra-agglomerate pores increased,
associated with an increase in pore size, and the large inter-agglomerate pores were reduced in
amount, but were not completely eliminated.
Fig. 4.5 Pore size distribution of printed, pyrolyzed and sintered preforms CN1.
- 4 Results -
- 35 -
Fig.4.6 Preforms printed and sintered at 1400 °C for 1 h of (a) CN2 and (b) CNA.
Fig. 4.7 Pore size distribution of printed, pyrolyzed and sintered preforms CN2.
10 µm 10 µm
(a) (b)
- 4 Results -
- 36 -
Fig. 4.8 Pore size distribution of printed, pyrolyzed and sintered preforms CNA.
4.1.2 Wetting of Al melt on Nb2O5 and NbO2
Figs. 4.9 and 4.10 show the photographs of Al/Nb2O5 and Al/NbO2 samples during the wetting
experiments at different temperatures from 700 – 1300 °C. It can be observed that the spreading
process of molten Al on the Nb2O5 and NbO2 substrates took place with increasing temperature
above 1150 °C. The Al-droplet height decreases and the interfacial diameter increases.
Fig. 4.11 shows the variations in the wetting angle of the molten Al on Nb2O5 and NbO2
between 700 °C and 1300 °C. The wetting angles are strongly temperature dependent, and decrease
from ~ 125° to ~ 30° as the temperature increases from 700 °C to 1300 °C. It is worth noting that
the wetting angles are smaller than 90° at T > 1150 °C.
Fig. 4.12 shows the wetting angle of molten Al on Nb2O5 and NbO2 as a function of time at the
temperature of 1200 °C, which decreases from ~ 75° to ~ 25° in 60 min.
- 4 Results -
- 37 -
Fig. 4.9 Photographs of Al/Nb2O5 samples during the wetting experiments (wetting angle) at
different temperatures.
700 °C
1300 °C 1200 °C
1150 °C 1100 °C
1000 °C 900 °C
800 °C
5 mm
- 4 Results -
- 38 -
Fig. 4.10 Photographs of Al wetting experiment on NbO2 at different temperatures.
700 °C
1300 °C 1200 °C
1150 °C 1100 °C
1000 °C 900 °C
800 °C
5 mm
- 4 Results -
- 39 -
Fig. 4.11 Variations in the wetting angle of the molten Al on Nb2O5 and NbO2 between 700 °C and
1300 °C.
Fig. 4.12 Variation in the wetting angle of the molten Al on Nb2O5 and NbO2 with time at
temperature of 1200 °C.
- 4 Results -
- 40 -
4.1.3 Microstructure of reaction composites
Due to effective interconnected porosity (Fig. 4.2 and 4.6) and improved wettability of NbO2
preforms by Al-melt (Fig. 4.12), pressureless infiltration could be achieved at 1200 ºC. The reaction
between molten Al and NbO2 preforms resulted in the formation of NbAl3/Al2O3 composites,
NbO2 (s) + 13/3 Al (l) NbAl3 (s) + 2/3 Al2O3 (s) ΔGr (1400 °C) = - 298.197 kJ/mol (4.3)
Phase composition examined by XRD reveals that pressed and infiltrated CN1 composite
contained NbAl3 and -Al2O3. In Fig. 4.13 is presented a representative scanning electron
micrograph of NbAl3/Al2O3 composite (CN1) structure: the composite consists of dark alumina
grains and bright NbAl3 phase; the average grain size of Al2O3 is ~ 2.5 µm and the grain size of
NbAl3 is between 5 µm and 10 µm. The composite is dense with density of 3.7 g/cm3 and an open
porosity of less than 0.7 %.
Fig. 4.13 SEM micrograph of the polished cross section of pressed and infiltrated CN1 composite
(light gray: NbAl3, dark gray: Al2O3) 3.
50 µm
- 4 Results -
- 41 -
Phase composition examined by XRD reveals that printed and infiltrated CNA composite
contained NbAl3, -Al2O3 and residual Al. Fig. 4.14 shows SEM micrographs of printed and
infiltrated composite CAN: the average grain size of Al2O3 is ~ 2.5 µm and the grain size of NbAl3
is between 5 µm and 10 µm. The measured density of the composite is 3.6 g/cm3, and the measured
residual open porosity is ~ 5 %.
Fig. 4.14 SEM micrograph of the polished cross section of printed and infiltrated CAN (Light gray:
NbAl3, dark gray: Al2O3 and residual Al.
4.2 Nb-Al-C system
4.2.1 Printed, CIP-ed and sintered Nb2AlC
Fig. 4.15 shows the particle size distribution of powder mixture NNA1 ready for 3DP. The
prepared powder exhibits a bimodal size distribution with maxima of particle sizes of 2 µm and 12
µm. SEM analysis shows fine NbC particles surrounded by coarse Nb and Al particles.
20 µm 1 mm
(a) (b)
- 4 Results -
- 42 -
Fig. 4.15 Particle size distribution of powder mixture NNA1 for 3DP.
Fig. 4.16 shows the position of printed samples in the build bay of 3D-printer. The geometric
density of printed samples was measured as a function of their position in build bay, Fig. 4.17. The
geometric density of printed samples decreases from ~ 2.8 to ~ 2.2 g/cm3, as their position is varied
from left to right. 3D-printer spreads dry powder from the feed box to cover the surface of the build
platform in thin layers. The particle size of the powder mixture NbC/Nb/Al lied in the range
between ~ 1 µm and ~ 100 µm, Fig. 4.15. Due to the Van der Waals attractive force between
particles, the flowability of fine powders is substantially poorer than of large particles. During the
powder spreading finer NbC particles (d50 ~ 0.9 µm) filled the interstices between larger Nb (5 – 45
µm) and Al (< 45 µm) particles, and the fraction of NbC continuously decreased from the left to the
right along the build platform, resulting in a density gradient in the 3D-printed samples. The
formation mechanism of the density gradient is discussed in detail in chapter 5.1.1.
- 4 Results -
- 43 -
Fig. 4.16 Position of printed samples in build bay of 3D-printer.
Fig. 4.17 Effect of position on the density of the printed Nb-Al-C samples.
- 4 Results -
- 44 -
The printed sample had a porous structure with mean open porosity of ~ 36 % and mean pore
size of ~ 4 µm. After CIPing at a pressure of 200 MPa the open porosity decreased to ~ 20 % with a
mean pore size of ~ 0.3 µm. Typical microstructure along fracture surface of the printed and CIP-ed
samples is shown in Fig. 4.18 (a) and (b), respectively.
Fig. 4.18 Facture surface of prepared Nb-Al-C samples: (a) printed; (b) printed and CIP-ed at a
pressure of 200 MPa.
Fig. 4.19 shows the pore size distribution of the printed and CIP-ed samples. The printed sample
exhibits two broad peaks at 0.2 − 10 µm and between 40 and 90 µm. After CIPing at 50 MPa, the
large pores in the range of 40 − 90 µm are first eliminated. Higher compaction loads of 100, 150
and 200 MPa lead to a decrease in the average pore size, which is in a narrow range between 0.3
and 0.5 µm.
Fig. 4.20 shows the measured linear shrinkage as a function of applied CIP pressure along
different directions between the green and the CIP-ed stage. The linear shrinkage increases as the
applied CIP pressure increases. The shrinkage exhibits anisotropic behavior with the maximum
shrinkage along the height direction and minimum shrinkage along the length direction.
20 µm 20 µm
(a) (b)
- 4 Results -
- 45 -
Fig. 4.19 Pore size distribution of printed and CIP-ed Nb-Al-C samples.
Fig. 4.20 Effect of applied CIP pressure on the linear shrinkage of Nb-Al-C samples.
- 4 Results -
- 46 -
Fig. 4.21 shows the volume shrinkage of CIP-ed samples as a function of applied CIP pressure.
The volume shrinkage increases from ~ 11 % to ~ 24 % with increasing applied CIP pressure. The
geometric density of the CIP-ed samples continuously increases from 2.7 to 3.6 g/cm3 as the applied
CIP pressure increases, Fig. 4.22.
Pressureless reactive sintering of 1450 °C for 30 min resulted in the formation of NbC, NbAl3
and Nb2AlC as confirmed by XRD, Fig. 4.23, and EDS. A gradient was found with a NbC rich
surface and Nb2AlC and NbAl3 defected in the core region. After reactive sintering, a substantial
increase of the open porosity to ~ 60 % was observed. It could be assumed that interaction between
the Al-Nb reaction and the binder burnout led to porous microstructure during the pressureless
sintering, which is discussed in detail in chapter 5.1.2.
Fig. 4.21 Effect of applied CIP pressure on the volume shrinkage of Nb-Al-C samples.
- 4 Results -
- 47 -
Fig. 4.22 Effect of applied CIP pressure on the density of Nb-Al-C samples.
Fig. 4.24 shows the pore size distribution of the reaction sintered Nb-Al-C samples prepared at
different CIP pressures. Reactive sintering resulted in an increase in large pore sizes of 1 – 10 µm,
compared with the pore size distribution of the CIP-ed green bodies (Fig. 4.19).
4.2.2 Hot-pressed Nb2AlC
The phase compositions examined by XRD are summarized in Table 4.1. Dense single-phase
Nb2AlC was obtained by hot-pressing at 1650 °C for 90 min under 30 MPa. The density measured
by Archimedes method is 6.44 0.22 g/cm3 which is close to the theoretical density of 6.5 g/cm3. A
polished and etched surface of Nb2AlC is shown in Fig. 4.25: the average grain size of Nb2AlC is ~
17 µm.
- 4 Results -
- 48 -
Fig. 4.23 XRD patterns for the surface and the center of the sintered sample.
Fig. 4.24 Pore size distribution of reactive sintered Nb-Al-C samples prepared at different CIP
pressure.
- 4 Results -
- 49 -
Table 4.1 Summary of phase analysis of the Nb-Al-C samples hot-pressed at the temperature range
of 600 −1650 °C 2.
Temperature (°C) Phase Compositions
600 NbC, Nb, Al
700 NbC, NbAl3, Nb
900 NbC, NbAl3, Nb
1100 NbC, NbAl3, Nb, Nb2Al, Nb2AlC
1300 NbC, NbAl3, Nb2Al, Nb2AlC
1500 Nb2AlC, Nb2Al, NbC
1650 Nb2AlC
Fig. 4.25 SEM micrograph of Nb2AlC sample etched cross section after hot-pressing at 1650 °C
under a pressure of 30 MPa for 90 min 2.
20 µm
- 4 Results -
- 50 -
4.2.3 Thermal properties
The average CTE of hot-pressed Nb2AlC in the temperature range of 30 − 1050 °C is 8.1 × 10-6 K-1,
which can readily be approximated by a linear dependence of expansion on temperature, as shown
in Fig. 4.26. In the present work, the thermal diffusivity of 0.08 × 10-4 m2/s was measured. The
specific heat capacity of Nb2AlC was 382.6 J (kg K)-1 [Bar02-1]. According to the equation (3.4),
the thermal conductivity of Nb2AlC at room temperature is calculated as 20 W (m K)-1.
Fig. 4.26 Temperature dependence of thermal expansions of Nb2AlC.
4.2.4 Mechanical properties
Fig. 4.27 shows that the Vickers hardness of the hot-pressed Nb2AlC is strongly dependent on
indentation load. Salama et al. [Sal02] measured the hardness of hot isostatically pressed Nb2AlC
and the value was 6.1 ± 1 GPa. In the present work, the Vickers hardness decreases as the
indentation load increases from 1 N to 300 N and appears to approach an asympotic value of ~ 4.5
GPa, Fig. 4.27.
- 4 Results -
- 51 -
Fig. 4.27 Dependence of Vickers hardness of Nb2AlC on the applied load 2.
No cracks are observed to emanate from the corners of indentations and the material was pushed
out around the indent, Fig. 4.28 (a). Extended delamination and laminate kinking were observed in
region where the crack passed through Nb2AlC grains, Fig. 4.28 (b). Like for all other MAX phases
[Rag97, Rag99-2, Bar00-2], even at the highest indentation loads applied (300 N) no cracks from
the corners were detected. The delamination, kinking and pull-out of the Nb2AlC grains (Fig. 4.28 b)
around the indent is unusual for nonmetals. For metals, duo to plastic deformation the materials rise
above the undamaged surface after the indentations; as opposed to this, ceramics materials exhibit
sink-in after the indentations [Mar82, Zen96, Rag97]. Damage mechanisms around hardness
indentations in Ti3SiC2 were reported by El-Raghy et al. [Rag97]. Typical damage mechanisms in
Ti3SiC2 (similar to Nb2AlC in this work) were detected by El-Raghy et al. [Rag97]: grain buckling;
kinking of microlaminates; delamination on the basal planes of Ti3SiC2 crystals; crack deflection
along the basal planes; laminate fracture; grain push-out/pull-out. These multiple damage
mechanisms suggest that MAX phases offer microscale plasticity.[Rag97].
- 4 Results -
- 52 -
Fig. 4.28 Indent morphology of Nb2AlC after indentation with load of 100 N.
Fig. 4.29 shows the four-point bending strengths as a function of indentation loads. The
measured four-point bending strength for hot-pressed Nb2AlC is ~ 380 MPa (unpolished samples)
and ~ 440 MPa (polished samples), respectively. After indentation with load of 100 N (indent
diagonal: ~ 204 µm), the measured bending strength is ~ 418 MPa and no strength degradation is
observed. At the indentation load of 200 N (indent diagonal: ~ 295 µm), the measured bending
strength is 383 MPa, which is about 86% of the strength of the undamaged sample. Even after
indentation under the load of 300 N (indent diagonal: ~340 µm, about 10% of the sample’s width),
the measured bending strength is ~ 377 MPa and no drastic decrease in bending strength is
observed, Fig. 4.29, indicating a pronounced damage tolerance as reported in literature for other
MAX phases [Bar00-1, Sal02, Koo03, Bar04]. Focusing on the work of Salama et al. [Sal02],
polycrystalline, fully dense, predominantly single-phase samples of Nb2AlC with an average grain
size of 14 ± 2 µm were fabricated by reactive hot isostatic pressing of Nb, graphite, and Al4C3 at
1600 °C for 8 h and 100 MPa. Its average four-point bending strength is ~ 420 MPa. After
indentation with load of 300 N, the residual bending strength of ~ 230 MPa is about 55% of the
strength of the undamaged sample, Fig. 4.29.
100 µm
(b)
10 µm
(a)
- 4 Results -
- 53 -
Fig. 4.29 Indentation load dependence of residual bending strength of Nb2AlC 2.
After indentation and four-point bending tests, a fracture surface is shown in Fig. 4.30. The
dashed line indicate the extension of the damage zone under the Vickers indentation. The measure
fracture toughness is ~ 5.9 MPa m1/2. Assuming semicircular surface crack geometry and taking the
values for σc of 418 MPa, 383 MPa and 377 MPa, critical defect sizes ac of 63 µm, 76 µm and 78
µm, respectively, are calculated from the Griffith’s relation (ac = (KIC/c)2) which are in good
agreement with the observed expansion of the damage zone.
The fracture mode is predominantly transgranular with the crack path showing a complete non-
planar morphology, Fig. 4.31. Extensive delamination and laminated kinking give rise for a high
crack resistance. Zhou and Sun [Zho01-1] investigated the deformation behavior of Ti3SiC2 under
room temperature compression: kinking, delamination of individual grains, dislocation slip and
intergranular fracture resulted in microscale plasticity in Ti3SiC2.
Fig. 4.32 shows the residual bending strength as a function of quenching temperature, which
contrasts with the work of Salama et al. [Sal02]. Up to a temperature difference of 600 °C strength
only displayed a minor reduction where as at higher temperature differences thermal shock caused a
considerable reduction of strength.
- 4 Results -
- 54 -
Fig. 4.30 Fracture surface of Nb2AlC tested in four-point bending after Vickers indentation with
300 N load 2.
Fig. 4.31 Fracture surface showing delamination and laminate kinking of the Nb2AlC grains 2.
Table 4.2 summarizes the measured properties of hot-pressed Nb2AlC. For comparison, data of
Nb2AlC [Bar02-1, Sal02], Ti2AlC [Bar00-1, Bar00-2, Wan02-1], Cr2AlC [Tia06, Lin05], and
Ta2AlC [Hu08-1] are also included.
10 µm
- 4 Results -
- 55 -
Fig. 4.32 Effect of quenching temperature on the bending strength of Nb2AlC 2.
Table 4.2 Properties of Nb2AlC phase materials compared to isomorphous Cr, Ti and Ta MAX
phases 2.
Properties Nb2AlC
present work
Nb2AlC
[Bar02-1, Sal02]
Ti2AlC
[Bar00-1, Bar00-2,
Wan02-1]
Cr2AlC
[Tia06, Lin05]
Ta2AlC
[Hu08-1]
Density (g cm-3)
Average grain Size (µm)
Coefficient of thermal expansion (×10-6 K-1)
Thermal conductivity at 25 °C (W m-1 K-1)
Vickers Hardness (GPa)
Bending strength (MPa): 4-point bending
Bending strength (MPa): 3-point bending
Fracture toughness (MPa m1/2)
Young’s modulus (GPa)
6.44 ± 0.22
17
8.1
20
4.5 ± 0.3
443 ± 28
481 ± 42
5.9 ± 0.3
294
6.37 ± 0.02
14
8.7
22
6.1 ± 1
413 ± 16 - -
286
4.11
45
8.2
46
2.8 -
275
6.5
305
5.21 -
13.3
17.9
3.5 -
378 -
278
11.46
3/15
8.0
28.4
4.4 ± 0.1 -
360 ± 19
7.7 ± 0.2
292
- 4 Results -
- 56 -
4.3 Ti-Al-O-C system
4.3.1 Microstructure
According to the XRD analysis, the as-fabricated composite was mainly composed of Ti3AlC2,
TiAl3, Al2O3 and residual Al and TiC. Fig. 4.33 shows a SEM-micrograph of the microstructure of
the Ti3AlC2/Al2O3/TiAl3 composite. Plate-like grains of Ti3AlC2 with a length of 10 – 50 µm are
distributed homogeneously in a matrix of Al2O3 and TiAl3, Fig. 4.33 (Table 4.3). EBSD confirmed
that the basal planes (0001) of Ti3AlC2 are parallel to the long axis of these particles. Yin et al.
[Yin07-1, Yin07-2] reported on the reaction mechanism of Ti3AlC2 reinforced composites by the Al
melt infiltration into the porous TiO2/Ti2O3/TiC preforms: firstly, reaction between infiltrated Al
and Ti2O3 may lead to the formation of Al2O3 and TiAl3:
32332 28 OAlTiAlAlOTi (4.4)
Subsequently, TiC may react with Ti-saturated Al solution or TiAl3 to form the ternary phase
Ti3AlC2 [Son04]:
AlAlCTiTiCTiAl 22 233 (4.5)
Combining the above reactions (4.4) and (4.5), the total reaction is given as
3232332 62 OAlTiAlAlCTiAlTiCOTi (4.6)
- 4 Results -
- 57 -
(a)
(b)
Fig. 4.33 Typical microstructure of Ti3AlC2/Al2O3/TiAl3 composite. 1: Ti3AlC2, 2: Al2O3, 3: TiAl3.
20 µm
12
3
200 µm
- 4 Results -
- 58 -
Table 4.3 EDS taken from 1, 2 and 3 regions of Fig. 4.33.
Atomic % O Al Ti Al/Ti O/Al
1 0.00 19.57 80.43 0.24 0
2 59.55 40.45 0.00 1.47
3 0.00 64.49 35.51 1.82 0
4.3.2 Fracture behavior
Fig. 4.34 shows a stress-strain curve of Ti3AlC2/Al2O3/TiAl3 composite in four-point bending test at
room temperature. No non-elastic contribution of deformation was observed up to the peak load of
fracture. The measured fracture toughness of Ti3AlC2/Al2O3/TiAl3 composite is 8.3 0.3 MPam1/2,
which is higher than that of Al2O3/TiAl3 composite (7.1 MPam1/2) [Tra03] and most other brittle
ceramics.
Fig. 4.34 Stress-strain curve of Ti3AlC2/Al2O3/TiAl3 composite in four-point bending test.
- 4 Results -
- 59 -
Direct observation of the whole crack propagation during testing provides further information
regarding the fracture mechanisms of the investigated composites, Fig. 4.35. SEM micrograph of
crack propagation of Ti3AlC2/Al2O3/TiAl3 specimens shows that the crack paths are significantly
tortuous in nature.
Fig. 4.35 SEM micrograph of crack propagation of the Ti3AlC2/Al2O3/TiAl3 specimens
(The arrow indicates the direction of crack propagation).
Detailed analysis shows the crack deflection (tilting and twisting) through the Ti3AlC2 grains
(Fig. 4.36 a), crack propagating along the Ti3AlC2-matrix interface (Fig. 4.36 b, c) as well as along
or through the Ti3AlC2 grains (crack branching) (Fig. 4.36 d). These phenomena of crack
propagation increase the crack length and absorbed fracture energy, resulting in enhanced fracture
toughness [Sar07, Yin07-1]. Similar phenomena such as crack bridging, delamination, deflection,
branching (translamellar fracture) and pull-out of MAX phase grains were reported by Sarkar et al.
[Sar07] and Yin et al. [Yin07-1], leading to enhanced crack growth resistance. These characteristics
of crack propagation can be traced to their nanolaminate crystal structure, mixed bonding (covalent,
ionic and metallic), unique deformation mechanisms [Bar99-1, Bar99-2, Zho01-2]. Barsoum et al.
[Bar99-1] have reported a deformation mode for Ti3SiC2: shear band formation by dislocation
arrays, cavitations, creation of dislocation walls and kink boundaries, buckling and delamination of
Ti3SiC2 grains, which is discussed in detail in chapter 5.2.1 and 5.2.2.
1 mm
- 4 Results -
- 60 -
Severe buckling and delamination of the plate-like Ti3AlC2 grains were observed in the vicinity
of the crack propagate path, Fig. 4.36. The experimental observations show transgranular fracture to
occur in Ti3AlC2 grains. Cleavage energies (Gc) for several crystal planes in this material were
calculated. D is defined as the plane between the Al layer and TiC6 octahedra, A and B are defined
as the planes between the Ti-I and C and Ti-II and C, respectively, all of them parallels to the (0001)
basal plane, and finally F is defined as the (10ī0) cleavage plane, Fig. 4.37. The corresponding
cleavage energy Gc for the fracture along the planes A, B, D and F are presented in Table 4.4. The
values of Gc for A(0001) and B(0001) are 4.83 J/m2 and 6.23 J/m2, respectively, much greater than
1.34 J/m2 for D(0001). The weak interaction between Ti and Al will promote the transgranular
fracture in Ti3AlC2 along the surfaces between the Al layer and the TiC6 octahedra along (0001)
basal planes. It is expected that the delamination observed in Fig. 4.36 (a) occurs on these Al
terminated (0001) planes parallel to the longitudinal grain orientation [Zha07-2]. Concerning the
bonding anisotropy, the corresponding value of Gc for the plane F (3.11 J/m2) is lower than the
values for A(0001) and B(0001) planes but is still more than two times larger than that of D(0001)
plane. This result can explain qualitatively, that a Ti3AlC2 grain with the basal plane perpendicular
to the crack path can stop the advance of the crack, Fig. 4.36 (b) (see arrow 1 and 2).
- 4 Results -
- 61 -
Fig. 4.36 In-situ fracture series for SEVNB specimen of Ti3AlC2/Al2O3/TiAl3 composite in four-point
bending test.
Table 4.4 DFT-calculated cleavage energies for Ti3AlC2.
Plane Gc (J/m2)
B (0001) 6.23
A (0001) 4.83
D (0001) 1.34
F (10ī0) 3.11
30 µm 30 µm
(b) (a)
(c) (d)
30 µm 20 µm
1
2
- 4 Results -
- 62 -
Fig. 4.37 Layered crystal structure of Ti3AlC2 (left); bonding charge density of Al/Ti terminated
(10ī0) surfaces in Ti3AlC2 (right) (Red, green and blue color means high, middle and low electron
density, respectively. Bonding charge calculation provided by Dr. Pavel Leiva-Ronda).
Fig. 4.37 (right) presents electron charge density distribution in a plane perpendicular to the
(0001) cleavage planes in the Ti3AlC2 crystal. The charge density distribution calculation suggests
the interaction between Ti and Al to be significantly weaker than the Ti-C interaction. Thus, crack
propagation along the interfaces between the Al and the TiC6 octahedra layer on the (0001) basal
plane is likely to govern transgranular delamination which agrees with the experimental observation
of crack tilting and twisting. Zhou et al. [Zho01-3] employed ab initio calculations based on density
functional theory to reveal the electronic structure and bonding properties of Ti3AlC2. Zhou et al.
[Zho01-3] reported that titanium, carbon and aluminum atoms form Ti(2)-C-Ti(1)-C-Ti(2) and
Ti(2)-C-Ti(1)-C-Ti(2)-Al chains. The interatomic distance between Ti(1) and C and between Ti(2)
and C was 2.2068 Å and 2.0886 Å, respectively; the distance between Al and Ti(2) is 2.8783 Å
[Zho01-3]. Thus, the weak bonding exists between Al and the Ti(2)-C-Ti(1)-C-Ti(2) chain. These
x
z
c
Ti3AlC2 D
A
B
Ti-I
Ti-II
Ti
C
Al
F
- 4 Results -
- 63 -
results [Zho01-3] agree with the present work. According to the charge density distribution around
titanium, carbon and aluminum atoms and difference in electronegativity between these atomes, Ti-
C bond can be characterized as a mixture of ionic and covalent bonding; and Ti-Al bond is a mixed
ionic, covalent and metallic bonding [Zho01-3].
- 5 Discussion -
- 64 -
5 Discussion
5.1 3DP multistep processing of composites 5.1.1 3DP In the 3D-printer dry powder is spread from the feed box to cover the surface of the build platform
in thin layers. For fine powders (particle size < 10 µm), their high surface area and the Van der
Waals attractive force between particles cause extensive agglomeration. Cohesive strength of the
unpacked powder also increases, and as a result, the flowability of fine powders is substantially
poorer than large particles [Yoo96]. Thus, the fraction of fine powder particles continuously
decreases from the left to the right along the build platform, resulting in the density gradient in
green bodies. In this work, the particle size of the powder mixture NbC/Nb/Al lied in the range
between ~ 1 µm and ~ 100 µm (see Fig. 4.15). During the powder spreading finer NbC (d50 ~ 0.9
µm) particles filled the interstices between larger Nb (5 – 45 µm) and Al (< 45 µm) particles, and
the fraction of NbC continuously decreased from the left to the right along the build platform.
Therefore, the printed green samples exhibited a density gradient.
The post-printing step, CIPing, was employed to achieve high and uniform packing density (see
Fig. 4.22 and 4.24). After CIPing, however, the linear shrinkage exhibited anisotropic behavior with
the maximum shrinkage along the Z-axis of the 3D-printer (see Fig. 4.20). The anisotropic
shrinkage was already introduced during powder spreading and/or printing process [Kha96]. During
powder spreading the fine powder consisted of agglomerates which caused inhomogeneities in the
powder bed [Yoo96]. The smallest building unit in a 3D-printed sample can be characterized as a
primitive, which means an agglomerate of powder particles (or granules) formed by a single binder
droplet [Lau92]. The interaction between powder particles and binder solution means a
rearrangement of the powder particles, resulting in an inhomogeneity in the packing density of the
powder [Gir95, Gir96]. In addition, the packing density of the printed samples is affected by the
orientation of the samples in the build bay of 3D-printer. This is accompanied by printed bands
along the X-axis (gantry direction of travel), continuous strips along Y-axis (cartridge direction of
travel) and laminated layers along Z-axis [Cha07]. Thus, the stitching between printed bands is
strongest along the Y-axis and weakest along the Z-axis [Gir95], resulting in the shrinkage
anisotropy observed (see Fig. 4.20). It is necessary that a post-printing stage CIPing is performed to
- 5 Discussion -
- 65 -
increase the packing density, to narrow the pore size distribution in order to increase the
densification rate upon sintering.
5.1.2 Reaction and microstructure control The carbothermal reduction of niobium oxide resulted in the release of carbon monoxide (CO), Eq.
(4.1) and (4.2). The partial pressure of CO could be determined by thermodynamic calculations, as
shown in Fig. 5.1 (a). A high CO partial pressure of 4.6 × 10-2 MPa was calculated at about 1150 °C,
which may result in improved vapor transport during sintering [Sil01, Rea84, Rea86, Rea87,
Qua89]. TGA results confirmed that considerable weight loss occurred above approximately
1150 °C. Improved vapor transport during sintering can accelerate the neck growth and particle
coarsening and inhibit further densification, resulting in porous structure [Rea84]. In the present
work, improved vapor transport resulted in the formation of porous microstructure of sintered NbO2
preforms. A phase stability diagram was determined by thermodynamic calculation, Fig. 5.1 (b):
NbO2 begins to form at temperature above 700 °C and formation of NbO and NbC begins at
temperature exceeding 1100 °C, which are in good agreement with experimental observation of
formation of niobium carbide at temperature above 1150 °C.
- 5 Discussion -
- 66 -
Fig. 5.1 (a) Partial pressure of CO for Eq. (3.1); and (b) phase stability diagram in the system
Nb2O5 - NbO2 - NbO - NbC - C associated with the reactions upon reduction 3 (calculated by means
of equiTherm Version 5.04i [Bar97-2]).
(a)
(b)
- 5 Discussion -
- 67 -
Capillary force driven spontaneous wetting and infiltration of a porous ceramic skeleton by
liquid metal requires an open porosity with suitable pore size and shape in the preform [Hil88]. Fig.
5.2 shows typical microstructures of pre-sintered CN1 preforms, which were fabricated by uniaxial
pressing and 3DP, respectively. The as-printed preform exhibits a non-uniform microstructure due
to the agglomeration of powder particles, resulting in incomplete Al-infiltration. During 3DP the
interaction between the ceramic powder bed and the binder liquid determines the microstructure and
dimension of a single primitive [Lau92]. “The printed object is constructed by stitching this
primitive together with adjacent primitives between lines and layers” [Moo01]. Therefore, the
properties of powder bed and binder liquid play an important role in determining microstructure of
the printed object. In general, fine powders can spontaneously agglomerate due to van der Waals
forces [Aks84]. During 3DP, agglomerates result in a non-uniform powder bed structure with a
local packing density gradient [Moo01]. Additionally, the surface tension forces of the binder
exceed the cohesive strength of the powder bed, causing particle rearrangement and anisotropic
pore structure during printing, Fig. 5.3 [Yoo96]. In some cases, the printing inhomogeneities
(defects) introduced by powder-binder interaction can not be completely eliminated by post-
processing [Yoo96, Lau92]. Fracture origins (intergranular defects) were observed on the fractured
surfaces of 3D printed alumina [Gir96]. Therefore, a more uniform packing density of powder bed,
and a successful interaction between the powder particles and the binder solution are the
preconditions for homogenous microstructure in the 3D-printed samples [Yoo96, Gir95, Gir96].
- 5 Discussion -
- 68 -
Fig. 5.2 Microstructure of sintered CN1 preforms prepared under different processing: (a) printed;
(b) pressed.
Fig. 5.3 Decrease in the packing density of printed powder bed as a result of binder-powder
interaction. Left portion of the micrograph shows the unprinted region while the other half shows
lower packing due to rearrangement of granules [Yoo96].
100 µm 100 µm
(a) (b)
- 5 Discussion -
- 69 -
According to XRD results (Table 4.1), no reaction occurred at temperatures below 700 °C. At
700 °C and 900 °C, no reaction occurred with NbC. When the temperature increased to 700 °C,
aluminum was completely reacted with niobium to form the NbAl3.
33 NbAlNbAl (5.1)
Above 900 °C, NbAl3 is supposed to react with Nb to form Nb2Al:
AlNbNbNbAl 23 35 (5.2)
When the temperature was raised to 1100 °C, Nb2AlC was detected. Within the temperature
range from 1300 °C to 1500 °C, the amount of Nb2AlC increased with the consumption of NbAl3,
Nb2Al and NbC [Hu08-2, Hu08-3]. After heating treatment at 1650 °C only Nb2AlC phase was
formed. The reaction to form Nb2AlC can be described by Eq. (5.3) 2:
AlCNbNbCAlNbNbAl 223 552 (5.3)
Salama et al. [Sal02] reported that after heating treatment at 1600 °C for 8 h dense Nb2AlC
samples with an average grain size of 14 ± 2 µm were obtained, and no grain growth was observed
even after heating treatment at 1600 °C for 16. In the present work, the average grain size of
synthesized Nb2AlC is ~ 17 µm. However, Ti2AlC samples with an average size of about 300 µm
(aspect ratio: 10) were fabricated by hot pressing at 1600 °C for 4 h [Bar97-1, Sal02]. Thus,
compared to other MAX phases, the grain growth of Nb2AlC is extremely sluggish [Sal02].
It is important to note that the pressureless sintered samples were porous with average open
porosity of ~ 60 % and cracks were observed after sintering. Cracking was not found when the
samples were fabricated by uniaxial pressing using powder mixture NNA2 without binder, followed
by CIPing and reactive pressureless sintering. Thus, we assume that the binder burnout is the main
reason which caused cracking. According to DT/TGA results, the burnout of dextrin is performed at
a temperature range of 250 − 800 °C. However, aluminum was completely reacted with niobium to
NbAl3 up to 700 °C (Table 4.1 and Eq. 5.1). The fractional volume change upon reaction is
calculated from:
- 5 Discussion -
- 70 -
)()(3
)()(3)( 3
0 NbVAlV
NbVAlVNbAlV
V
V
mlm
mlmm
(5.4)
where the Vm of various phases is: Vm(Al)l = 11.2 cm3/mol at 700 °C [Hat84, Yin06]; Vm(Nb) = 10.8
cm3/mol; Vm(NbAl3) = 38.3 cm3/mol. A volume shrinkage upon reaction of ~ 14 % is determined.
Simultaneously, the dextrin may have to be removed as gas by thermal decomposition, resulting in
a volume expansion. Therefore, it could be assumed that interaction between the Al-Nb reaction and
the binder burnout led to cracking during the pressureless sintering.
After reactive pressureless sintering, significant differences in phase composition existed
between the surface and the center of fabricated samples (see Fig. 4.23). In the case of Ti3SiC2, it is
noted that the evaporation of Si may cause the higher content of TiCx phase at the surface of
fabricated sample through reactive sintering of Ti/Si/2TiC (molar ratio of 1:1:2) [Li99]. Li and
Miyamoto [Li99] pointed out that the evaporation of Si led to a Si-deficient liquid phase, especially
in the surface region, which, inhibited the formation of Ti3SiC2. Table 5.1 shows the vapor pressure
of aluminum at different temperature [Hat84]. In this work, the evaporation of Al may result in the
higher content of NbC at the surface of pressureless sintered sample, especially when the
temperature is above 1300 °C. Therefore, it is difficult to synthesize single phase and dense Nb2AlC
ceramic from NbC/Nb/Al/Dextrin using 3D-printing, followed by CIPing and reactive pressureless
sintering. From the present experiments, dense Nb2AlC ceramic could be synthesized from the
green compact NbC/Nb/Al without binder using reactive hot-pressing. However, this technology
can only produce components with simple geometries. The future work could be focused on a
combination of 3D-printing and HIPing, in order to fabricate dense complex-shape ceramic parts.
For example, the feasibility of fabricating high-density parts from Inconel 718 powder using 3DP
was assessed, and subsequently HIP-ed to achieve full density [Sic08].
- 5 Discussion -
- 71 -
Table 5.1 Vapor pressure of aluminum [Hat84].
Temperature (°C) Vapor pressure (Pa)
827 7.4×10-6
927 3.7×10-3
1127 0.3
1327 7.8
1527 98
5.1.3 Wetting and infiltration Wetting of ceramic substrates by liquid metals is an important aspect and has a significant impact
on fabrication of ceramic/metal composites. In the present work, temperature (see Fig. 4.11), time
(see Fig. 4.12) and interfacial reaction strongly influenced the wettability of Nb-O preforms by
molten Al. Wetting can be improved by an interfacial chemical reaction (negative Gibbs free
energy). The reaction product niobium aluminide (atomic ratio of Al/Nb: 2.8, Table 5.2) was
formed at the interface during wetting at 1200 °C, Fig. 5.4. The formation of interfacial reaction
product lead to the further decrease of wetting angle between molted-Al and Nb-O preforms with
time (see Fig. 4.12). At a temperature above 1150 °C, the improved wettability with wetting angle
smaller 90° between molted-Al and Nb-O preforms, and homogeneous microstructure of Nb-O
preforms (see Fig. 4.2 and 4.6) were found to provide adequate condition for pressureless
infiltration of Al melt.
Time, pore size, pore shape, porosity, viscosity and surface tension of Al melt have a strong
effect on the infiltration kinetics. Infiltration depth as a function of time at 1200 °C was calculated
using the available literature and experimental data according to Eq. (2.8):
21
cos25.61
r
thp
p (2.8)
- 5 Discussion -
- 72 -
Taking values for porosity p of 0.63, mean particle radius of NbO2 preform r of 1 µm, porosity
shape factor of 0.5 – 1 [Lap00], viscosity of the Al-melt at 1200 °C = 0.78 mPas according to
Eq. 5.5 [Hat84], surface tension of Al melt at 1200 °C = 0.79 J/m2 according to Eq. 5.6 [Hat84],
)/5.1984exp(1492.0 T [mPas] (5.5)
310)(152.0868.0 mTT [J/m2] (5.6)
wetting angle θ (°) calculated as a function of time t (sec.) according to the results of wetting test:
t013.072 (5.7)
the infiltration depth h was calculated as a function of time for:
)013.072cos(0002.0 tth [m] (5.8)
Fig. 5.4 Interfacial microstructure for the sample of molten on Nb2O5 at 1200 °C for 1 h in vacuum
(< 10 Pa).
500 µm 100 µm
(a (b)
- 5 Discussion -
- 73 -
Table 5.2 EDS results taken from 1, 2, 3 and 4 region of Fig. 5.4 (b).
Atomic % O Al Nb Al/Nb
1
2
3
4
-
70.67
73.64
3.49
73.37
-
-
96.51
26.63
29.33
26.36
-
2.8
XRD, SEM and EDS results have demonstrated that the porous NbO2 was infiltrated and
reacted with Al resulting in the formation of NbAl3/Al2O3 composites. The formation of these
phases can be explained with the help of the ternary Nb-Al-O phase diagram, Fig. 5.5 [Zha94]. The
total porosity in the preform defines the amount of infiltrated Al. The infiltrated Al is available for
the redox reaction assuming that there is no residual Al melt:
23232 23133 xNbOOAlNbAlAlNbOx (5.9)
The fractional volume change upon reaction is described:
)(13)()3(
)(13)()3()()(2)(3
2
22323
0 AlVNbOVx
AlVNbOVxNbOxVOAlVNbAlV
V
V
mm
mmmmm
(5.10)
Vm(i) are the molar volumes of the various phases i (Vm(NbAl3) = 38.3 cm3/mol; Vm(Al2O3) = 25.7
cm3/mol Vm(NbO2) = 21.2 cm3/mol; Vm(Al)l = 11.6 cm3/mol at 900 °C [Yin06]. For the case that no
residual NbO2 is available, the volume shrinkage of ~ 22 % is calculated. According to:
)(13)(3
)(13
2 AlVNbOV
AlV
mm
mp (5.11)
the total porosity εp of ~ 70 % is calculated. The volume fractions of NbAl3 to Al2O3 in the reaction
composite are 69 % and 31 %, respectively. The exothermic reaction between the infiltrating metal
melt and the ceramic preforms results a volume shrinkage. However, the linear change of
component less than 1 % can be achieved by reactive infiltration processing [Mül99]. Rigid
structure of preforms can enhance the stability of the component shape during the reactive melt
- 5 Discussion -
- 74 -
infiltration processing; in addition, “local volume change is compensated by flow of excessive melt
into (volume contraction of solid phase) or out of the component (volume expansion of solid
phase)” [Yin06, Kum99].
Fig. 5.5 Ternary Nb-Al-O phase diagram at 1100 °C [Zha94, Sch98-2, Sch00].
Yin et al. [Yin07-2] reported that the reaction mechanism of reactive infiltration of Al melt into
TiO2/TiC preform according to DTA results, Fig. 5.6 [Yin07-2]: two endothermic peaks at 678 °C
and 1303 °C indicate the melting of Al and TiAl3, respectively; two exothermic peaks at 973 °C and
1040 °C may be associated with reaction of the formation of TiAl3/Al2O3 and Ti3AlC2, respectively.
According to the Ti-Al-C phase diagram [Pie94], TiAl3 is stable with the existence of Ti3AlC2 at
1300 °C, and not stable with TiC, which also suggests TiAl3 may react with TiC to form Ti3AlC2
[Yin07-2].
- 5 Discussion -
- 75 -
Fig. 5.6 Differential thermal analysis results of a sintered TiO2/TiC preform infiltrated with Al melt
[Yin07-2].
5.1.4 Surface finish and accuracy of 3DP
In the present work, metal melt infiltration process resulted in dense composite materials with
non-improved surface roughness. Speed and accuracy are the functional requirements of an RP
system. The prototype quality is evaluated by surface finish, dimensional and form accuracy
obtained from RP process. The important challenges are improving the surface roughness and
dimensional accuracy of 3D-printed parts. The surface quality plays an important role in
improving the dimensional accuracy, optimizing the surface structure or textures, reduction of
surface defects and post-processing, and enhancing the mechanical properties of 3D-printed
components [Art96].
Melcher [Mel09] has studied the surface roughness of 3D-printed objects using special parts,
Fig. 5.7 [Mel09]: the surface roughness of three different planes (0°-plane, 45°-plane and 90°-plane)
was measured by laser scanning microscopy; the surface roughness of printed, sintered, glass-
infiltrated and Cu-O-infiltrated samples was examined, respectively, Fig. 5.8 [Mel09]. The
measured surface roughness results were summarized [Mel09]: average roughness of sintered
samples is ~ 50 µm lower for all surfaces than that of green samples; for sintered samples, the
- 5 Discussion -
- 76 -
average roughness of the 0°-surface is slightly smoother than the 45°- and 90°-surfaces are;
considering the standard deviation no significant difference between sintered and infiltrated samples
can be drawn; infiltration process resulted in dense composite materials with non-improved surface
roughness. For 3D-printed objects, surface roughness of more than 40 µm was reported, which is
much more than the roughness for processes such as SLS, LOM and FDM [Ipp95, Kar98].
Fig. 5.7 Testing part for surface finish measurement [Mel09]: (A) CAD model; (B) sintered Al2O3;
(C) Cu-O-infiltrated.
Fig. 5.8 Surface finish of testing parts in green, sintered and infiltrated state depending on different
planes (0°/45°/90°) [Mel09].
- 5 Discussion -
- 77 -
In general, parts fabricated by 3DP exhibit ~ 40 – 60 Vol. % open porosity. If such parts are
sintered to full density they will have high dimensional change of ~ 15 – 20 % linear shrinkage
[All00, Zha09]. Due to the high dimensional accuracy (linear shrinkage smaller than 0.1 %) for
typical application of 3DP such as rapid tooling and manufacturing, a combination process of
sintering and melt infiltration is used for densification. Metal parts for tooling with high
dimensional accuracy was fabricated by melt infiltration into the 3D-printed and lightly sintered
preforms [All00]: debinding/sintering and melt infiltration process resulted in ~ 2 % linear
shrinkage and ~ 0.3% linear expansion, respectively, leading to a total shrinkage value of about
1.7%, which can be adjusted by modifying the CAD model [All00]. 3DP multistep processing (3DP,
sintering and melt infiltration) can be used to fabricate metal tooling parts with good surface finish
and high dimensional accuracy [All00]. In former studies, Yin et al. have studied the dimensional
accuracy of 3D-printed, sintered and Al-infiltrated composite: total shrinkage of the final Al-
infiltrated composite was less than 3.2 % compared with the CAD model used for 3DP [Yin07-2].
5.1.5 Comparison and application
During 3DP, agglomerates result in a non-uniform powder bed structure and non-uniform
microstructure of printed sample (see Fig. 5.2 (a)). Good flowability of powders used for 3DP plays
an important role in improving their homogeneity, packing density and green density during the
powder spreading [Pau96]. Different methods were used to optimize the flowability of powders:
powders can be granulated (spray dried granules); spherical particles can be used [Gir96];
particles/granulates can be coated with organic additives [Bea97]; furthermore, electrical [Mel91]
and magnetic [Mel92] fields, vibration mechanism [Sac00] and mechanical agitation [Bun95] were
developed and applied. In order to enhance the homogeneity, packing and green density of powder
bed, uniaxial pressing was used after the spreading of each powder layer [Pau96]; the press-rolling
technique was used to create well packed powder layers [Yoo96]. In addition, post densification
process such as isostatic pressing, melt infiltration can be used to achieve dense
materials/composites with enhanced physical and mechanical properties [Bea97, Cim95]. Freeform
fabrication of ceramic based composites such as TiC/Cu [Ram05], Al2O3/Cu-O [Mel06], SiSiC
[Tra06], TiAl3/Al2O3 [Yin06], Ti3AlC2/TiAl3/Al2O3 [Yin07-1], and Al2O3/glass [Zha09] using 3DP
multistep processing was demonstrated. Comparison of typical mechanical properties such as
fracture toughness and bending strength of ceramic materials fabricated by 3DP and other technique
- 5 Discussion -
- 78 -
are summarized in Table 5.3. A combination process of 3DP and melt infiltration can produce the
parts with complex geometry and offers a high potential to make accessible applications, Fig. 5.9.
Table 5.3 Comparison of fracture toughness and bending strength of ceramic materials fabricated
by 3DP and other technique.
Materials Processing Proportion (Vol. %)
KIC
(MPa m1/2) Bending strength
(MPa) Ti3AlC2 [Wan02-2]
TiAl3 [Mil01]
Ti3AlC2/TiAl3/Al2O3
Ti3AlC2/TiAl3/Al2O3 [Yin07-1]
TiAl3/Al2O3 [Mül99]
Ti3AlC2/TiC/Al2O3 [Che06]
Ti3AlC2/Al2O3 [Che04]
Hot pressing
Ace melting
3DP
3DP
Reactive casting
Combustion reaction
Hot pressing
-
-
35/30/10
35/30/10
70/30
-
90/10
7.2
2
8.3
8.1 – 9.7
6
5.8
8.7
340
162
-
320
> 400
466
425
- 5 Discussion -
- 79 -
Fig. 5.9 Complex geometry parts by 3DP: a Al2O3-based moulding dies [Rep04, Mel06]; b glass-
infiltrated half skull [Zha09]; c infiltrated turbine wheel [Zha09]; d glass-infiltrated jaw; e macro-
cellular SiSiC [Sch10].
2 cm 2 cm
2 cm
2 cm
a
b c
ed
2 cm
- 5 Discussion -
- 80 -
5.2 Mechanical behavior of MAX phase composites
5.2.1 Deformation and damage mechanisms
Zhang et al. [Zha04] have studied the deformation and damage mechanisms of Ti3SiC2 induced by
indentation, which are summarized schematically in Fig. 5.10 [Zha04]. Duo to the typical nano-
laminate crystal structure, sliding is the basic deformation mode in Ti3AC2 (A: Al or Si), Fig. 5.10
(a) [Zha04, Bar99-2]. With increased stress the maximum shear stress will occur, leading to another
two very important deformation modes buckling (Fig. 5.10 (b)) and kinking (Fig. 5.10 (c)) in
Ti3AC2 grains [Zha04, Bar99-1]. Sliding along grain boundaries can lead to the intergranular
cracks/fracture, Fig. 5.10 (d) [Zha04]. In general, the formation of buckling can result in two
damage modes: cleavage fracture and delamination, Fig. 5.10 (e) and (f) [Zha04, Bar99-1]. In
addition, kinking can cause crack propagation along the kinking boundaries and delamination
cracking, Fig. 5.10 (g) [Zha04, Bar99-1]. Therefore, MAX phases Ti3AC2 (A: Al or Si) offer the
multiple deformation and damage modes to make a contribution to microscale plastic deformation
[Zha04].
Fig. 5.10 Schematic of deformation and damage mechanisms of Ti3AC2 (A: Al or Si) [Zha04].
(a) Sliding (b) Buckling (e) Fracture
(g) Delamination/Fracture
(c) Kinking
(d) GB Cracking
(f) Buckling/Delamination
- 5 Discussion -
- 81 -
Barsoum et al. [Bar99-1] have studies the formation mechanisms of deformation and damages
modes, and summarized a dislocation-based model: “the basic elements of the model are shear
deformation by dislocation arrays, cavitation, creation of dislocation walls and kink boundaries,
buckling, and delamination” [Bar99-1]. The model can be used to explain experimental
observations. Fig. 5.11 shows the fracture surface of Ti3AlC2 reinforced composite after four-point
bending, containing the typical damage mechanism of MAX phase such as buckling and
delamination.
Fig. 5.11 Fracture surface of Ti3AlC2/Al2O3/TiAl3 composite after four-point bending test showing
fracture mechanism of buckling, delamination and cleavage fracture.
5.2.2 Quasi-plasticity
The plastic behavior of MAX phases Ti3AC2 (A: Al or Si) was explained by their layered structure
and the metallic nature of the bonding in the Al and Si layers [Bar99-1, Gil00, Bar99-2]. In the case
of Ti3SiC2, the plasticity is induced by the multiple basal plane (001) slip at room temperature
[Low98, Zho01-1]. Mechanisms of plastic deformation in Ti3SiC2 at room temperature involve
relief of local stress and strain fields from kink band (boundaries) formation, buckling and
delamination of individual grains [Bar99-1, Sar07]. Similar to Ti3SiC2, the kinking of the
microlaminates and the absence of cracking at the kinks (see Fig. 4.28) suggest that Nb2AlC
fabricated in the present work exhibits quasi-plasticity at room temperature [Rag00-2]. In addition,
dense Nb2AlC with Vickers hardness (Hv) of ~ 4.5 GPa and Young’s modulus (E) of ~ 294 GPa
was fabricated in the present work using reactive hot-pressing. This low Hv/E ratio suggests that the
20 µm
- 5 Discussion -
- 82 -
mechanical behavior of Nb2AlC is similar to that of ductile metals [Low98]. Contact damage after
Vickers indentation showed characteristics of quasi-plastic materials (see Fig. 4.28 and 4.30).
Compared to brittle ceramics, the fabricated Nb2AlC is significantly more damage tolerant. These
results are similar to the results described in the review of Barsoum for other ternary MAX phases
[Bar00-1]. The mechanical behavior of MAX phase materials can be derived from their unique
characteristics [Bar99-1, Bar99-2, Koo03, Zhe04]:
1) Only basal plane dislocations exist, which are mobile and multiply, even at temperatures as
low as 77 K [Bar99-1, Bar99-2, Sal02, Bar04].
2) Dislocations can arrange themselves either in arrays (pileups) or walls (tilt and twist
boundaries) normal to the arrays [Far98, Far99, Bar99-1, Bar00-1, Koo03, Bar04].
3) Due to the high c/a ratios, typical characteristics such as glide, formation of kink bands
(KBs) and delamination play an important role in the deformation of MAX phases [Koo03,
Far98, Far99, Bar99-1, Bar99-2].
Hess and Barrett [Hes49] developed a dislocation-based model that can be used to explain the
formation of KBs, which is summarized schematically in Fig. 5.12 [Bar99-1, Bar04, Hes49]: with
further deformation maximum shear stresses occur at two sections of L/4 and 3L/4, Fig. 5.12 (a, b);
above a critical value these shear stresses are sufficient to trigger a pair of dislocations of opposite
sign that move in opposite directions, Fig. 5.12 (c); dislocations moving leads to the formation of
kink bands between the unkinked crystal, resulting in kink boundaries BC and DE, Fig. 5.12 (d)
[Bar99-1, Bar04]. Hess and Barrett have also pointed out that KBs are expected only in crystals,
such as hexagonal metals or alloys having an axial c/a ration greater than ~ 1.73 [Hes49]. With c/a
ration of ~ 4.5 [Bar00-1], it is not surprising that Nb2AlC deforms by KBs.
- 5 Discussion -
- 83 -
Fig. 5.12 Model of the formation of kink bands [Hes49, Bar99-1, Bar04].
5.2.3 Crack propagation and structure modeling
Coefficients of thermal expansion mismatch:
Residual stresses originate upon cooling from the processing temperature and are due to the
coefficients of thermal expansion (CTE) mismatch between different phase compositions. The
fracture toughness increment, ΔK, of a dispersion reinforced matrix composite arising from the
thermal residual radial stress σr can be calculated by Tay90, Li07-2:
)(2
2dD
K r
(5.12)
f
dD
085.1 (5.13)
d is the average grain size of dispersed phase, D the average distance between dispersed particles,
and f the volume fraction of dispersed particle. r is the residual radial stress in the matrix at the
point with a distance of R from the centre of the dispersed particle Sel61, Li07-2:
- 5 Discussion -
- 84 -
3
)21(2)1(
2
R
r
EE
ETE
pmmp
pmr
(5.14)
where Δα=αp-αm. α is the coefficient of thermal expansion, the subscript m and p refer to matrix and
dispersed particle, respectively, ν is the Poisson’ ratio, E is the Young’s modulus, ∆T is the
temperature difference, r is the radius of dispersed particle [Sel61, Li07-2]. The type of the stress,
tension or compression, depends on the sign of Δα. The coefficient of thermal expansion of TiAl3 (α
~ 13 × 10-6 K-1) is higher than both for Ti3AlC2 (α ~ 9.0 × 10-6 K-1) and Al2O3 (α ~ 8.3 × 10-6 K-1),
Table 5.4. For the case that TiAl3 matrix (m) contains either Ti3AlC2 or Al2O3 dispersed particle (p),
then αp < αm, which can result in a compressive stress in dispersed particle, and a tensile stress in
matrix [He09]. The presence of tensile stress may cause microcracks in matrix around the dispersed
particle, when the dispersed particle size exceeds a critical value. As a result, the microcrackings
can lead to significant crack branching and deflection, resulting in enhanced resistance to crack
propagation and higher toughness of composite [Eva84, He09].
The effect of process temperature on the grain size of Ti3AlC2 and R-curve behavior of
Ti3AlC2/TiC/Al2O3 were studied by Yin et al. [Yin07-1]: an average grain size of 5 µm in length
and 2 µm in thickness at a process temperature of 1300 °C can be achieved; at the higher process
temperature of 1400 °C led to an increase in the grain size of 50 µm in length and 5 µm in thickness
[Yin07-1]. R-curve behavior was observed in the Ti3AlC2/TiC/Al2O3 samples [Yin07-1]: fracture
toughness of fabricated samples at 1300 °C increased from 7 MPa m1/2 to 8.6 MPa m1/2 as the crack
length increased up to 2.1 mm; for samples fabricated at 1400 °C, fracture toughness increased from
9.6 and 34.8 MPa m1/2 with a crack length of 2.5 mm, Fig. 5.13 [Yin07-1].
- 5 Discussion -
- 85 -
Fig. 5.13 Variation of fracture toughness as a function of crack extension (R curves) for Ti3AlC2
reinforced TiAl3/Al2O3 prepared at 1300 °C and 1400 °C [Yin07-1].
Table 5.4 Properties E, CTE and ν of TiAl3, Ti3AlC2 and Al2O3.
Property TiAl3
[Mil01]
Ti3AlC2
[Bar00-1, Fin00]
Al2O3
[Che04]
E (GPa)
CTE (10-6 K-1)
υ
156
13
0.16
297
9
0.2
386
8.3
0.21 – 0.27
- 5 Discussion -
- 86 -
Crack deflection and bridging:
The intergranular fracture propagating along the TiAl3/Ti3AlC2 and TiAl3/Al2O3 interfaces results in
a crack deflection and an increase of the composite toughness. In addition, the crack wake shielding
such as crack bridging can be observed by Ti3AlC2 grains (see Fig. 4.36 b, arrow 1 and 2). For
small-scale bridging, the fracture toughness increment can be estimated from [Bar02-2, Li07-2]:
0EfqKb (5.15)
where E is the elastic modulus of the composite; f, q, σ0, and χ is the volume fraction, characteristic
dimension, yield stress and dimensionless function representing the work of rupture of the
reinforcement, respectively [Bar02-2, Li07-2]. Taking values for E of ~184 GPa for composite
[Yin07-1], f of ~ 0.35 for Ti3AlC2, q of ~ 15 µm, σ0 (compressive yield strength) of ~ 560 MPa for
Ti3AlC2 [Tze00], and χ with an estimated value of ~ 0.4 [Li07-2], the toughness increment due to
crack bridging by Ti3AlC2 is approximately 14.7 MPa m1/2, which was calculated according to Eq.
5.15.
Chemical bonding and electronic structure
Deformation and damage mechanisms of MAX phases depend not only on the grain structure
(sliding, kinking and buckling) [Bar99-1, Zha03], but also on their chemical bond structure and
stress-strain deformation at the atomic level [Med08]. Simulation methods can be used to study
electronic structure, chemical bond structure, phase transition, mechanical properties of MAX
phases [Zha07-2]. For example, the atomic and electronic structures of some MAX phases have
been studied and reported [Ahu00, Zho01-3, Mag05]. In addition, ab initio calculations of cleavage
characteristics of MAX phases can be used to study their deformation and damage mechanisms of
MAX phases [Med08]. In Ti3AC2 (A = Si, Al) there are three types of chemical bonds, TiI-C, TiII-
C and TiII-A (A = Si, Al). The calculated cleavage energies in Ti3AC2 (A = Si, Al) are summarized
in Table 5.5. The calculated results agree with the experimental study of Ti3SiC2 [Bar99-1], which
found the nature of the Ti-Si bond to be relative weak when compared with the Ti-C bond.
Medvedeva et al. [Med08] have pointed out that a significant stretching of the Ti-Si bonds occurred
under tensile stress according to the results of ab initio full-potential linearized plane wave
calculations, which can be used to explain the damage mechanisms of Ti3SiC2. For the Ti3AlC2
- 5 Discussion -
- 87 -
system, TiII-Al is the weakest and TiII-C is the strongest of the three chemical bonds. Therefore, it
can be predicted that when the tensile stress is applied on Ti3SiC or Ti3AlC2, the break point is TiII-
Si or TiII-Al, respectively. In addition, the cleavage energy of TiII-Al bond is less than that of the
TiII-Si bond (see Table 5.5) [Zha07-2]. The difference in electronegativity can be used to
understand this phenomenon: the electronegativities of Ti, Al and Si are 1.54, 1.61 and 1.90,
respectively [Dea99, Zha07-2]; the difference in electronegativity of Ti-Si (0.36) is much greater
than that of Ti-Al (0.07) [Zha07-2]. The Ti-Si bond is stronger than the Ti-Al bond, as a result,
shear and deform of Ti-Si will be more difficult, which agree with the experimental results that
Ti3SiC2 has a larger shear and bulk modulus than Ti3AlC2 [Rad06, Fin00].
Table 5.5 DFT-calculated cleavage energy Gc (J/m2) of Ti3AC2 (A = Al, Si).
System TiII-Si/Al TiII-C TiI-C Reference
Ti3AlC2
Ti3AlC2
Ti3SiC2
Ti3SiC2
1.34
2.07
2.88
3.16
6.23
6.44
6.33
6.16
4.83
4.68
5.07
7.16
Present work
[Zha07-2]
[Zha07-2]
[Fan06]
MAX phases are nanolaminates, which can be characterized by interleaved layers with high and
low electron density [Mus06]. This fact is connected with the fracture mechanism and some
electronic structure calculations using the density functional theory have been performed for these
systems on the last years [Med08, Wan08, Mus06, Zha07-2]. Between the remarkable results
corresponding to the cases of Ti3SiC2 and Ti3AlC2, it can mention that the preferential habit
cleavage is the A terminated (0001) plane for TiII/A bonding with A: Al or Si [Med08, Zha07-2].
From the experimental aspect, nano-laminated Ti3AlC2 offers a significant toughening
performance in Ti3AlC2-Al2O3-TiAl3 composites which showed non-catastrophic failure behavior.
Nb2AlC fabricated in the present work exhibits quasi-plasticity at room temperature and good
damage tolerant. Similar deformation and damage mechanisms can be observed in Ti3SiC2 samples
under compression and indentation loads [Zha04]. The load state obtained by bending test in this
study and the composite mechanical properties allow bending to large angle and consequently the
onset of buckling and delamination for the ternary carbide. Zhang et al. [Zha04] pointed out that
- 5 Discussion -
- 88 -
sliding, buckling and kinking are three basic deformation mechanisms under indentation tests.
Transgranular and intergranular cracking result in local damage in Ti3SiC2 polycrystals, where the
formation of delamination cracks and cleavage fracture proceed along the basal (0001) plane
[Zha04, Koo03, Bar99-1].
- 6 Summary and Conclusions -
- 89 -
6 Summary and Conclusions
Binary and ternary MAX-based composites were fabricated using 3DP multistep processing
technique. While 3DP provides an opportunity to produce porous ceramic preforms with a high
degree of freedom in geometry and shape, a subsequent liquid metal infiltration into these preforms
offers a way to fabricate dense materials through exothermic reaction. Upon sintering prior to melt
infiltration a porous microstructure with interconnected porosity in the preform can be achieved.
Furthermore, a carbothermal reduction process of Nb2O5 to NbO2 was found to provide adequate
condition for pressureless infiltration of Al melt.
Nanolaminate structure of MAX phase Nb2AlC grains give rise for extended plasticity resulting
in excellent damage tolerance and thermal shock resistance. The high capacity of Nb2AlC for
absorbing and distributing damage during Vickers indentation has been demonstrated. Under
compression-shear deformation Nb2AlC exhibits a quasi-plasticity deformation behavior, which can
be explained by the multiple basal plane slip between microlamellae, intergrain sliding, lamellae, or
grain push-out. A kinking-based model [Hes49, Bar99-1, Bar04] can explain the quasi-plasticity
and damage tolerance triggered by the nanolaminate structure of Nb2AlC.
Nano-laminated Ti3AlC2 offers a significant toughening performance in Ti3AlC2-Al2O3-TiAl3
composites which showed non-catastrophic failure behavior. Similar deformation and damage
mechanisms can be extensively observed in Ti3SiC2 samples subjected to compression and
indentation loads [Zha04]. Ab initio calculations of cleavage energy and electron density in Ti3AlC2
crystal confirmed the experiment-based deformation and damage mechanism of Ti3AlC2 in the
composites.
- Einleitung -
- 90 -
Einleitung
Keramische Werkstoffe und Keramik-Verbundwerkstoffe spielen eine wichtige Rolle für
Anwendungen als Leichtbauteile im Automobil- und Luftfahrtsektor. Materialforscher und
Hersteller suchen nach keramischen Materialien, die bessere Eigenschaftskombinationen aufweisen.
Intermetallische/Keramik-Verbundwerkstoffe wurden entwickelt, die die günstigen Eigenschaften
der Keramik, wie z. B. hohe Verschleißbeständigkeit, niedrige Dichte und gute Korrosions- und
Oxidationsbeständigkeit, mit hoher Duktilität und Zähigkeit der metallischen Komponente
kombinieren. Die Materialgruppe der ternären Carbide und Nitride besitzt eine
Nanoschichtmikrostruktur der allgemeinen Summeformel Mn+1AXn (oder MAX), wobei n 1, 2, oder
3 ist, M einen Übergangsmetall ist, A ein Element der IIIA- und IVA- Metalle ist, wie z. B.
Aluminium und Silizium, und X entweder Kohlenstoff oder Stickstoff ist. Diese bieten ein hohes
Potenzial für neuartige technische Anwendungen, die einen erhöhten Anspruch hinsichtlich der
mechanischen Leistungsfähigkeit verlangen. Hergestellt werden MAX-Materialien hauptsächlich
über die Heißpresstechnik. Ein äußerer Druck wird benötigt, um die Festkörper-Reaktion zwischen
den Pulver-Komponenten zu beschleunigen. Daher können nur einfache Formkörpergeometrien
hergestellt werden, was die Anwendungsgebiete der Produkte einschränkt.
Rapid Prototyping bietet ein breites Spektrum an Formgebungsverfahren, die komplexe Formteile
direkt durch Computer Aided Design (CAD)-Daten erzeugen können. Beim dreidimensionalen
Drucken (3D-Drucken) wird das 3D-Objekt durch Aufspritzen flüssigen Binders auf ein
Pulvermaterial schichtweise aufgebaut. 3D-Drucken bietet die Möglichkeit, poröse keramische
Vorkörper unterschiedlicher Geometrie herzustellen. Durch eine anschließende Schmelzinfiltration
können dichte Materialien erzeugt werden. Zu den Vorteilen der Kombination von 3D-Drucken und
Schmelzinfiltration gehören eine bessere Kontrolle der Mikrostrukturentwicklung und
Eigenschaften, wie auch die Verwendung von kostengünstigen Precursoren.
Die vorliegende Arbeit behandelt die Erforschung und Entwicklung neuartiger Verarbeitungsketten
zur Herstellung von MAX-Phasen-basierten Verbundwerkstoffen durch den Einsatz des
dreidimensionalen Druckens. Auf Basis vorläufiger thermodynamischer Berechnungen wurden die
folgende Systeme ausgewählt: Nb-Al-O, Nb-Al-C and Ti-Al-O-C. Der Arbeitsplan basiert auf der
Formgebung mittels dreidimensionalem Drucken und Umsetzung der porösen Formkörper in einen
- Einleitung -
- 91 -
dichten Verbundwerkstoff mittels reaktiver Schmelzinfiltration. Die Eigenschaften der gedruckten
Formkörper sowie der durch Schmelzinfiltration erzeugten Verbundwerkstoffe wurden ermittelt.
Die wissenschaftliche Herausforderung in dieser Arbeit liegt in der Steuerung des Benetzungs- und
Infiltrationsverhaltens der Metallschmelze sowie der Flüssig-Fest-Reaktion zur Bildung homogener
und dichter Verbundwerkstoffe. 3D-gedruckte MAX-Phasen Komposite zeigen eine ausgezeichnete
Schadenstoleranz durch lokale Verformungsmechanismen in der Nanolaminatstruktur.
- Zusammenfassung -
- 92 -
Zusammenfassung
Binäre und ternäre MAX-Phasen-basierte Komposite wurden über dreidimensionalen Drucken (3D-
Drucken) hergestellt. 3D-Drucken bietet die Möglichkeit, poröse keramische Vorkörper mit einem
hohen Freiheitsgrad in der Geometrie herzustellen. Durch anschließende reaktive
Schmelzinfiltration wurden dichte Materialien erzeugt. Durch Sintern der 3D-gedruckten Grünlinge
konnte ein poröses Gefüge mit einem durchgängigen Porennetzwerk erreicht werden, dass günstige
Bedienung für die Infiltration mit metallischen Schmelzen bietet.
Die Nanolaminatstruktur der MAX-Phase Nb2AlC bietet die Voraussetzung für quasi-plastische
Verformung, die zu ausgezeichneter Schadenstoleranz und Temperaturwechselbeständigkeit führt.
Die Schadenstoleranz wurde über lokale Verformungsexperimente (Vickerseindrückmethode)
nachgewiesen. Bei einer Druckscherverformung weist Nb2AlC ein plastisches
Verformungsverhalten auf, was durch die Basalgleitung zwischen Mikrolamellen dominiert wird,
die zu ausgeprägten Lamellen-Knick-Vorgängen (Kinking) [Hes49, Bar99-1, Bar04] führt.
Ti3AlC2-Nanolaminat basierte Verbundwerkstoffe bieten eine Bruchzähigkeitssteigerung in den
Ti3AlC2-Al2O3-TiAl3 Kompositen, die kein katastrophales Bruchversagen aufweisen. Ähnliche
Verformungs- und Schadensmechanismen wurden in Ti3SiC2 Proben unter Druckbelastung und
nach Vickerseindrücken beobachtet [Zha04]. Ab-initio-Rechnungen der Spaltungsenergie und
Elektronendichteverteilung im Ti3AlC2 Kristall bestätigen die experimentellen Ergebnisse der
Verformungs- und Schadensmechanismen-Hypothese von Ti3AlC2.
- References -
- 93 -
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- List of publications/Veröffentlichungen -
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List of publications/Veröffentlichungen
1. W. Zhang, R. Melcher, N. Travitzky, R.K. Bordia, P.Greil
Three-Dimensional Printing of Complex-Shaped Alumina/Glass Composites
Adv. Eng. Mater., 11 [12], 1039–43 (2009).
2. W. Zhang, N. Travitzky, C.F. Hu, Y.C. Zhou, P. Greil
Reactive Hot Pressing and Properties of Nb2AlC
J. Am. Ceram. Soc., 92 (10) 2396-2399 (2009).
3. W. Zhang, N. Travitzky, P. Greil
Formation of NbAl3/Al2O3 Composites by Pressureless Reactive Infiltration
J. Am. Ceram. Soc., 91 (9) 3117-3120 (2008).
4. R. Melcher, W. Zhang, N. Travitzky, P. Greil
3D-Printing of Al2O3/Cu-O composites
Ceramic Forum International Special Edition - Rapid Prototyping, 83 (13) 18-22 (2006).