Opera&onal  Performance  Summary:  •  Our  research  will  revolu/onize  the  efficiency  ac/ve,  high  power  and  

temperature  sensi/ve  devices  from  nano  to  bulk  scales,  including  high  frequency  transistors  thermoelectric  devices  and  opto-­‐electronic  devices  for  RADAR,  pulsed  power,  compu&ng,  and  energy  harves&ng  

Connec&on  to  Na&onal  Defense/Intelligence:  •  We  are  improving  energy  efficiency  in  current  and  future  applica/ons  

of  importance  to  na/onal  security  and  defense;  our  research  is  supported  by  AFOSR,  ARO,  ONR    

 

Products  and  Developments:  •  Compact  thermal  conduc/vity  diagnos/c/measurement  system:  from  

lab  to  package  

Technical  Approach:  Pump-­‐probe  spectroscopy  using  a  range  of  femtosecond  to  con/nuous  wave  lasers  to  measure  thermal  conduc/vity  and  thermal  boundary  conductance  from  bulk-­‐  to  nano-­‐scales  

 

Summary:  Iden/fy  energy  efficient  materials  on  nano-­‐to-­‐bulk  length  scales  by  manipula/ng  atomic  proper/es  and  thermal  conduc/vity   Conclusion/Future  Direc&on:  

•  Engineer  novel  materials  at  atomic  scale  to  set  new  records  and  extremes  of  thermal  conduc/vity  

•  Material  engineering  for  material  stability  and  durability  in  harsh,  corrosive,  high  impact,  extreme  environments  

•  Improved  efficiency  of  turbine  engines,  opto-­‐  and  nano-­‐electronic  devices  and  energy  harves/ng  modules  based  on  op/mizing  opera/onal  temperature  and  thermal  management  

•  Real  /me  diagnos/cs  of  high  frequency  (RF),  high  power  device  thermal  breakdown  and  failure    

•  Take  technology  out  of  lab  and  into  field,  on-­‐chips  and  on-­‐soldiers  •  Tac/cal  sensing  of  reac/ons/chemicals  in  the  field  

Key  Center  at    U.Va.  And  Collabora&ons:  U.Va.  Thermal  Management  and  Diagnos/cs  Group  

Hopkins  (MAE),  Haj-­‐Hariri  (MAE),  Norris  (MAE),  Ghosh  (ECE),    Stan  (ECE),  Skadron  (CS),  Zhigilei  (MSE),  Poon  (PHYS)  

     

SP#Tsunami#3.0#W,#80#MHz#90#fs#pulse#width#

Isolator#λ/4#

E.O.#Modulator#

BiBO#

λ/2#

Delay#line#(~7#ns)#

Red#filter#

Dichroic#

Blue#filter#

LockQin#amplifier#

Photodiode#

Camera#to#image#sample# Pump%

Probe%

Semi-infinite substrate

Nanostructure

Thin metal film transducer

Pump

Probe

Thermal penetration depthℓ

temperature between 180 and 387 K and largely insensitiveto film thickness in the range 22 to 106 nm as shown inFig. 3(b). We also note that the chosen substrate (additionalfilms were deposited on glass and silicon, as opposed to theITO and PEDOT:PSS coated glass slides described above)or heat treatment (annealed or unannealed) did not lead tostatistically significant changes in thermal conductivity.

In Ref. [17], Olson and Pohl used low temperatureheat capacity measurements to determine the Einsteintemperature of C60=C70 fullerite microcrystals, !E ¼35 K, which corresponds to a frequency of kB!E=@ ¼4:58" 1012 rad s#1, where @ is Planck’s constant dividedby 2!. With this value and the Einstein model of thermalconductivity,

"E ¼ 2k2B@ N1=3

!!E

x2ex

ðex # 1Þ2 ; (2)

whereN is the fullerene density and x ¼ !E=T, they foundexcellent agreement between the model and their data.Following the reverse procedure and fitting the Einsteinmodel of thermal conductivity to our temperature-dependent thermal conductivity data yields !E ¼ 22 K,which corresponds to a frequency of 2:88" 1012 rad s#1.This suggests that the presence of the molecular tail is notonly responsible for lowering the sound speeds of PCBMmicrocrystals, but also lowering the characteristic fre-quency of their highly localized vibrations.

To put the exceptionally low thermal conductivity ofPCBM into perspective, in Fig. 4, we plot the room-temperature thermal conductivities of several amorphousand crystalline materials as a function of their atomicdensity. While previous reports have made similar com-parisons with regard to mass density [6], plotting thermalconductivity as a function of atomic density allows easieridentification of trends among crystalline and amorphousmaterials, respectively. The outliers (P3HT, C60=C70, andPCBM) are nominally microcrystalline, exhibit some ofthe highest atomic densities, and simultaneously, some ofthe lowest conductivities. In this respect, it is interesting tonote that some of the best thermal conductors, as well asthe best thermal insulators, are carbon allotropes or carbonbased materials [37].In summary, we have reported on the thermal conduc-

tivities of [6,6]-phenyl C61-butyric acid methyl ester(PCBM) thin films from 135 to 387 as measured by timedomain thermoreflectance. Thermal conductivities wereshown to be independent of temperature above 180 Kand <0:030& 0:003 Wm#1 K#1 at room temperature.The longitudinal sound speed as measured by picosecondacoustics was 2300& 100 m s#1, 30% lower than that inC60=C70 fullerite compacts. Using Einstein’s modelof thermal conductivity, we found the Einstein character-istic frequency of microcrystalline PCBM is 2:88"1012 rad s#1. Through a comparison of our data to previousreports on C60=C70 fullerite compacts, we have argued thatthe molecular tails on the fullerene moieties in our PCBMfilms are responsible for lowering both the apparent soundspeeds and characteristic vibrational frequencies belowthose of fullerene films. In turn, the room-temperaturethermal conductivities of PCBM thin films are the lowestreported of any fully dense solid.J. C. D. and P. E. H. acknowledge funding from the

National Science Foundation (CBET-1134311). Y. S. andM.C.G. would like to thank the NASA Langley ProfessorProgram, NSF I/UCRC program, and the University ofVirginia Energy Initiative for financial support. Thiswork was supported in part by the Laboratory DirectedResearch and Development (LDRD) program at SandiaNational Laboratories.

*phopkins@virginia.edu†mgupta@virginia.edu

[1] D. G. Cahill, MRS Bull. 37, 855 (2012).[2] A. Einstein, Ann. Phys. (Berlin) 340, 679 (1911).[3] P. Debye, Ann. Phys. (Berlin) 344 789 (1912).[4] C. Kittel, Introduction to Solid State Physics (Wiley,

Hoboken, New Jersey, 2005), 8th ed.[5] W. Kim, R. Wang, and A. Majumdar, Nano Today 2, 40

(2007).[6] K. E. Goodson, Science 315, 342 (2007).[7] D. G. Cahill and R.O. Pohl, Annu. Rev. Phys. Chem. 39,

93 (1988).

1022 1023

0.01

0.1

1

10

100

1000

Atomic Density (cm-3)

The

rmal

Con

duct

ivity

(W m

-1 K

-1)

AmorphousCrystalline

DiamondCopper

AluminumSilicon

GermaniumLead

SiO2

Aerogels PCBM

C60/C70WSe2

α:carbonP3HT

5000

0.005

FIG. 4 (color online). Room-temperature thermal conductivityof various materials plotted as a function of their atomic density.The values for diamond, copper, aluminum, silicon, germanium,and lead are from Ref. [31], SiO2 and aerogels from Ref. [16],amorphous carbon from Ref. [35], WSe2 from Ref. [18],C60=C70 from Ref. [17], P3HT from Ref. [38] and PCBM isfrom the present work. Not only does PCBM exhibit the lowestconductivity, but it is among the densest of the materials, secondonly to diamond.

PRL 110, 015902 (2013) P HY S I CA L R EV I EW LE T T E R Sweek ending

4 JANUARY 2013

015902-4

THERMAL-BARRIER COATINGS FOR MORE EFFICIENT GAS-TURBINE ENGINES

892 MRS BULLETIN • VOLUME 37 • OCTOBER 2012 • www.mrs.org/bulletin

the underlying metal parts, any TBC failure can endanger the engine. 14 Furthermore, because of the coupled diffusional and mechanical interactions between the oxide ceramic coating and the underlying alloys at these high temperatures, it is essential to consider TBCs as a complex, interrelated, and evolving mate-rial system, consisting not only of the oxide ceramic coating (topcoat) itself but also the underlying superalloy engine part, and two other layers in between. These include a metallic bond-coat layer that is more oxidation resistant than the superalloy, and a thin, thermally grown oxide (TGO) layer that forms between the topcoat and the bond coat as result of bond-coat oxidation in-service. The bond-coat composition is designed to result in a TGO made of α -Al 2 O 3 —a mechanically robust, effective barrier to oxygen diffusion. Figure 4 illustrates this multilayer structure in a typical TBC system.

During service, several kinetic processes occur in parallel. Interdiffusion between the bond coat and the underlying super-alloy occurs, driven by chemical potential gradients; Al diffuses from the bond coat to form the TGO; and microstructural,

chemical, and phase changes occur in all the materials, including in the ceramic topcoat itself, changing their very properties. Since all of these are thermally activated processes, the rates at which they occur are expected to increase exponentially with temperature, albeit with different activation energies. Furthermore, the processes generally lead to degradation and failure of the coating.

TBCs are also multifunctional: they must provide thermal insulation to protect the under-lying superalloy engine parts, have strain compli-ance to minimize thermal-expansion-mismatch stresses with the superalloy parts on heating and cooling, and must also refl ect much of the radiant heat from the hot gas, preventing it from reaching the metal alloy. Furthermore, TBCs must maintain thermal protection for prolonged service times and thermal cycles without failure. Typically, these times are 1000s of hours for jet engines being cycled numerous times between a maximum temperature of ∼ 1300°C and room temperature (takeoff/landing and on-ground), and 10,000s of hours for power-generation engines with fewer thermal cycles (maintenance shut-downs). However, the latter are now being increasingly employed to stabilize the electric grid connected to renewable sources (wind, solar), and, thus, these engines experience more frequent cycles to compensate for the inherent intermittency of renewables. TBCs need to do this without separating from the engine parts while also withstanding extreme thermal gradients ( ∼ 1°C µ m –1 ) and energy fl uxes ( ∼ 1 MW m –2 ). Not only are these demands extremely exacting but also are often confl icting: TBCs must have both

low thermal conductivity and low weight; they must remain intact while withstanding large stress variations, both due to heating and cooling as well as under thermal shock; they must be chemically compatible with the underlying metal and the TGO; and they must operate in an oxidizing environment at maximum pressures of ∼ 10 atmospheres and maximum gas velocities exceeding Mach 1.

These demands and the desire to operate at higher tempera-tures reliably for longer times are driving new TBC innovations based on exploration of the underlying materials, processing sciences, and mechanistic understanding of degradation/failure and its mitigation. Several but not all of these key areas are highlighted in this issue of MRS Bulletin .

In this issue Ceramic topcoat processing Sampath et al. describe the oxide ceramic topcoat deposition processes and microstructures. Unlike more traditional thin fi lms used in microelectronics and in materials-growth studies,

Figure 1. Cutaway view of Engine Alliance GP7200 aircraft engine, photograph of a turbine blade ( ∼ 10 cm long) with thermal-barrier coating (TBC) from the high-pressure hot section of an engine, and a scanning electron microscope (SEM) image of a cross-section of an electron beam physical vapor deposited 7 wt% yttria-stabilized zirconia TBC. (Engine image courtesy of Engine Alliance, turbine blade photograph courtesy of YXLON, and the SEM micrograph is from Reference 44.) TGO, thermally grown oxide.

!'

Figure'3:''Diagram'of'the'chemical'solution'deposition'(CSD)'process.'''Solution'precursors'mixed'via'conventional'wet'chemistry'techniques'are'deposited'on'a'substrate'via'either'spin'or'dip'coating.''The'coated'substrates'are'then'pyrolyzed'at'low'temperatures'for'a'specific'amount'of'time'to'remove'the'organic'components'of'the'precursor'solution,'leaving'behind'

the'desired'components'of'the'target'stoichiometry'in'a'semi>crystallized'state.''A'final'annealing'step'can'then'be'performed'to'continue'the'crystallization'process'as'desired'(ie,'control'of'grain'and'phase'formation'within'the'material)

While chemical solution deposition is not a new technology, it offers distinct advantages over existing coating technologies that are currently in use. CSD offers a means to produce high-quality films from inexpensive precursors over large areas with film thicknesses anywhere from tens of nanometers to tens of microns [4– 6]. Additionally, CSD allows for two methods of coating [7], spin and dip coating, the latter of which can be used to coat internal features. Currently, the coating of internal features is only achievable through the use of the most advanced deposition systems utilizing directed vapor coating.

Additionally, it has been shown that the incomplete pyrolyzation (removal of organic components within the precursor solution) of liquid precursor in APS coatings leads to the formation of vertical cracks through the porous network in the 7YSZ top-coat, resulting in a high degree of strain tolerance that is desirable for TBCs [3]. The effects of incomplete pyrolyzation on porosity have also been investigated in CSD films [8], suggesting that similar control over cracks and other defect structures to attain desirable attributes may be possible in Strontium Niobate TBCs.

Finally, CSD-fabricated Strontium Niobate films have been shown to form a highly-oriented, well crystalized structure regardless of the substrate that the film is deposited on (ie, poly or single-crystalline substrate, lattice-matched or not) [9]. This fact makes Strontium Niobate extremely versatile as a highly-insulative material layer that can be deposited on almost any film or substrate structure, such as different superalloys or bond coats.

The overarching goal of this proposal work is to assess the commercial potential of multilayer Strontium Niobate for use as a high-performance thermal barrier coating (TBC). Specifically, this work will focus on assessing the validity of chemical solution deposition (CSD) as a scalable fabrication technique for TBCs. We will investigate the connection between synthesis conditions, microstructure and performance of this material as a TBC in high temperature environments.

2

SrTi

O3

5 nm

Sr2N

b 2O

7

n=5

n=4

1 nm

(b)

FIG. 1. (a) X-ray diffraction patterns of samples with 5 % lanthanum(x = 0.1) and the following thicknesses (nm); (i) 130, (ii) 220, (iii)400, (iv) 800. ⇤ denotes reflections from substrate planes and ⇧denotes substrate reflections from Cu K� radiation. (b) Atomic-resolution HAADF-STEM images illustrating the layered crystalstructure of strontium niobate. The grain is imaged along a h101itype direction with the b-axis normal to the SrTiO3 substrate. Arrowsalong the left image indicate the position of defects in the layeringsequence. The image on the right shows a higher magnification of asection of the grain overlaid with h101i projections of the Sr (yellow)and Nb (orange) atom positions in the Sr2Nb2O7 (n = 4 NbO6 octa-hedra) and Sr5Nb5O17 (n = 5 NbO6 octahedra) crystal structures.

substrates at 3000 RPM for 30 seconds and pyrolyzed on ahot plate in air at 300�C for five minutes. Sr2Nb2O7 andSr1.9La0.1Nb2O7��

films were formed by annealing theas-pyrolyzed films to 1000�C in an air atmosphere for 5minutes by directly inserting the samples into a preheatedfurance. After the final deposition and crystallization anneal,the La-containing films were post-annealed in a dry 3%H2/N2 atmosphere to promote solubility of the lanthanum and

to activate electronic carriers for thermoelectric applications.The coating/pyrolysis/crystallization process was repeated toincrease film thickness with as-crystallized individual layerthickness of 27 nm.

The phase-assemblage and orientation of the films was con-firmed via x-ray diffraction (XRD, Philips MPD) using Cu-K↵ radiation, shown in Figure 1(a). XRD reveals highly(0k0)-oriented films for each thickness. Minority peaks as-sociated with the 002 and 151 reflections were also observed,but constituted less than 3% of the diffracting volume basedupon comparison of measured intensities and those expectedfrom a random powder pattern. In previous work [20] thesefilms were shown to possess a random in-plane orientationand therefore are fiber-textured. Also, while we utilize sin-gle crystalline substrates in this study, (010)-fiber texture wasidentified on a non-lattice-matching and polycrystalline sub-strates, demonstrating a preferred (010) out-of-plane orienta-tion regardless of substrate [20].

Scanning transmission electron microscopy (STEM) wasused to confirm the layered structure of the films (seeFig. 1(b)). We employed an FEI-Titan G2 instrument, op-erated at 200 keV and equipped with a high angle annulardark field (HAADF) detector. High-resolution STEM imag-ing highlights the naturally layered crystal structure of thismaterial system and confirms that these layers are well alignedwith the substrate. Figure 1(b) shows detail of an individualSr2Nb2O7 grain that was tilted, in the microscope, to a h101i-type orientation, allowing the layered structure to be imageddirectly. These observations identify the presence of defectsin the layering sequence. In the ideal Sr2Nb2O7 crystal struc-ture, the NbO6-octahedra are arranged in slabs that are 4-octahedra wide along the b-axis [16, 21]. However, more gen-erally within the SrnNbnO3n+2 homologous series (for whichSr2Nb2O7 corresponds to n = 4), slab widths of both 4 and5 NbO6-octahedra have been observed [22, 23]; for instance,the Sr5Nb5O17 structure (the n = 5 member of the homol-ogous series) consists entirely of slabs that are 5 octahedra-wide [24]. Examples of individual n = 4 and n = 5 slabsare indicated on Figure 1(b) with the superimposed Sr andNb atom positions from h101i projections of Sr2Nb2O7 andSr5Nb5O17. The arrows on Figure 1(b) indicate the distribu-tion of n = 5 slabs across one grain. At this defect density (9slabs of the n = 5 phase versus 34 slabs of n = 4 present withinthe left image), we calculate that there is a small amount ofoxygen reduction by � = 0.05 in the Sr2�x

Lax

Nb2O7��

sam-ples. To confirm our results were not sensitive to this low levelof reduction, an additional sample (x = 0.1, approximately220 nm thick) was prepared through CSD, but done so in ox-idizing atmospheres to discourage the formation of the n = 5slabs [20].

The thermal conductivities of the samples were determinedusing TDTR [25] where we fit the experimental data to amulti-layer thermal model [26–28]. To provide the transducerfor our optical measurements, we deposited an aluminum filmon the samples by electron beam evaporation. The film wasapproximately 90 nm thick, as confirmed by picosecond ultra-

Nano%Macro%

Heat%Turbine%Blade% “Nano4bulk”%coa7ng%

Atomic%vibra7ons%

Scale 1 inch

600 800 1000 1200 1400 1600 1800 20000

200

400

600

800

1000

1200

Future trends given current SOA

Inefficiency losses

Spec

ific

Cor

e Po

wer

(kW

/(kg/

s))

Turbine Rotor Inlet Temperature (deg. Celcius)

Ideal performance

Superalloytemperature

Activecoolinggains

Potential for “Nano-bulk”

coatings

Characteristic lengths (m)

Cha

ract

eris

tic ti

mes

(s)

1 nm (10-9)

10 nm (10-8)

100 nm (10-7)

1 µm (10-6)

10 µm (10-5)

100 µm (10-4)

fs (10-15)

ps (10-12)

ns (10-9)

µs (10-6)

> µs

“Hot”&&electrons&

Atomis0c/quantum/coherent&transport&

!

“Hot”&acous0c&phonons&

“Hot”&op0cal&phonons&

Interfacial&sca9ering&and&thermal&boundary&resistance&

Nano/micro&contacts&

Packaging/“sourceAtoAsink”&thermal&management&

Professor  Patrick  E.  Hopkins  Director,  Experiments  and  Simula/ons  in  Thermal  Engineering  (ExSiTE)  Group  

Department  of  Mechanical  and  Aerospace  Engineering,  School  of  Engineering  &  Applied  Science