Advanced High Temperature Alloys

Advanced High 

Temperature Alloys 

Prof. Dr.-Ing. Uwe Glatzel 

Metals and Alloys 

University Bayreuth 

SS 2008 

1 

University Bayreuth, Advanced High Temperature Alloys Uwe Glatzel, Metals and Alloys

Lecturer: 

Prof. Dr.-Ing. habil. Uwe Glatzel 

• born Dez. 1960 

• Physik-Diplom (B.Sc. and M.Sc) in Tübingen 

(exchange year in Corvallis, Oregon, USA) 

• PhD thesis at the Institute for Metals Research, Technical 

University Berlin, Prof. Monika Feller-Kniepmeier 

• post-doc (1 Jahr) at Stanford University 

• Habilitation TU-Berlin 

• Gerhard-Hess award of the German Science Foundation 

(DFG) for young scientist (400.000 €) 

• 1996-2003 full professor for Metals and Alloys, Jena 

• since April 2003 Bayreuth (Chair for Metals and Alloys) 

postal address: 

Ludwig-Thoma-Str. 36b phone: +49 (0) 921 - 55-5555 

D-95447 Bayreuth, Germany e-mail: uwe.glatzel@uni-bayreuth.de 

2 


Literature 

• R. Bürgel, Handbuch Hochtemperatur-Werkstofftechnik, Vieweg 

• R.C. Reed, The Superalloys - Fundamentals and Applications, Cambridge Univ. Press 

• H. Frost, M. Ashby, Deformation-Mechanism Maps, Pergamon Press 

• G. Meetham, M. Van der Voorde, Materials for High Temperature Engineering 

Applications, Springer 

• J. Betten, Creep Mechanics, Springer 

• Askeland: Materialwissenschaften, Spektrum Lehrbuch; 1994 

• Callister: Materials Science and Engineering - An Introduction, Wiley, New York, 1999 

• H. Schumann, Metallographie, Deutscher Verlag für Grundstoffindustrie, Leipzig 

• F. Vollertsen, S. Vogler, Werkstoffeigenschaften und Mikrostruktur, Hauser Verlag 

• P. Haasen, Physikalische Metallkunde, Springer-Verlag, Berlin 

• H.-J. Bargel, G. Schulze, Werkstoffkunde, VDI-Verlag, Düsseldorf 

• P. Sarrazin, A. Galerie, J. Fouletier, Mechanisms of High Temperature Corrosion, Trans 

Tech Publications 

lecture notes: http://www.uni-bayreuth.de/departments/metalle/ then follow "Lehre" 

then you will find the link to the lecture notes of this lecture 

University Bayreuth, Advanced High Temperature Alloys 3 

Uwe Glatzel, Metals and Alloys

What You Should Know: 

• basic thermodynamics 

• introduction to diffusion 

• introduction to dislocations 

• phase diagrams 

• theory of elasticity 

• ... 

• basic materials science courses 

4 


Contents 

1. Introduction, Basics 

2. Stability of Microstructure 

3. Mechanical Properties 

a) Static 

b) Cyclic (Fatigue) 

4. High Temperature Corrosion 

5. High Temperature Alloys 

6. Lost Wax Investment Casting 

5 


Introduction 

• only alloys will be looked at (no ceramics, no 

polymers). 

• no coatings (BUT in principle: all high 

temperature systems are coated!), simply not 

enough time. 

6 


Polymer 

Glass 

Metals 

Ceramics 

Composits 

Maximum Temperatures for 

Applications of Different Materials 

Group 

maximum service temperature 

[°C] 

up to 300 

up to 800 

Fe-Basis (coated) up to 1100 

Fe-ODS up to 1300 

Ni-base up to 1200 

Pt-base up to 1600 

refractory metals in inert 

atmosphere above 1600 

MoSi 2 up to 1800 

SiC up to 1600 

(SiC/C) up to 1600 

deformation/damage mechanism 

melting, decomposing (pyrolyze) 

viscous flow 

creep, dislocation climb, 

grain boundary sliding 

viscous flow, glass transition 

temperature, grain boundary 

sliding 

complex 

7 


Nutzbare Festigkeit 

Overview Materials 

500 Temperatur [°C] 1500 2000 

Quelle: 

Plansee AG, 

Reutte, 

Tirol, 

Austria 

8 



Taking Density into Account 

500 Temperatur [°C] 1500 2000 

9 



Oxidation Resistance 

500 Temperatur [°C] 1500 2000 

10 


Material Choice 

• temperature 

• environment 

• moving/non-moving part 

• design complexity (how to manufacture) 

• price constrictions (depending on application 

of system). Reduction of 1 kg in weight: 

– car ~ 0 - 5 € 

– plane ~ 100 – 500 € 

– aerospace ~ 100.000 - 500.000 € 

11 


• temperature: 

Influence of ... on ... 

– phase transitions, volume fractions, ... 

– diffusion (� recrystallization, dislocation climb, diffusional creep, ... ) 

– thermal fatigue (TF) 

• mechanical: 

– creep 

– fatigue (low cycle, LCF, high cycle fatigue, HCF) 

• environment: 

– oxidation 

– corrosion 

• combinations: 

– thermo-mechanical fatigue (TMF) 

– stress corrosion cracking, stress oxidation, ... 

12 


Basics 

Boltzmann-statistics: energy of 

movement increases with temperature 

Thermodynamics ↔ Kinetics 

3 

u total = 2⋅ 

u kin = 2⋅ 

k B ⋅T 

= 3⋅ 

k B ⋅T 

atom 

atom 2 

Utotal mol 

= 3⋅ 

R ⋅T 

� 0,33 eV, bzw. 32 kJ/mol bei 1000°C 

3 

u kin = k B ⋅T 

atom 2 

ε& = ε& 

Q 

R T 

0 e ⋅ 

− 

⋅ 

Arrhenius-plot 

13 


c 

v 

Vacancy Concentration 

F = U - T·S � non-zero vacancy concentration is 

in thermodynamic equilibrium 

= 

c v 

e 

T[°C] 

T/T m 

Qvac 

− 

R⋅T 

20 

0.17 

10 -23 

Q vac nickel = 1,36 eV (energy necessary to create one vacancy) 

300 

0.33 

3·10 -12 

450 

0.42 

10 -9 

800 

0.62 

10 -6 

equilibrium vacancy concentration for nickel 

1000 

0.74 

10 -5 

1200 

0.85 

7·10 -5 

1454 

1.00 

3·10 -4 

14 


Nickel Vacancy Concentration 

vacancy concentration 

10 0 

10 -5 

10 -10 

10 -15 

10 -20 

10 -25 

Nickel Vacancy Concentration 

0 200 400 600 800 1000 1200 1400 1600 

temperature [°C] 

15 

University Bayreuth, Advanced High Temperature Alloys Uwe Glatzel, Metals and Alloys 

Tm

J = −D 

⋅∇c 

vacancy diffusion or 

volume diffusion 

Diffusion 

1. Fick's law 

[J] = (atoms) · m -2 ·s -1 

[D] = m 2 ·s -1 

[c] = (atoms) · m -3 

16 


Qvac Qmigration QSD Coefficient of Diffusion 

D 

= 

energy to create a vacancy 

activation energy to migrate a vacancy 

activation energy for volume diffusion 

Q SD = Q vac + Q migration 

D 

0 

⋅e 

−( 

Q 

vac 

+ Q 

k⋅T 

migration 

−Q 

k⋅T 

Q SD ≈ 17 ·k B ·T m � Q SD nickel ≈ 2.5 eV = 244 kJ/mol 

(for a perfect crystal; defects will lower the activation energies) 

17 


) 

= 

D 

0 

⋅e 

SD

Dependence Melting Point and 

Enthalpy of Vacancy Creation 

element 

Pb 

Al 

Cu 

Ag 

Ni 

Pt 

Mo 

W 

T m 

[°C] 

327 

660 

1 085 

1 235 

1 455 

1 768 

2 623 

3 422 

17·R·T m 

0.88 

1.36 

1.99 

2.21 

2.53 

2.98 

4.23 

5.40 

Q vac 

[eV] 

0.57 

0.68 

1.29 

1.12 

1.32 

3.00 

4.00 

crystal 

structure 

18 


fcc 

fcc 

fcc 

fcc 

fcc 

fcc 

bcc 

bcc

Q SD versus T m 

400 kJ/mol 

0.137 kJ/(mol·K) 

≈ 17 · kB ·NA 19 


Coefficient of Diffusion 

Steep slope indicates a 

high activation energy. 

Small elements diffuse 

faster. 

Diffusion in fcc crystals 

slower than in bcc crystals. 

20 


Coefficient of Diffusion with Defects 

Coefficient of diffusion of Th 

in W. 

Overall velocity for diffusion 

depending on grain boundary 

thickness, grain size and 

dislocation density. 

21 


dashed line: 

Pipe Diffusion 

diffusion in crystal by the velocity of pipe diffusion 

⎛ atoms ⎞ 

⎜ ⎟ 

⎝ time ⎠ 

D eff = D SD + a disl. · ρ ·D disl. 

adisl. area of dislocation core 

( ≈ 5 b2 ≈ 0.3 nm2 ) 

ρ dislocation density 

Ddisl. pipe diffusion along 

dislocation core 

atom flux ~ D·area 

22 


D 

SD 

grain 

⋅d 

~ D 

grain 

SD 

= D 

⋅d 

2 

grain 

identical atom fluxes if: 

disl. 

⋅ b 

2 

⎛ 

⎜ 

⎝ 

n 

⋅ 

d 

grain 

atoms ⎞ 

2 

⎟ ~ Ddisl. 

⋅b 

time ⎠disl. 

= D 

disl. 

⋅ b 

2 

⋅ρ 

⋅n

Grain Boundary Diffusion 

dashed line: diffusion in crystal by the velocity 

of grain boundary diffusion 

D eff = D SD + π · δ / d·D grain bound. 

with: 

δ effective grain boundary 

thickness ( ≈ 2 b ≈ 0.5 nm) 

d grain size 

Ddisl. pipe diffusion along 

dislocation core 

23 


Activation Energies Indicating 

Mechanism Changes 

~ Q SD 

Single crystal aluminium, oriented such that {111} slip is activated. 

Lytton, Shepard and Dorn, Trans. AIME 212 (1958) 220 

24 


Diffusion in Ordered Structures 

(Intermetallic Phases) 

• High binding energies � high activation 

energies � low coefficient of diffusion 

• For example NiAl: very low enthalpy of ordered 

B2 structure � low enthalpy outweighs entropy 

� ordered up to melting 

temperature 

T m Ni = 1454°C 

T m Al = 660°C 

T m NiAl = 1638°C 

25 


∂ 

∂ 

c 

t 

= D⋅ 

Δc 

solution to these 

boundary conditions: 

Second Fick's Law 

Can be concluded directly from first Fick's law. 

Similar in heat transfer systems, electrical 

potential, ... . 

( c − c ) 

c( x, 

t) 

= c1 

− 1 0 

⎛ 

⋅Γ⎜ 

⎜ 

⎝ 2 

x ⎞ 

⎟ 

D t ⎟ 

⎠ 

0.5 1 1.5 2 

26 


1 

0.8 

0.6 

0.4 

0.2 

f 3 (x) 

f 2 (x) 

f1( x) 

f 1 (x) 

= 1− 

Γ 

f2 ( x) 

( x) 

f3( x) 

⎛ 

= 1− 

Γ⎜ 

⎝ 

⎛ 

= 1− 

Γ⎜ 

⎝ 

x 

0. 

5 

⎞ 

⎟ 

⎠ 

x 

0. 

05 

⎞ 

⎟ 

⎠

Thermal Conductivity 

Most simple, stationary case: no heat radiation, constant 

temperatures in front and back of component. 

α heat transfer coefficient 

λ coeff. heat conductivity 

λ = a · c p · ρ 

a coeff. temperature conductivity 

c p heat capacity 

ρ density 

27 


Temperature Distribution with 

Thermal Barrrier Coating (TBC) 

hot air 

cooling air 

Wärmedämm- TBC Haftvermittlerschicht bond coat Grundwerkstoff substrate 

schicht 

28 


material/property 

ferritic steel 

austenite steel 

Ni-base alloys 

Mo 

Ti alloys (α-rich) 

Al 

Al 2 O 3 

Thermal Conductivity 

λ 

⎡ W ⎤ 

⎢ 

⎣m 

⋅ K ⎥ 

⎦ 

45 

15 

11 

145 

7 

210 

25 

c p 

⎡ J ⎤ 

⎢ ⎥ 

⎣kg 

⋅ K ⎦ 

⎡ g ⎤ 

⎢ 3 

⎣cm 

⎥ 

⎦ 

29 


460 

500 

450 

240 

530 

890 

800 

ρ 

7.8 

8.0 

8.2 

10.2 

4.5 

2.7 

3.9 

⎡ −6 

⎢10 

⎣ 

a 

m 

s 

2 

⎤ 

⎥ 

⎦ 

13.0 

3.8 

3.0 

59.0 

2.9 

87.0 

8.4

Contents 




a) Static 





30 


Microstructure is NOT stable 

annealed deformed 

stress-relieved recrystallized 

31 


Recrystallization 

time dependence of 

recrystallization can be 

approximated by 

Avrami-Johnson-Mehl 

function: 

32 


f 

r 

= 

1− 

e 

− 

⎛ 

⎜ 

⎝ 

t 

t0 

⎞ 

⎟ 

⎠ 

n

Grain Coarsening 

• driving force: reduction of grain boundary 

energy 

• T > 0.7 · Tm • no pre-deformation necessary 

• self-similar system 

• Ostwald ripening d ~ t1/3 (big grains eat up 

small grains) 

• new grains have low dislocation density 

33 


Grain Coarsening 

monomodal 

bimodal (some grain 

boundaries are pinned, 

e.g. by precipitates) 

34 


Precipitate Hardening 

Requirements: 

• solid solution at higher 

temperatures (ability to 

homogenization heat 

treatment) 

• during cooling a two-phase 

region should be reached 

• in general: cooling rate as 

high as possible, thereafter 

annealing (in the two-phase 

region) to let grow the 

precipitates 

35 


Thermodynamic ↔ Kinetic 

36 


Example: Al-Cu Alloy 

37 


Time Dependence of 

Precipitation Hardening 

nucleation, growth, coarsening 

d T precipitate size λ T distance between precipitates 

f T volume fraction of precipitates 

T = const. 

38 


misfit δ 

Coherent - Semicoherent - Incoherent 

δ 

: = 

1 

2 

a 

T 

− a 

( a + a ) a a a 

T 

M 

M 

39 


≈ 

a 

T 

− a 

M 

M 

≈ 

a 

T 

− a 

T 

M 

≈ 

Δa

Energy Consideration 

ΔG total = ΔG vol + ΔG boundary + ΔG strain + ΔG defect 

total change in free enthalpy 

strain enthalpy (elastic energy + dislocation line energy) 

reduction of enthalpy by precipitation coupled with a defect 

enthalpy of phase boundary (scales with surface) 

enthalpy of formation of matrix to precipitate (scales with volume) 

40 


Heterogeneous Nucleation 

41 


TEM-Micrograph of TiC Precipitates at 

Dislocations in an Austenitic Steel 

42 


Ostwald-Ripening of Precipitates 

d 3 -d 0 3 ~D⋅t here for T > Tm 

γ' particle size in IN 738 LC at 

T = 920°C. 

particle coarsening constant of 

(50 nm) 3 /h 

43 


Contents 




a) Static 





44 


Room Temperature (RT) 

High Temperature (HT) 

Deformation 

• most alloy properties at room temperature are 

time and rate independent (elastic constants, 

tension stress, ... ), tension stress experiment. 

• For T > 0.4 · Tm the properties (deformation) will 

be time and rate dependent, creep experiment. 

Kaltverfomung 

Kriechen 

Verformungsverfestigung 

stark 

zeitlich begrenzte 

Festigkeitssteigerung, 

reduziert Zeitstandfestigkeit, 

kann Rekristallisation zu 

Feinkorn bewirken 

Feinkornhärtung 

Festigkeitsabnahme bei 

feinem Korn 

� Grobkorngefüge 

erforderlich 

mittel bis stark 

mittel bis stark 

45 


mittel 

Mischkristallverfestigung 

mittel 

Teilchenhärtung 

mittel bis stark

Change in Materials Properties 

with Temperature 

Material properties of steel and 

Ni-alloys at elevated 

temperatures. Comparison 

between short-term and longterm 

parameters. 

46 


shear modulus G 

Elastic (E-)Modulus and 

Poisson's Ratio 

G 

= 

E 

2⋅ 

( 1+ 

ν) 

47 


High Temperature Deformation 

• dislocation glide (Peierls stress, in fcc and hcp very small and for T > 

0.15 Tm negligible) 

• cross slip of screw dislocations and dislocation interactions (for a low 

stacking fault energy � larger dislocation spacing � thermal 

activation necessary, T > 0.2 Tm , influence on deformation rate) 

• climb of edge dislocations to overcome obstacles: 

diffusion at complete 

dislocation line 

� T > 0.4 Tm 48 


Dislocation Climb 

climb of edge dislocations to 

annihilate each other. 

arrangement in low energy 

configurations (sub-grain 

boundaries), climbing around 

abstacles (leaving the glide 

plane) 

movement of screw 

dislocations with kink 

49 


Back Stress 

Dislocations climb allows annihilation of 

dislocations and to establish a constant internal 

back stress of: 

σ = α⋅ 

G ⋅b 

⋅ ρ 

σ dislocation = and 

i 

G ⋅ b 1 

⋅ 

2⋅ 

π r 

G shear modulus, α constant 0.3 - 1, b 

magnitude of Burgers vector 

50 


ρ 

= 

1 

r

Creep Experiment 

behavior of pure metals: 

primary secondary tertiary: 

51 


Creep Experimental Setup 

up to 1400°C 

Constant 

temperature 

and stress 

52 


Creep Experimental Setup for 

Electrical Conductivity Material 

up to Melting Temperature 

Pyrometer from left, optical strain 

measurement from right, both contact-free. 

53 


strain [%] 

7 

6 

5 

4 

3 

2 

1 

0 

Interrupted creep tests 

single crystal (SX) nickel base superalloy (habilitation thesis Glatzel) 

strain rate [1/s] 

[001] orientation, 1123K, 650MPa 

0 10 20 30 40 50 60 70 

10 -5 

10 -6 

10 -7 

time [h] 

1123K, 650 MPa 

[001] orientation, 1123K, 650MPa 

0 10 20 30 40 50 60 70 

time [h] 

0 1 2 3 4 5 6 

University Bayreuth, Advanced High 

strain 

Temperature 

[%] 

Alloys 54 

Uwe Glatzel, Metals and Alloys 

strain rate [1/s] 

8x10 -6 

6x10 -6 

4x10 -6 

2x10 -6 

0 

logarithm of strain rate versus strain 

(most valuable information for 

materials scientist)

Different Creep Stages 

• primary creep: strain rate dε/dt decreases � 

material hardens 

• secondary creep stage: strain rate constant � 

hardening and softening are in equilibrium � 

dislocation multiplication and annihilation in 

equilibrium � disl. density ρ = const. 

• tertiary creep: necking (creep pores) develop 

� local stress and strain rate increases 

drastically. 

55 


Problem with Low Creep Rates 

Life time of stationary gas turbines > 20 years. Assuming a 

maximum deformation of 3%, this leads to a strain rate of 

about 5·10 -11 s -1 . Reliable data in labs can only be obtained 

down to 5·10 -9 s -1 (1 μm change with l 0 = 25 mm after 2 h, 

one creep experiment per year!). 

Therefore within university labs we are two and more 

orders of magnitude too fast than real life in a stationary 

gas turbine! 

56 


aw data creep curves: 

Engineering Creep Curves 

time to failure: 

isochrone time to failure: 

� time - strain 

� isochrone strain 

57 


Tension ↔ Creep Experiment 

58 


ε& 

steady state 

ρ 

≈ 

⎛ 

⎜ 

⎝ 

Natural Creep Law 

= ρ⋅ 

b ⋅ v 

σ 

G ⋅ b 

v ~ σ 

external 

⎞ 

⎟ 

⎠ 

1 

external 

ε& 

2 

3 

σ 

G 

external ~ 2 

� natural creep law 

⋅b 

59 


ε& 

= 

A ⋅σ 

Norton Creep Law (Empirical) 

n 

external 

power law break 

down (plb) 

T = const. 

diffusional creep 

⋅e 

−Q 

creep 

R⋅T 

dislocation 

climb 

with the Norton creep exponent "n" and 

Qcreep ≈ Qself diffusion 

stress dependence: 

T = const. 

60 


ε& 

ε& 

Diffusional Creep 

• Nabarro-Hering creep (pure volume diffusion) 

NH 

= 

D 

self diffusion 

2 2 

h 

σ⋅ 

Ω 

k ⋅T 

• Coble creep (grain boundary diff.) 

C 

= 

δ⋅D 

grain boundary 

2 3 

h 

σ⋅ 

Ω 

k⋅T 

h grain size, δ thickness of grain boundary 

61 


ε& = ε& 

+ ε& 

diffusion creep 

Combined NH and Coble Creep: 

NH 

C 

D 

Dself 

diffusion Dgrain 

boundary D 

2 ~ ⋅ 

2 

3 

k T h 

h kT 

h 

Ω σ 

σ⋅ 

Ω ⎛ π⋅ 

δ⋅ 

⎞ 

= ⋅ ⋅ ⎜ + 

⎟ 

⋅ ⎝ 

⎠ 

eff 

= 

D 

self 

diffusion 

real geometry (non-cuboidal grains) 

+ 

π⋅ 

δ⋅ 

62 


D 

grain boundary 

h 

eff 

2

Temperature Dependence of 

Stationary Creep Rate 

σ = 28 MPA = const. 

Austenitischer Stahl 800H 

fcc alloys: 

63 


ε& 

s 

= A ⋅ γ 

3, 

5 

SF 

⎛ σ ⎞ 

⋅ ⎜ ⎟ 

⎝ E ⎠ 

n 

⋅ e 

−Qc 

R⋅T

Activation Energy for Creep 

slope = 1 

64 


Constant Load ↔ Constant Stress 

n 

n 

n ⎛ F ⎞ ⎛ F⋅ 

( 1+ 

ε) 

⎞ 

ε& = ε& 

0 ⋅σ 

= ε& 

0 ⋅⎜ 

⎟ = ε& 

0 ⋅ 

= ε0 

A ⎜ 

A ⎟ & 

⎝ ⎠ 

0 

failure 

ln ε& = ln 0+ 

n · ln σ0 + n · ln (1+ε) = const. + n · ln (1+ε) 

ε& 

( ) n 

1+ 

ε 

in case the gauge length 

deforms uniform with 

constant volume 

65 


⎝ 

⎠ 

⋅σ 

As method to determine n only 

applicable if secondary creep 

state lasts to at least 10% 

n 

0 

⋅

Ashby Deformation 

Mechanism Maps 

66 


Ashby Deformation 

Mechanism Maps 

67 


Deformation Mechanisms: 

Elastic Deformation: Spontaneous and reversible deformation. In the elastic region: σ = E·ε (rule of 

thumb: εe, max ≈10-3 , but definitely 0.4⋅Tm ) and lower stress levels dislocation climb plays the 

major role => time dependent constant strain rate (dε/dt) ss ~ σn , with a Norton stress exponent in-between 3 

und 8. 

Diffusional Creep: In-between 0 K und 0.8⋅Tm and very low stress levels: Coble-creep (grain boundary 

diffusion). Below 0.4⋅Tm not measurable. For geological times a time dependent deformation can be 

determined. Transition to Nabarro-Herring creep (volume diffusion) is dependent on grain size and grain 

boundary thickness. The transition temperature from coble to Nabarro-Herring creep can be explained by 

the different activation energies of volume and grain boundary diffusion. 

68 


a) interaction dislocation 

and impurity (low temp.) 

b) stationary dislocation 

pinned by impurities 

(Cottrell clouds) 

c) pulled off Cortrell clouds 

(Lüders bands) 

d) gliding dislocation trails 

impurities behind (viscous glide) 

Creep of Alloys 

e) impurities faster than dislocation (very high temp., no hardening) 

f) annihilation due to dislocation climb 

solid solution 

69 


σ 

i 

= α⋅ 

G ⋅b 

⋅ ρ + σ

σ 

i 

Precipitation Hardening 

= α⋅ 

G ⋅b 

⋅ ρ + σ + σ 

solid solution 

precipitate 

threshold stress concept (with n ≈ 3 - 4 and Q creep = Q self diffusion ): 

cutting 

mechanism 

bypass by Orowan 

climb over obstacles 

ε& 

ss 

⎛ σ − σ 

= A ⋅ ⎜ 

⎝ E 

temperature 

0 K up to T s 

0 K up to T s 

> 0.4⋅T s 

70 


0 

⎞ 

⎟ 

⎠ 

n 

⋅ e 

−Qc 

R⋅T 

coherent and semicoherent 

phase 

boundaries 

yes 

yes 

yes 

in-coherent phase 

boundaries 

no 

yes 

no

Hardening Mechanisms as 

Function of Precipitate Size 

d T0 initial precipitate size 

σ 1 and σ 2 arbitrary stress levels 

passing by: 

climbing: 

ε& ~ dT 

71 


ε& 

~ 

1 

2 

dT

Pinning of Dislocations by 

Carbides in Austenitic Steel 

T = 1000°C, σ = 25 MPa, carbides of the type TiC und M 23 C 6 

72 


Very High Volume Fractions 

Volume fractions of 70% are only achievable with non-spherical precipitates. 

Spacing between precipitates is getting smaller � Orowan stress 

σOrowan ≈ G·b/L necessary. For small strains precipitates are not cut by 

dislocations. With G = 90 GPa, b = 0.25 nm, L ≈ 75 nm => σOrowan ≈ 300 MPa 

nickel base superalloys 

ODS alloys: 

σ Orowan ≈ 

G ⋅b 

⋅ 

73 


d 

part. 

f 

vol.

yield 

stress 

Dispersion Hardening 

precipitate strengthened 

(ODS-Alloys) 

dispersion strengthened 

temperature 

back-side pinning of dislocation by 

ODS-particle (Rössler + Arzt) 

74 


Creep Damage 

a) cracks at grain boundaries b) cavities (micropores) at grain boundaries 

75 


Creep Damage 

nucleation, not detectable with OM 

micropore, difficult to detect 

fracture 

micro cracks 

pear necklace like chain of 

micropores (easy detectable) 

76 


t 

f 

Extrapolation of Time-to-Fracture Data 

(Larson-Miller plot, Larson-Miller parameter) 

Monkmann-Grant relation with constant K and exponent m ≈ 1: 

K 

= 

ε& 

m 

ss 

Qcreep 

− 

R⋅T 

ε& ss = B⋅ 

e or: 

ln( t f ) = K − m ⋅B1 

− m ⋅B 

2 ⋅ 

or: ln( t ) = K − m ⋅ln( 

ε& 

) 

77 


f 

ln( ε& ss) 

= B1 

− B2 

⋅ 

1 

T 

= −C 

+ P ⋅ 

with material dependent constants C and P 

1 

T 

ss 

1 

T

Comparison of CMSX-6, 

LEK 94 and CMSX-4, 

patent Wöllmer, Glatzel, 

Mack, Wortmann 

Larson-Miller-Plot 

P=T⋅[20 + ln(t f )]⋅10 -3 (T in K, t f in h) 

78 


stress [MPa] 

500 

230 

120 

Comparison LEK 94 with 

CMSX-4 and CMSX-6 

24 K 

ΔT = 10 K 

10 K 

Larsen-Miller-parameter 

P = T (20+log tB ) 10 -3 

25 26 27 28 29 30 31 32 

CMSX-6 [Wortmann 88] 8.0 g/cm 3 

CMSX-4 [Erickson 94] 8.7 g/cm 3 

CMSX-4 [Frasier 90] 8.7 g/cm 3 

LEK-2 8.5 g/cm 3 

LEK-4 8.2 g/cm 3 

LEK-5 8.2 g/cm 3 

LEK-3 8.1 g/cm 3 

LEK-6 8.3 g/cm 3 

LEK-1C 8.4 g/cm 3 

LEK-1B 8.3 g/cm 3 

LEK-1A 8.2 g/cm 3 

29 K 

Not corrected 

regarding density! 

79 


Contents 




a) Static 





80 


Time Dependent Variation of Stress 

and/or Temperature and/or ... 

Wöhler diagram for T < 0.4·Tm . Z time fatigue limit, D endurance 

fatigue limit 

a) type I metal (bcc) b) type II metal (fcc) endurance limit at 2·10 7 

81 


Change in Wöhler Diagram with 

Temperature and Holding Time 

82 


Thermal Fatigue 

Thermal breathing of turbine blade: 

a) heating phase: edges reach high temperatures faster than interior 

b) cooling phase: edges cool faster than interior 

c) repeated thermal cycles lead to thermal fatigue cracks at edges 

83 


Thermal Strains and Stresses : 

ε thermal = α thermal · ΔT, or: σ thermal = E · ε thermal 

� σ thermal = E · α thermal · ΔT 

84 


Lower E-Modulus is Helpful: 

� orientation of single crystals in direction reduces thermal stresses 

85 


TMF and many other Time 

Dependent Test Techniques 

Can not be covered in this lecture! 

86 


Contents 




a) Static 





87 


High Temperature Corrosion 

• oxidation: external and internal, passivation 

• carburization (internal carbides) 

• nitration: internal, seldom nitrite passivation 

• sulfurization: external (sometimes 

passivation), seldom internal 

Worldwide 1 ton iron per minute corrodes to rust (low 

temperature aqueous corrosion). 

88 


Ellingham-Richardson-Diagram 

right hand and lower axes 

� O2 partial pressure at T = 0. 

As an example p O2 of 

10 -15 Pa = 10 -20 bar = 10 -17 mbar 

is shown as a dashed line. 

only the oxides below this line 

are thermodynamic stable. 

89 


Time Dependent Oxidation 

90 


Oxidation Mechanisms 

• logarithmic (not shown) � low temperature oxidation 

which eventually comes to a stop or no measurable increase 

in oxide scale thickness (e.g. Al, Cr, Mg). 

• parabolic mass change (Δm/A) 2 ~ t. Diffusion through 

oxidation layer (either oxygen or metal). Most favorable 

oxidation behavior. 

• linear mass change: oxide layer with cracks � continuous 

contact with metal (e.g. Ta, Nb). 

• mass loss: volatile oxides � catastrophic oxidation (e.g. V, 

Mo, W, Cr, Pt). You can see it inside a broken light bulb. 

91 


Pilling-Bedworth Ratio 

PB = (volume of oxide of one metal atom)/(volume of metal atom) 

Oxide 

PB 

MgO 

0.81 

Al 2 O 3 

1.28 

ZrO 2 

1.56 

NiO 

1.65 

FeO 

1.70 

TiO 2 

1.73 

92 


CoO 

1.86 

ideal is 1.1 to 1.3 

Cr 2 O 3 

2.05 

FeCr 2 O 4 

2.10 

SiO 2 

2.15 

Ta 2 O 5 

2.50 

Nb 2 O 5 

Of course thermal expansion coefficients also play a major role for the stability of oxide scales. 

2.68

Alloying Effects: 

different elements have 

different oxygen affinity 

concentration changes 

diffusion rates are different 

oxide layer contains other 

metals 

93 


Example Ni-Cr-Al 

� Ni Cr 10 Al 5 

oxide layer and 

internal 

oxidation occurs 

94 


Contents 




a) Static 





95 


High Temperature Alloys 

T > 500°C, Application in: 

• energy generation 

• engines (cars, trains, airplanes, ships, ... ) 

• chemical industry 

• metallurgy 

• mechanical engineering 

96 


ele 

m. 

Ti 

V 

Cr 

Mo 

structure 

α 

hdp 

β krz 

krz 

krz 

krz 

T trans. T m 

[°C] 

882 

1855 

1910 

1863 

2623 

Overview Metals 

ρ 

[g/cm 3 ] 

4.5 

4.5 

6.1 

7.2 

10.2 

max. O-solubility 

[at.%] 

31.9 

8 

17 

0.0053 

0.03 

advantages/disadvantages 

+ low density 

+ high melting point 

+ abundant available 

+low α th. (~ 10 -5 K -1 ) 

− now alloy known with adequate strength for temperatures > 600°C 

− high oxygen and nitrogen solubility > 700°C, increased brittleness 

− linear oxidation > 800°C 

− low thermal conductivity 

− ignition hazard 

− catastrophic oxidation; T m (V 2 O 5 ) = 658°C 

− very brittle at RT; conventionally not processable 

+ very high creep strength 

+lowα th , high thermal conductivity, good thermal fatigue strength 

− very brittle at RT 

− catastrophic oxidation; T m (MoO 5 ) = 795°C 

− no long lasting coating available 

W krz 3422 19.3 

≈ 0 + highest melting point of metals (only C with even higher Tm) + very high creep strength 

+low α , high thermal conductivity, good thermal fatigue strength 

th 

− very brittle at RT 

− catastrophic oxidation > 1000°C durch hohe WO -Abdampfrate 

3 

− no long lasting coating available 

University Bayreuth, Advanced High Temperature Alloys 97− 

very high density 

Uwe Glatzel, Metals and Alloys

elem. 

Fe 

Co 

Ni 

Pt 

structure 

α krz 

γ kfz 

δ krz 

ε hdp 

α kfz 

kfz 

kfz 

T trans. 

T m 

[°C] 

912 

1395 

1538 

422 

1495 

1455 

1772 

Overview Metals 

ρ 

[g/cm 3 ] 

7.9 

7.7 

7.4 

8.8 

8.7 

8.9 

21.5 

max. Osolubility 

[at.%] 

0.0008 

0.0098 

0.029 

≈ 0 

0.048 

0.05 

≈ 0 

advantages/disadvantages 

+ very good corrosion resistance by alloying with Cr or (Cr + Al) 

+ γ-structure can be stabilized down to RT (by Ni) 

+ very good processable and weldable 

+ low cost (~ 1 €/kg) 

− strength at high temperatures (> 700°C) limited 


+ Co-alloys castable in air good weldability 

− only moderate hardening available 

− Ni-additions necessary to stabilize fcc structure, reduces strength 

+ broad possibilities for alloying, high strength increase possible 


+ processable 

− relatively low melting point 

−α th. high, low thermal conductivity 

+ high corrosion and oxidation resistance 


− very high density 

− very expensive (~ 33 €/g) 

98 


10.000 h Life Time 

99 


Example of Intermetallic 

Phases (Ni-Al-System) 

100 


phase 

Ni 3 Al 

NiAl 

structure 

L1 2 

L1 0 

Ni-Al Intermetallic Phases 

T trans. 

T m 

[°C] 

1383 

1638 

ρ 

[g/cm 3 ] 

7.5 

5.85 

advantages/disdavantages 

+ anomalous temperature dependence of strength 

+ same structure base than Ni matrix (fcc) 

+ stable for larger Al variations > 1 wt.% Al 

+ ductile as single crystal 

− high density 

− brittle as polycrystal (can be hindered by boron doping (grain 

boundary strengthener) 

−Al-content not sufficient to build stable Al 2 O 3 -layer � reduced high 

temperature oxidation resistance 

+ very good oxidation resistance, since 30 wt.% Al 


+ low density 

+ ordered structure up to melting point 

+ high thermal conductivity 

+ low coefficient of thermal expansion 

− extremely brittle at temperatures below 500°C (von Mises criterion 

not fulfilled) 

− low strength at high temperatures 

101 


NiAl, B2 Ordered 

Intermetallic Phase 

• At a first sight very interesting (see 

advantages) but despite many efforts and many 

100 Mio. US$ research money spent, up today 

no bulk usage of NiAl has been achieved. 

• BUT: aluminum coatings leading to NiAl 

layers is heavily used. 

102 


Contents 




a) Static 





103 


MTS-Fabrik Wolfsbach 

Spatenstich: 20.02.2008 Richtfest: 06.06.2008 

Start der Produktion: ~ 12/2008 

104 


MTS-Fabrik, Juni 2008 

105 


MTS-Fabrik, Juni 2008 

106 


Herstellungsprozess einer 

Turbinenschaufel 

FPI 

X-Ray 

Feinguss, Wachsausschmelzverfahren, lost wax investment casting, ... 

107 


Archäologischer Fund (Bibracte) 

datiert auf ~ 50 v.Chr. 

Mit Wachs gefüllte Gießform 

Kleiderspange 

108 


Projekt der Bayerischen Forschungs- stiftung 

(BFS) und der Oberfrankenstiftung (OFS) 

Sprecher: 

109 


Zahlen 

• 3.800 m 2 Fabrik, 10.400 m 2 Grundstück, 

~ 25 Mio. € Investitionsvolumen. 

Identische, zweite Fabrik in Planung. 

• Finanzvolumen des BFS/OFS-Projekts: 

2,8 Mio. €, Fördersumme 

�1 Mio € BFS, 0,4 Mio. € OFS. 

• Bei den Forschungsstellen Einstellung von 

5 wissenschaftliche Mitarbeiter/innen und 

2 Techniker/innen über 3 Jahre. 

110

Advanced High Temperature Alloys

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