(Rapid Prototyping von Keramik/Metall-Verbundwerkstoffen) · (Rapid Prototyping von...

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


- 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


- 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]:


- 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].


- 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


- 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


- 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].


- 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


- 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].


- 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

-


- 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.


- 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


- 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


- 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)


- 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].


- 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


- 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.


- 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.


- 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


- 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.


- 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


- 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


- 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)


- 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


- 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)


- 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).


- 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


- 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


- 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 -

- 113 -

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).

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