Naval Research Reviews, a theme-issue quarterly journal,

publishes articles about Department of the Navy Science

and Technology programs conducted by academia,

government laboratories, for-profit and nonprofit

organizations, and industry.

The journal is available on the World Wide Web at

http://www.onr.navy.mil/onr/pubs.htm

Chief of Naval Research

RADM Paul G. Gaffney, II, USN

Executive Director and Technical Director

Dr. Fred Saalfeld

Vice Chief of Naval Research

BGen Timothy Donovan, USMC

Director, Corporate Communications

Liane Young

Publications Coordinator

Cynthia Nishikawa

Scientific Editor

Colin Wood

Managing Editor*

Kathy Parrish

Technical Writer/Editor*

Saul Oresky

Design and DTP*

Jonna Atkinson

Donna Gloystein

Jan D. Morrow

*Naval Research Laboratory

800 North Quincy Street

Arlington, VA 22217-5660

yyz{{|
Cover: Photograph of blue and blue/green LEDs courtesy of Cree Research, Inc.

Colin Wood

Office of Naval Research

From the Guest Editor . . .

aster, cheaper, smaller” is now

the standing order for satellite

and probe development at the

National Aeronautics and Space “F Administration.

development require electronics performance that can

only be delivered by wide bandgap semiconductors.

What is it about these newer semiconductors that

makes them so necessary and attractive? Large

bandgaps mean that leakage and avalanche onset by

generation-recombination currents are insignificant

at normal temperatures. With three-to-tenfold higher

thermal conductivities, heat dissipation is more

efficient, and dielectric constants are approximately

30% lower, thereby reducing parasitic capacitances.

Most importantly, however, is that breakdown voltages

are an order of magnitude higher, allowing higher

voltage swings.

Before these properties can be exploited, however,

several challenges remain to be solved in the area of

materials.

• Silicon carbide (SiC), although it can be grown

in bulk crystals at atmospheric pressure, requires

enormous temperatures, which in turn limits the

size of crystals and leads to many large defects,

known as micropipes.

• GaN bulk growth is not yet a reality, owing to

the enormous rate of decomposition above

1,000 °C. Other potential solutions, such as

compliant substrates, lateral growth, and pendeoepitaxial overgrowth (that is, epitaxy sideways

over a substrate, which reduces defect density

by five orders of magnitude) are under ONRfunded investigation.

• Acceptor dopants are “deep” in SiC, deeper still

in GaN, and too deep in AlN. Thus, such

components as lasers and bipolar transistors,

which depend on low resistivity p-type films, are

most challenging to produce. Solutions to this

particular challenge may come from recent ONR

initiatives in spontaneous and piezo-polarizationinduced hole conduction.

World experts address these and other related wide

bandgap materials and device challenges in this

edition of Naval Research Reviews.

If only life were that simple for Department of

Defense research and development. Although we

prefer “more affordable” to “cheaper,” comparatives

such as “lighter,” “more powerful,” “more efficient,”

“broader band,” “higher frequency,” “lower noise,”

“stealthier,” and “higher resolution” abound in the

modern Navy’s research, development, and

engineering (RD&E) lexicon. Defense electronics

RD&E does not pass muster unless it fits these, and

many more, criteria.

In truth, we in the DoD should use mostly

superlatives, e.g., most efficient, for that is what we

mean: only the best is good enough for the U.S.

military. ONR’s function is to ensure that the best is

what they get.

More than 80% of command, control,

communications, computers, and intelligence (C4I)

system electronics performance is achieved with

silicon (Si; specifically, complementary metal-oxide

on silicon (CMOS)). Similarly, the highest

performance microwave, millimeter wave, and optoelectronic functions depend upon cubic III-V

semiconductors: GaAs, InP, and their alloys with

AlAs, etc. Development, perfection, and insertion of

these compounds were largely the results of ONR’s

(especially Max N. Yoder’s) forward thinking and

oversight, and of sustained investment by ONR and

DARPA over the last 30 years.

For more than 20 years, Max N. Yoder has been

pushing the wide bandgap boat out into a very

skeptical sea. Although the true potential of wide

bandgap semiconductors such as GaN and SiC was

recognized more than 50 years ago, until recently

the materials problems were thought insurmountable.

Multifunction electromagnetic systems under
Volume 51, Number 1 1

SiC - the compound with many faces...

SiC - The Power Semiconductor for the 21st

Century: A Materials Perspective

by W.J. Choyke and Robert P. Devaty

SiC transit: gloria mundi...

SiC Power Devices

by Anant K. Agarwal, Sita S. Mani, Suresh

Seshadri, Jeffrey B. Casady, Phillip A. Sanger,

Charles D. Brandt, and Nelson Saks

A “MODFET” is not a 1960s-style festival...

GaN-Based Modulation-Doped FETs and

UV Detectors

by Hadis Morkoç

How do they DO that?

Recent Progress in the Growth and Control

of Defects and Electronic Properties in

III-Nitride Thin Films

by Kevin Linthicum, Drew Hanser, Tsvetanka

Zheleva, Ok-Hyun Nam, and Robert F. Davis

2 “I’m a LEO-...”

GaN-Based Solutions for Communications

and Sensing

by Umesh K. Mishra and Stacia Keller

Growing healthy GaN devices in our “Victory

Garden”

GaN-Based Materials Growth for Microwave

Electronic Device Applications

by Daniel D. Koleske, Alma E. Wickenden,

Richard L. Henry, Mark E. Twigg, Steven C.

Binari, Paul B. Klein, Jaime A. Freitas, Jr.,

James C. Culbertson, and Mohammad Fatemi

Profiles in Science

Dr. W.J. Choyke

13

24

44

61

54

72
2 Naval Research Reviews

The

Compound

with Many

Faces

2 Naval Research Reviews
Volume 51, Number 1 3

– S.O.

 n the technological stage, silicon carbide (SiC) has in the last

50 years moved from its previous supporting actor roles in the dramas of

the early 20th Century Industrial Age into the spotlight of the Technological

Age thrillers now being cast for the new millennium. Although this naturally

occurring compound has probably existed since before our solar system

was born, it has truly come into its own only in the last 10 years. Its over

170 polytypes, each with its own unique optical and electronic properties,

make this material a versatile leading actor in the power semiconductor

theater. SiC, with its wide bandgap, has much higher thermal conductivity

and much higher electrical breakdown than Si, which means that it is ideal

for smaller structures that dissipate more power. That’s the charm for

power applications: smaller structures mean faster switching rates. In fact,

while Si devices shut down slowly in a ping-ponging on-off process, SiC

devices do not do that, and are therefore four to five times more efficient.

That yields energy savings that are big box office for high power applications.

Professors W.J. Choyke and Robert P. Devaty of the University of

Pittsburgh explore the material science behind this compound’s road to

stardom. Along the way, they review its history, explain polytypism, and

describe SiC’s band structures, lattice vibrations, and impurity and defect

states. It’s a standing ovation for this comeback star.

O

Volume 51, Number 1 3
4 Naval Research Reviews

W.J. Choyke and Robert P. Devaty

Department of Physics and Astronomy, University of Pittsburgh

Pittsburgh, Pennsylvania

SiC — The Power Semiconductor

for the 21st Century:

A Materials Perspective

Introduction

Isn’t it ironic that silicon carbide (SiC), by far the oldest of

the compound semiconductors, has had to wait until the 21st

Century to realize its great promise? Astrophysicists

discovered during the last 25 years that grains of material

that formed around stars other than our Sun before our solar

system was born 4.6 billion years ago have actually managed to reach our planet. Perhaps the best studied of these

grains is SiC! Most of this SiC is believed to have originated

around stars on the asymptotic giant branch, whereas a

small fraction of it can be traced to ejecta from supernova.

History

The first hint that there could even be a bond between

carbon and silicon was first enunciated by the great Swedish

chemist Jöns Jakob Berzelius in 1824. Sixty-seven years

later in Monongahela, Pennsylvania (near present-day

Pittsburgh), Eugene G. Acheson melted a mass of carbon

and aluminum silicate by passing a high current through a

carbon rod immersed in this mixture. After the mixture had

cooled to room temperature, Acheson found small, bright

blue crystals in the vicinity of the carbon rod. He fully

expected to have made a compound of aluminum and

carbon, and since Al2O3 had been called corundum, he

named his “new” compound carborundum. Acheson had an

excellent chemist on his staff who very quickly determined

that the composition of this “new” compound was in reality

70% silicon and 30% carbon by weight. It was the compound SiC whose existence was first hinted at by Berzelius.

For reasons that defy logic, SiC nevertheless became known

as CarborundumTM the world over.

Acheson was quick to send some of his new SiC crystals to

Professor B.W. Frazier at Lehigh University in Bethlehem,

Pennsylvania. Frazier was a gifted crystallographer, and he

quickly determined by means of optical goniometry that

Acheson’s SiC crystals came in many crystallographic

structures called polytypes. To date, over 170 polytypes of

SiC have been discovered. Since polytypism radically

influences the properties of SiC, we will now spend a little

time analyzing the nature of polytypism in SiC.

What are Polytypes?

Let us consider Fig. 1 where we designate a Si-C atom pair

in an A-plane double layer in close packing as Aa, in the B

plane as Bb, and in the C plane as Cc. It is possible then to

generate a series of structures by varying the stacking

Fig. 1 - Close-packed planes perpendicular to the c

axis (principal axis) in hexagonal or rhombohedral

lattices or [111] in diamond or zincblende

lattices. Note that the z axis, mentioned in the

text, is perpendicular to the x-y plane.
Volume 51, Number 1 5

sequence along the principal crystal axis (the z axis,

perpendicular to the x and y axes in Fig.1). For AaBbCc

AaBbCc...stacking, we generate the zincblende (cubic)

form of SiC known as 3C-SiC. On the other hand, if we

stack the atom pairs as AaBbAaBb...we generate the

wurtzite (hexagonal) form of SiC known as 2H-SiC. Other

common forms of SiC may be generated by the following

stacking sequences along the principal axis: 4H-SiC

(AaBbAaCcAaBbAaCc...), 6H-SiC(AaBbCcAaCcBb

AaBbCcAaCcBb...) and 15R-SiC(AaBbCcAaCcBbCc

AaBbAaCcAaBbCcBbAaBbCcAaCcBbCcAaBbAaCcAa

BbCcBb...). The notation is clear: C for cubic, H for

hexagonal, and R for rhombohedral. For simplicity, when

dealing with SiC, one normally reduces the notation Aa, Bb,

Cc to A, B, C. It is interesting to note that the largest unit

cell of a SiC polytype that has thus far been identified has an

astounding length of about 1200 Å along the principal axis

(the z axis, also commonly referred to as the c axis). The

reason for the stability of so many SiC polytypes is still a

question of great interest and current study.

The long slender unit cells of SiC are not just a scientific

curiosity but are in fact the reason that we should think of

Fig. 2 - This schematic illustrates inequivalent sites in 6H-SiC. N1, N2, and

N3 are nitrogen atoms substituting on three inequivalent carbon sites in

the SiC lattice. N1 is substituting on a hexagonal site, h, and N2 and N3

are substituting on two quasicubic sites, k1 and k2. The inequivalence of

the sites is illustrated in the table by the distance of each of the nitrogen

atoms to carbon-like and silicon-like planes. Note that the length of the

unit cell is only 15.1 Å as indicated on the left of the figure.

SiC as a large family of semiconductors. Each polytype of

SiC has some unique optical and electronic properties. The

designer of a particular SiC device should consider very

carefully in which polytype optimum characteristics may be

obtained. One reason for this diversity is that the number of

inequivalent Si or C sublattice sites varies among the

polytypes. For example, there is one in 3C and 2H, two in

4H, three in 6H, and five in 15R. This means that if we

substitute nitrogen (N), which acts as a donor in SiC, or

aluminum (Al), which acts as an acceptor in SiC, in the 4H

lattice, we obtain not just one donor/acceptor but rather two

donors/acceptors. In 6H-SiC, we would have three donors/

acceptors, and in 15R SiC, five donors/acceptors. Why is

this so? Figure 2 is a schematic representation of the atomic

arrangements of the Si and C atoms in the (1120) plane of

the 6H-SiC hexagonal pyramid. In general, when representing the atomic stacking of any SiC polytype in the (1120)

plane, the complicated stacking sequences can be expressed

in terms of simple zigzags. As shown in Fig. 2, 6H-SiC can

be illustrated by a zig of three lattice positions (ABCA) to

the right and a zag of three lattice positions to the left

(ACBA). In short hand, we write (22) for 4H, (33) for 6H,

and (232323) for 15R, and so on.
6 Naval Research Reviews

Looking more carefully at Fig. 2, we designate a carbon

atom on a particular plane perpendicular to the c axis with a

small dot and the associated silicon atom with a larger gold

circle. A point h designates a lattice position in which the

carbon or associated silicon atom finds itself in a

“quasihexagonal” stacking environment with respect to its

neighboring stacking planes. Similarly, we see that k1 and k2

represent lattice points in which the carbon and associated

silicon atoms find themselves in “quasicubic” stacking

environments. In addition, as the inset shows, if we substitute a nitrogen atom at the h, k1, and k2 carbon sites, the

atomic distances for the substitutional N1 nitrogen to the

nearest carbon-like or silicon-like planes distinctly differ

from those for the substitutional N2 and N3 nitrogen atoms.

The N2 and N3 nitrogen atoms sense different environments!

This shows why the hexagonal site h and the two cubic sites

k1 and k2 in 6H-SiC are expected to have, and do have,

slightly different electronic properties.

Figure 3 shows the stacking sequences and selected physical

properties of the five polytypes of major interest to current

research and development. The lattice constant in the

zincblende 3C-SiC modification and the lengths of both the

pyramid face and the principal axes of the unit cells in the

case of the hexagonal lattices for 15R, 6H, 4H, and 2H-SiC

are all obtained from X-ray measurements. Once the unit

cells for the various polytypes are known, the number of

atoms per unit cell is determined as well as the number of

inequivalent sites. The space groups for the various

polytypes are a consequence of the atomic arrangements of

the Si and C atoms in each polytype. Finally, the indirect

exciton energy gap at 2K is obtained from optical measurements. The latter leads us directly, in the following section,

into a consideration of the band structure of SiC polytypes.

Band Structure

During this decade, much progress has been made by

theorists in elucidating the electronic and vibrational

structure of the polytypes of SiC, some of whose physical

properties are shown in Fig. 3. At the same time, improved

crystal quality and purity have made experiments possible to

verify the theory and point to new directions that require a

great deal more study. It has been known since the late

1950s that SiC is an indirect semiconductor. To illustrate

what we mean, refer to Fig. 4. This is a schematic represenFig. 3 - A summary of some of the physical properties of 3C, 15R, 6H, 4H, and 2H-SiC. If these

polytypes are all represented in the (1120) plane of a hexagonal pyramid, then the repeat distance

along the c axis would be cn.
Volume 51, Number 1 7

tation of indirect exciton recombination for a semiconductor

with the valence band (VB) maximum at k = 0 and conduction band (CB) minima at k = kCB. In a direct semiconductor

such as GaAs or GaN, the extrema of the valence and

conduction bands are located in the same place in k space at

the center of the Brillouin zone, the Γ point, and the

absorption or emission of photons can take place without the

assistance of lattice phonons. We term such transitions

direct. In SiC, Si, and Ge, the CB minima are not at the

same place in k space as the VB maximum and we require

the extra help of phonons to permit transitions from the VB

to the CB. Such transitions are termed indirect. As Fig. 4

shows, there are various possibilities for making these

transitions. One could either scatter an electron with a

phonon from the CB minimum to an intermediate state at k0

(k = 0) in the CB and then make a direct transition to the VB

maximum or one can scatter a hole via a phonon to an

intermediate state in the VB at kCB and then make a direct

transition to the CB minimum.

Most experimental information about the band structure of

SiC has come from low-temperature optical experiments.

Near zero degrees Kelvin, if one absorbs a photon with

energy larger than the bandgap and raises an electron from

the VB to the CB, leaving behind a hole in the VB, one has

created an excited state of the crystal. After a very short

time, the free electron in the CB as well as the free hole in

the VB not only scatter to the minimum of the CB and

maximum of the VB, respectively, but lower their energy

even further by having a slight coulomb attraction for each

other. We call such an electron-hole pair a free exciton. In

low-temperature experiments, in which we attempt to

determine the bandgaps of the polytypes, we really measure

the energy that it takes to overcome the bandgap minus the

energy that is given back due to the slight coulomb binding

of the electron-hole pair (exciton). Hence, we measure the

exciton bandgap (EGX) values that were given in Fig. 3. We

now need to explain the existence of the “no-phonon”

transition indicated in Fig. 4. If we have impurities or

defects that have energy levels in the bandgap, it is very

probable at low temperatures that the exciton is captured

and bound to these centers. This localization permits a small

number of transitions to the VB to proceed without the need

of phonon participation. In the case of low-temperature

photoluminescence (LTPL), we indeed see very sharp nophonon lines and they are immensely helpful in the analysis

of the data.

A comprehensive treatment of the fundamental theoretical

aspects of SiC has recently been given in the first six articles

in “Silicon Carbide:...Vols. I and II” [1]. One is struck by

the progress that has been made since the start of the current

renaissance of SiC in the late 1980s. Of particular interest

for device development is the guide that theory has given in

locating the positions of the conduction band minima in the

zincblende Brillouin zone of 3C-SiC, in the hexagonal

Brillouin zones of 2H, 4H, and 6H-SiC, and in the rhombohedral Brillouin zone of 15R-SiC. Theory has also given

valuable guidance as to the structure of the VB near its

maximum at k = 0 in several of the common polytypes.

Finally, theory has made great strides in calculating the

tensor components of the electron-effective masses in the

common polytypes of SiC. The hole effective masses have

also been estimated but there is unfortunately no experimental confirmation available to date.

Figure 5 summarizes the best current information from

experiment and theory concerning the CB minima as well as

the VB maxima for 2H, 4H, 15R, 6H, and 3C-SiC. Numerical values obtained purely from theory are marked as such.

We have arranged the polytypes with the 2H wurtzite

structure at the top and the zincblende 3C structure at the

bottom. This ordering in terms of “hexagonality” shows a

number of interesting trends. As we go from 2H to 3C-SiC,

the forbidden gap becomes progressively smaller.

Lambrecht et al. [2] have shown how this comes to pass

from their band theoretical considerations. At the top of the

VB, there is a threefold degeneracy (neglecting spin) and we

expect the spin-orbit interaction (∆S-O) in the cubic,

Fig. 4 - An illustration of possible electronic transitions

from the maximum of the valence band to the minimum

of the conduction band.
8 Naval Research Reviews

zincblende, lattice to split this degeneracy into fourfold and

twofold degenerate bands at k = 0. For the hexagonal and

rhombohedral polytypes, we have in addition a crystal field

that results in an additional splitting (∆C-F) such that we

have three nondegenerate bands at the VB maximum at k =

0. We see from Fig. 5 that in the hexagonal and rhombohedral polytypes, we measure a spin-orbit splitting of about

seven milli-electron-volts (meV), whereas in the cubic

modification 3C-SiC, we find a spin-orbit splitting of 10

meV [3, 4]. Theory [2] predicts a linear crystal-field

splitting from zero to 144 meV as we go from cubic SiC

(3C-SiC) to 100% hexagonal SiC (2H-SiC). Until now, we

have only been able to measure the crystal-field splitting in

6H-SiC [5], and the experimental value of 44 meV is in

good agreement with the value expected from theory.

Band calculations are used by several authors, in slightly

different approaches, to obtain the effective mass tensor

components for the bottoms of the conduction bands in the

polytypes under discussion. A knowledge of the electron

effective masses is extremely important for device modeling. Various optical techniques as well as cyclotron resonance have been used to obtain electron effective mass

values in 3C-SiC. Here the various experimental techniques

are in excellent agreement. Except for 4H-SiC, for which

optical detection of cyclotron resonance has given reliable

values of the electron effective mass tensor components, we

have run up against serious interpretational problems in

absorption experiments on shallow donors, mainly due to a

lack of a proper and sufficiently general effective mass

theory. Figure 6 is an adaptation from a table in the paper of

Wellenhofer and Rössler [6] where all the current experimental and theoretical values for the effective mass components at the respective CB minima for 3C, 2H, 4H, 6H, and

15R-SiC are given. For clarity we have left out the references to the various calculations and experiments, but all the

references may be found in Wellenhofer and Rössler [6]. As

one of our next important tasks, it will be most useful to get

accurate and reliable measurements of all the band parameters in the pure wurtzite 2H form of SiC. The simplicity of

the 2H modification will facilitate further intrepretation in

the more complicated hexagonal and rhombohedral

polytypes. Unfortunately, this will require much larger, less

defective, and much purer single crystals of the 2H polytype, presenting a super challenge to the crystal growers!

Lattice Vibrations

We have already mentioned that in LTPL, many sharp lines

are observed, most of which result from transitions involving phonons (lattice vibrations). These experiments are a

rich source of information on phonons, especially in indirect

semiconductors such as SiC. However, the phonons that we

see in these spectra are characteristic of very special

positions in k space. In spectra due to free excitons or

excitons bound to shallow donors, we sense the phonons

characteristic of the k space positions of the CB minima in

2H, 4H, 6H, 15R, and 3C-SiC.

Other common techniques for the study of vibrational

properties in semiconductors are Raman scattering and

neutron scattering. Until very recently, no single crystals of

SiC of sufficient size have been available to make neutron

scattering practical. In 1968, Feldman, Parker, Choyke, and

Patrick [7, 8] introduced a new method of using Raman

scattering to obtain phonon dispersion curves based on the

existence of polytypes. Measurements on a polytype often

yield results that are determined partly by the properties

Fig. 5 - A compilation of salient features of the band structures of 2H, 4H, 15R, 6H, and 3C-SiC.

P y olytype Hexagonalit

Exciton

Bandgap

EGX (2K)

Position of Conduction Band

Minimum in the Brillouin Zone

# of CB

minima

∆S-O

(meV)

∆C-F

(meV)

21 0 H 3 ) .33 K 2— (HEXAGONAL 144

(Theory)

4 5 H 0 5 . 3 ) .26 M 3— (HEXAGONAL 72

(Theory)

1 4 5R 0 6 . 2 ) .98 X 37 (RHOMBOHEDRAL 57

(Theory)

6 3 H 0 3 .3 3.02 U (L - M)

HEXAGONAL 67 44

30 0 C 2 ) .39 X 30 (ZINCBLENDE 1 0
Volume 51, Number 1 9

common to all polytypes and partly by the unique structure

of the polytype on which the measurements are made. An

example is the first-order Raman spectrum of 6H-SiC [7].

The property assumed to be common to all polytypes is, in

this case, the phonon spectrum in the axial direction [8].

Most of the lines in the Raman spectrum are, nevertheless,

characteristic of the 6H-SiC structure, which determines the

points of the common dispersion relation spectrum accessible to Raman measurements. The relationship of the

Raman spectrum to the phonon spectrum is clear from a

large zone [9] point of view, which permits the assignment

of Raman frequencies to positions on dispersion curves,

even though these modes have zero wave vector (approximately) in the Brillouin zone. The procedure is now termed

“zone folding” and is used extensively in the analysis of

data obtained from semiconductor superlattices grown by

means of molecular beam epitaxy (MBE) or metal-organic

chemical vapor deposition (MOCVD) techniques. Figure 7

illustrates the existence of a common SiC phonon spectrum

(within ~2%) for polytypes 4H, 6H, 15R, and 21R-SiC.

Nakashima et al. [10] have added data for 8H and 27R-SiC,

which fit extremely well on the curves shown in Fig. 7. This

method has given us phonon dispersion curves comparable

in completeness and accuracy with those obtained for other

materials by neutron diffraction (but only for one direction

in momentum space). Nakashima and Harima [11] have

recently given a detailed review of the current applications

of Raman scattering to the study of the vibrational properties of the polytypes of SiC.

Impurity and Defect States

In a perfectly ideal world, we would first produce ultrapure

single crystals of a material before attempting any doping

with impurities. Fifty years ago we already knew that certain

impurities such as nitrogen, aluminum, and boron could be

readily incorporated into SiC. We also knew that nitrogen

made SiC n-type, and aluminum and boron made SiC ptype. What we didn’t know was how to produce a relatively

clean and defect-free single crystal. To make matters worse,

the best crystals were small platelets with mixed polytypes.

The cry then was, if we could only produce large boules of a

single polytype with a high degree of doping control, we

would have a marvelous industrial process. These things

have now come to pass! Two-inch wafers are rapidly

becoming available with a low enough micropipe density to

serve most applications. The n- and p-type doping ranges of

these boules can be controlled from roughly 5 × 1015 cm–3

Fig. 6 - Theoretical and experimental values for the effective electron masses in 3C, 2H, 4H, 6H, and 15R

SiC. The different columns compare different theoretical as well as experimental values obtained in a

variety of ways.

m 0.70 0.67 0.63 0.68 0.60 0.667 0.67 0.67 m⊥ 0.23 0.25 0.23 0.23 0.29 0.247 0.25 0.22

m⊥ 0.40 0.45 0.45 0.43 m 0.26 0.26 0.27 0.26

mM-Γ 0.66 0.57 0.58 0.57 0.58 m⊥ : 0.18 0.30 mM-K 0.31 0.32 0.28 0.28 0.31

mM-L 0.30 0.32 0.31 0.31 0.33 m : 0.22 0.48

mM-Γ 0.78 0.75 0.77 0.75 m⊥ : 0.35 0.24 0.25 0.42

mM-K 0.23 0.27 0.24 0.24

mM-L 1.2-2.0 1.95 1.24 1.83 m : 1.4 0.34 1.7 2.0

mX-Γ 0.67 m⊥ : 0.28 0.24

mX-K 0.22 mX-U 0.41 m : 0.53 0.38

3C SiC minimum at X

2H SiC minimum at K

4H SiC minimum at M

6H SiC minimum along M-L

15R SiC minimum at X

m(m0) THEORY EXPERIMENT

kz

kx

ky

Γ K X

U

kz

kx

ky

Γ

A

H

M K

L

kz

X

Γ

L

U

WK

ky

kx
10 Naval Research Reviews

to 5 × 1019cm–3. Device structures are presently made on

homo-epitaxial films grown on boule wafers. Currently used

epitaxial films range from a thickness of about 0.5 μm to

100 μm and a range of possible n- and p- type doping of less

than 1014cm–3 to greater than 1019cm–3.

The conductivity type of SiC, as we have already learned,

can be adjusted by doping with either donors such as

nitrogen and phosphorus or acceptors such as aluminum,

boron, or gallium. Most doping must be done during boule

or epitaxial growth since the diffusion coefficients for these

dopants are very small at workable temperatures. An

alternative is ion implantation, which has been used extensively in SiC, although it is far more difficult to use than in a

single atom material such as Si. Even at annealing temperatures of 1700 °C, residual defect centers from the implantation may still be present. Understanding and controlling

such centers are key to using ion implantation in SiC.

Nitrogen is the best studied donor in SiC. It is known to

substitute on the carbon sublattice and some have claimed to

achieve doping densities as high as 3 × 1020 cm–3. A very

high doping density is required for vertical power device

structures to obtain low power losses. Resistivity values of

about 2 × 10–3 Ω-cm have been reached in SiC boule wafers,

which are used as the substrates for these devices. A

resistivity of 2 × 10–3 Ω-cm compares favorably with values

that have been reported for Si. For high frequency devices,

on the other hand, semi-insulating substrates are desirable.

By compensating boules during growth with elements such

as vanadium, resistivities of about 109 Ω-cm at 500 K have

been reported. Another approach that is appealing and is

currently being actively pursued is producing bulk material

that is so pure that the intrinsic resistivity is high enough to

satisfy the demands of the high frequency designs and at the

same time preserve the good transport characteristics of the

material. In epitaxial growth, the residual nitrogen concentration is normally the limiting feature in making films with

carrier concentrations below 1014 cm–3. Very high voltage

applications (i.e., greater than 25 kV) are now driving efforts

to produce thick single-crystal films (100 μm) with ultra

high purity (less than 5 × 1013 dopants per cm3).

In the last few years, phosphorus donors have been incorporated into 6H-SiC during epitaxial growth by ion implantation and through neutron transmutation. Only in the CVDgrown material have we seen a bound exciton spectrum of

phosphorus. It gives the signature of a slightly shallower

donor than nitrogen. Ionization energies obtained from Hall

measurements on implanted and annealed material bear this

out. If we invoke site competition epitaxy (SCE) in analyzing the CVD growth, we conclude that the phosphorus

substitutes on the Si sublattice. Atomic size arguments and

evidence from electron spin resonance (ESR) measurements

also point in this direction.

Gallium, aluminum, and boron serve as “shallow” acceptors

in SiC and are most frequently introduced during the growth

process. However, implantation can be used to generate ptype material by overcompensating the nitrogen donor

background. Gallium and aluminum substitute on the Si

sublattice but the situation for boron appears more complex.

Boron is a maverick by forming a shallow acceptor as well

as a deep state that acts like an acceptor from the point of

view of donor-acceptor recombination spectra. Current

interpretations of ESR and electron nuclear double resonance (ENDOR) experiments point to the following: The

shallow boron state is due to a boron atom at an offcenter

position near a Si sublattice site. The deep center is due to a

boron atom at an offcenter position near a Si sublattice site

adjacent to a vacancy on the carbon sublattice site. Both the

deep and shallow boron centers are aligned along the

principal crystal axes. In contrast to the situation with

donors, the ionization energies of the acceptors are only

weakly dependent on polytype or on the hexagonal (h) and

Fig. 7 - A common SiC phonon spectrum (within

~2%) derived from first order Raman scattering data

in the polytypes 4H, 6H, 15R, and 21R-SiC.
Volume 51, Number 1 11

quasicubic (k) lattice sites. Hall measurements yield an

average value for the ionization energy of aluminum of

about 200 meV, and for the shallow boron acceptor, they

yield an average value of the ionization energy of about 300

meV. In CVD growth, the concentrations of shallow and

deep boron centers can be adjusted by use of SCE.

Figure 8 is an adaptation of a figure from the recent Ph.D.

dissertation of Thomas Troffer at the University of

Erlangen-Nürnberg. Here we summarize what is known

about the energy levels of the main impurities in SiC. The

band offsets between polytypes are believed to occur mainly

in the conduction bands. For the VB offset between 6H- and

4H-SiC, we take a value of 50 meV, leaving 192 meV for

the 6H- to 4H-SiC CB offset. The large CB offset accounts

for the fact that in 4H-SiC, one observes two acceptor levels

for Ti3+/4+, whereas in 6H-SiC, these levels are resonant with

the CB and, hence, are not seen. For the sake of clarity, the

energy positions of the donor and acceptor levels in the

forbidden gap have not been drawn to scale. Furthermore,

the quoted binding energies are rough values and are subject

to much fine-tuning!

We have already alluded to the fact that during ion implantation, defect states are introduced into the SiC lattice. Some

of these states have been extensively studied. The so-called

DI

 and DII centers [12] appear to be the most persistent and

can still be detected after a 1700 °C anneal. Their effect on

electrical properties of SiC is still not certain and is currently under study. High energy electron bombardments

(i.e., higher than 1 MeV) have led to annealing studies of

lattice defects from room temperature up to 1700 °C. The

main conclusions seem to be that up to 800 °C, single

vacancies can exist in SiC, whereas above this temperature,

only vacancy complexes are stable. During boule and

epitaxial growth, some of the same defect centers seen in

irradiated materials may also be generated. Normally these

centers are seen, primarily by means of LTPL, after too

rapid growth or too rapid cooling after growth. In recent

years, many deep levels have also been observed with LTPL

and deep level transient spectroscopy (DLTS) measurements

but have not as yet been identified. The study of defects in

SiC is just in its infancy and will be vigorously pursued

because of its clear importance to device performance.

Conclusion

We have seen that SiC has emerged in the last 50 years from

being a ceramic material used as a universal abrasive,

heating element, and reducing agent in metallurgical

processes to a semiconductor of great promise for high

power, high frequency devices and unique applications in

radiation and corrosive environments. To bring all this about

will require a continued, aggressive, and coordinated effort

Fig. 8 - Impurity centers in SiC polytypes 6H

and 4H. For the sake of clarity, the positions

of levels for the various donors and

acceptors are not drawn to scale. The valence

band offset is taken to be 50 meV, and the

conduction band offset is 192 meV.
12 Naval Research Reviews

of materials growth, materials characterization, and fundamental understanding.

Acknowledgment

We wish to thank Dr. S.G. Sridhara for his great help in

preparing the figures.

References

1. Silicon Carbide: A Review of Fundamental Questions and

Applications to Current Device Technology, Vols. I and II,

W.J. Choyke, H. Matsunami, and G. Pensl, eds. (Akad. Verl.,

Berlin, 1997), ISBN 3-05-501792-7.

2. W.R.L. Lambrecht, S. Limpijumnong, S.N. Rashkeev, and B.

Segall, “Electronic Band Structure of SiC Polytypes: A

Discussion of Theory and Experiment,” Phys. Stat. Sol. (B)

202, 5-33 (1997).

3. R.G. Humphreys, D. Bimberg, and W.J. Choyke, “Wavelength-Modulated Absorption in SiC,” Sol. State Commun. 39,

163-167 (1981).

4. R.P. Devaty, W.J. Choyke, S.G. Sridhara, L.L. Clemen, D.G.

Nizhner, D.J. Larkin, T. Troffer, G. Pensl, T. Kimoto, and H.S.

Kong, “Optical Properties of SiC: Some Recent Developments,” Materials Science Forum 264-268, 455-460 (1998).

5. W.J. Choyke, R.P. Devaty, and S.G. Sridhara, “A Survey of

Conduction and Valence Band Edges in SiC,” accepted for

publication in Physica Scripta (1998).

6. G. Wellenhofer and U. Rössler, “Global Band Structure and

Near-Band-Edge States,” Phys. Stat. Sol. (B) 202, 107-123

(1997).

7. D.W. Feldman, J.H. Parker, W.J. Choyke, and L. Patrick,

“Raman Scattering in 6H SiC,” Phys. Rev. 170, 698-704

(1968).

8. D.W. Feldman, J.H. Parker, W.J. Choyke, and L. Patrick,

“Phonon Dispersion Curves by Raman Scattering in SiC,

Polytypes 3C, 4H, 6H, 15R and 21R,” Phys. Rev. 173, 787-

793 (1968).

9. H. Jones, The Theory of Brillouin Zones and Electronic States

in Crystals (North Holland Publishing Company, Amsterdam,

1960).

10. S. Nakashima, Y. Katahama, Y. Nakakura, and A. Mitsuishi,

“Relative Raman Intensities of the Folded Modes in SiC

Polytypes,” Phys. Rev. B 33, 5721 (1986).

11. S. Nakashima and H. Harima, “Raman Investigation of SiC

Polytypes,” Phys. Stat. Sol. (A) 162, 39-64 (1997).

12. S.G. Sridhara, D.G. Nizhner, R.P. Devaty, W.J. Choyke, G.

Pensl, and T. Kimoto, “DII Revisited in a Modern Guise —

6H and 4H SiC,” Materials Science Forum 264-268, 493-496

(1998).

The Authors

For the biography of W.J. Choyke, see Profiles in

Science on p. 72 and the inside back cover.

After earning a Ph.D. in Physics from Cornell

University, Robert P. Devaty was a National Research

Council (NRC) Associate at the Naval Research

Laboratory. He joined the Department of Physics and

Astronomy at the University of Pittsburgh in 1984. In the

1980s, he applied far-infrared and infrared techniques to

the study of small particles and metal-insulator composite

materials under ONR support. After he was promoted to

Associate Professor, he changed the focus of his research

to optical studies ranging from the ultraviolet to the farinfrared of large bandgap semiconductors, particularly

silicon carbide. Together with W.J. Choyke, he participates

in a number of international collaborations, including

Gerhard Pensl’s group at the University of ErlangenNürnberg and Peter Déak’s theoretical group at the

Technical University of Budapest. He is co-editor of the

upcoming International Conference on Silicon Carbide

and Related Materials (ICSCRM’99).
– S.O.

P

 erhaps it is an overstatement to say that the way silicon carbide (SiC)

semiconductors pass electric current is one of this world’s glories, but the

compound’s increasing role in advancing high power applications may make a

little exaggeration excusable. In the following article, Anant Agarwal and other

industry and government researchers illuminate the materials science behind

this new-old class of semiconductors, especially its 4H-SiC and 6H-SiC

polytypes. They also describe the rapid progress in the development of p-n

diodes and prescribe steps for further development of SiC power devices.

The concurrent surge in the development of SiC-based power metal-oxide

semiconductor field effect transistors (MOSFETs) is

also summarized. They explain problems such as stepbunching and explore the SiC polytypes’ strengths

and limitations. Finally, the authors introduce a new

power switching device configuration, the JFET Controlled Thyristor, and detail its usability for high power

applications, such as those found in power plants and

scores of Defense-oriented projects. Of particular

interest to Defense S&T is the usability of SiC-based

devices in the entire range of all-electric tanks, planes,

and ships currently being designed.

 Silicon carbide was used in this moissanite gemstone. Photograph courtesy of Cree Research, Inc.

Volume 51, Number 1 13
14 Naval Research Reviews

SiC Power Devices

Anant K. Agarwal,* Sita S. Mani, Suresh Seshadri,

Jeffrey B. Casady, Phillip A. Sanger, and Charles D. Brandt

Northrop Grumman Science & Technology Center

Pittsburgh, Pennsylvania

Nelson Saks

Naval Research Laboratory

Washington, D.C.

Introduction

Today’s power electronic components are predominantly

made of silicon (Si). Silicon technology has matured over

the last 40 years, however, so that no major increase in

power handling capability of these components is expected

in the future. This limit can be overcome by the superior

electronic properties of silicon carbide, which offers at least

a tenfold increase in power handling capability. High-power,

high-speed 4H-SiC diodes and switches operating at high

temperatures (>350 °C) are expected to play enabling and

vital roles in the design of the future concept military

hardware. The system level benefits include a large reduction in the size, weight, and cost of the power conditioning

system. To ensure reliability and longevity of military

electronics, the junction temperature of conventional siliconbased devices is currently limited to the MIL-STD temperature range of -55 °C to 125 °C. High-temperature devices

such as the 4H-SiC rectifiers proposed here can be mounted

on high-temperature stations such as motors, generators, and

energy management systems to improve mobility, survivability, lethality, and fuel economy. These devices will help

in the realization of solid state power conditioning and

circuit protection systems used for motors, actuators, energy

storage, and pulse power systems. The high power p-i-n

rectifier proposed here will be critical to the development of

DC zonal electric power distribution being developed under

Power Electronic Building Blocks (PEBBs). These rectifiers

may lead to greatly reduced noise and thermal signatures of

ships, submarines, tanks, and armored personnel carriers.

SiC Material Properties

This article describes our approach to very-high-power SiC

devices for low-frequency switching applications. Reference

1 presents an overall review of SiC electronics. Although

more than 170 polytypes of SiC are known, only two (4HSiC and 6H-SiC) are available commercially. 4H-SiC is

preferred over 6H-SiC for most electronics applications

because it has a higher and more isotropic electron mobility

than 6H-SiC. Table 1 compares some key electronic

properties of 4H-SiC to Si and GaAs.

Table 1 - Key Electronic Properties of

P n roperty S s ilico G C aA 4H-Si

E 2 g (eV) 1 5 .1 1 5 . 3.2

μn (cm2

/Vs) 1 0 400 9 0 20 80

μp (cm2

/Vs) 4 0 50 4 0 0 14

ni (cm-3

) 5 at 300 K 1. × 1010 2.1 × 106 5 × 10-9

vnsat (× 107 c 0 m/s) 1 0 . 1 0 . 2.

Ecrit ( 5 MV/cm) 0 4 .2 0 0 . 3.

ΘK (W/cmK) 1 5 .5 0 9 . 4.

Si, GaAs, and 4H-SiC

The most important property of SiC from the power device

perspective is its indirect bandgap, which results in relatively high recombination lifetime (~1 μs) provided the

material is pure enough. This allows the development of

high-voltage p-n junction diodes and p-n-p-n thyristors that

cannot be produced from a direct bandgap material such as

GaN. The higher bandgap of 4H-SiC results in a very high

intrinsic temperature of 1650 °C for an extrinsic doping of

1016 cm-3. Another advantage of high bandgap is negligible

junction leakage currents up to 500 °C (μni). This allows

high-temperature operation without excessive leakage

current or thermal runaway. The high junction temperature

increases the heat flux that can be removed from the device,

* Currently with Cree Research, Inc., Durham, NC
Volume 51, Number 1 15

resulting in reduced cooling requirements. The high

breakdown strength, Ecrit, results in drift layers that are 10

times smaller than those of Si for a given blocking voltage,

thus reducing both the storage of minority carriers (i.e., Qrr

in diodes is reduced) and associated switching losses at a

given switching frequency. This means that the switching

frequency can be increased to between 50 and 100 kHz with

acceptable switching losses, which, then, would significantly reduce the size and weight of the magnetic components in an inverter system.

Finally, the thermal impedance of

the packages is reduced by the

high thermal conductivity of SiC,

which helps in spreading the heat

laterally. It is expected that the

size and weight of power electronics using SiC switching

devices and diodes will be significantly reduced by means

of passive cooling, use of smaller and lighter magnetic

components, and reduced snubbersize.

Figure 1 shows a three-phase 36-kV transfer switch composed of 10 4.5-kV silicon thyristors in series in each of the

three phases. This transfer switch is manufactured by

Silicon Power Corporation (SPCO) of Melvern, Pennsylvania and is used by big industrial plants to redirect power

from one grid to the other during power failures and other

such events. Each thyristor is mounted on a very big heat

sink and comes with its own gate-drive circuitry. The high

weight, large volume, high cost, and transportation of this

unit all present big problems. In a future implementation of

SiC technology, each bank of thyristors could be replaced

by a single SiC thyristor of 40 kV rating operating at 300 °C

on a smaller heat sink, thus reducing the weight, volume,

and cost of the system while improving the reliability.

Fig. 1 - A 36-kV three-phase transfer switch with 10

silicon thyristors in series in each of the three banks

(photo courtesy of Silicon Power Corporation).

SNUBBER. Passive circuit made of resistors

and capacitors. They are attached to

switches to safeguard them from extreme

voltage or current stress during switching

events.

p-n Diodes

Virtually all power converter applications require a switch

(such as a MOSFET, insulated gate bipolar transistor

(IGBT), thyristor, or MOS controlled thyristor (MCT)) and

a diode in shunt with it. The transistor is switched on and

off at a high frequency (2 to 10 kHz). Increasing the

switching frequency can reduce the size, volume, weight,

and cost of transformers and inductors in the system but will

also increase the switching losses.

Figure 2 presents an example of the

switching loss in which the current

voltage waveform of a diode is shown.

As the diode is turned off, a negative

spike in current results in high

switching losses because the voltage

across the diode rises to a high value. The charge released

during the negative current spike is known as the reverse

recovery charge, Qrr. This represents a big loss component

in present-day silicon diodes above 2000 V. Due to the

material properties of SiC (particularly its high breakdown

field strength), the reverse recovery charge can be reduced

by about 90% in SiC diodes.

Fig. 2 - Silicon diodes above 2000 V have

a large reverse recovery charge, which

increases the switching losses. SiC

diodes will virtually eliminate it.

Figure 3 illustrates SiC p-n diode construction. The diodes

are fabricated on n+ 4H-SiC wafers with 8 to 10 μm of ntype epitaxial layers doped at about 5 × 1015 cm-3. Unlike in

silicon, foreign atoms do not diffuse in SiC below 2000 °C.

Thus, impurity atoms are implanted using high-energy (400

keV) acceleration equipment called ion-implanters. The

anode and guard rings are formed by ion-implantation of Al

at 1000 °C to the depth of about 0.5 μm with a peak

chemical concentration of about 1 × 1020 cm-3 using a

boxlike profile. The implants are electrically activated at

1700 °C for 10 min in Ar ambient. A sintered titanium layer

forms the contact to the p+ region, whereas sintered nickel is

used to contact the backside.

time

Current

Si

SiC

Qrr

3-Phase Transfer Switch - courtesy of SPCO

Silicon

• thermally limited

• big and heavy

SiC

• passive cooling

• smaller and lighter
16 Naval Research Reviews

Fig. 3 - A schematic cross section of an SiC implanted diode

structure.

The reverse breakdown voltages on diodes of different sizes

(100-μm to 1-μm diameter) were measured. The average

value of the breakdown voltage was found to be dependent

on the size of the diode. For example, the 600-μm diameter

diodes broke down at voltages that were typically 20% to

30% lower than the average values for 100-μm diodes. This

perimeter- or area-related effect indicates a high dislocation

density (low 104 cm-2) in the material. The reverse leakage

current vs reverse voltage plot (Fig. 4(a)) shows a low

leakage current up to about 600 V and then a significant

increase up to the breakdown voltage of about 880 V. This

also indicates a trap or dislocation-related tunneling process

under high electric fields.

The room temperature forward I-V curve for the same diode

(Fig. 4(b)) shows high forward drop of 6.5 V at 10 A. The

forward drop consists of drops across anode contact

resistance, drop across probe and lead resistance, drop in the

drift layer, and a drop in the substrate including the backside

contact resistance. The substrates are typically 300 μm

thick with resistivity of about 0.03 ohm-cm (at room

temperature), representing a drop of about 0.45 V at a

current density of 500 A/cm2. A large fraction of the

forward drop arises from the poor contact resistance (~ 10-3

ohm-cm2) of sintered Ti to the implanted p+ regions. This

issue is being addressed through coimplantation of C and

Al, which has a potential for better electrical activation.

A Four-Terminal Switch

In addition to a diode, most of the power electronic applications require a switch. Many switches, such as the

MOSFET, IGBT, GTO, MTO, and MCT, have been developed using silicon technology. Recent development of a

hybrid MOSFET Turn-Off Thyristor (MTO) has been very

successful in very high power applications requiring

switching of more than both 4000 V and 100 A. In the

center of a silicon MTO (Fig. 5) is a silicon Gate Turn-Off

1E-10

1E-9

1E-8

1E-7

1E-6

1E-5

1E-4

1E-3

-1200 -1000 -800 -600 -400 -200 0

Current (A)

On-wafer DC measurement

600-μm dia. Al-implanted diode

TRPDIODE1

BO906-7

C1R4 Dev4

Reverse Voltage (V)

0

2

4

6

8

10

12

14

0 2 4 6 8

Current (A)

Current (I)

probe R

removed

probe & subs.

R removed

10 A

BO906-7

600-μm dia. pkg.

3 of 5 devices

Forward Voltage (V)

(a)

(b)

Fig. 4 - (a) Reverse and (b) forward I-V characteristics of

a 600-μm diameter p-n diode.
Volume 51, Number 1 17

Thyristor (GTO) capable of switching up to 4.5 kV and

several hundred amps. Ten MOSFET chips are attached to

the periphery of the GTO between the gate and the cathode

terminals of the GTO. The MOSFETs are used to help turn

the GTO off after it has been turned on. This approach

works well for silicon devices but needs to be slightly

modified for SiC because the silicon and SiC MOSFETs are

not expected to be reliable at temperatures above

200 °C.

Our approach to SiC power switching devices is based on a

new device called the hybrid JFET Controlled Thyristor

(JCT) (patent pending) shown in Fig. 6.

The advantages of the hybrid approach are

1. The high-power GTO and the low-power JFET can be

independently optimized and separately processed.

2. The hybrid assembly may be operated at junction

temperatures as high as 500 °C, resulting in reduced

cooling requirements and reduced weight and volume.

In the case of JCT, the four-layer, normally off thyristor

structure (p+npn+) (Fig. 7) is turned on in a conventional

manner by injecting gate current of appropriate polarity

through the GTO gate while keeping the JFET off (i.e., a

negative voltage is applied to the JFET gate). This structure

has been inverted from the conventional silicon structure

due to the high resistance of the p+ SiC substrates. The turnon transition usually requires a small gate current, typically

1% to 2% of the anode current. The p-

 blocking layer is

used for forward blocking. However, unlike the GTO,

where the structure is turned off by reversing the gate

current on the GTO gate, the JCT is turned off by means of

turning on the n-channel JFET (JFET gate is returned to

0 V). This shorts out the forward-biased p+n junction, and

the current is bypassed through the JFET (Fig. 6). This

avoids complex gate drive snubbers currently employed to

turn off silicon GTO devices due to very high gate currents

GTO

Gate

External

JFET

Anode

Cathode

Fig. 6 - The JFET Controlled Thyristor (JCT) is an

ideal power switching device in SiC. The proposed

approach is based on the hybrid assembly using 4HSiC GTO and 4H-SiC vertical JFETs. The hybrid JCT

is turned off by diverting a sizable part of the main

current through the external JFET.

Low inductance, fast turn-off

Fig. 5 - A photograph

of a silicon MOS turnoff thyristor (MTO)

showing a hybrid

assembly of silicon

gate turn-off

thyristors (GTO) and

silicon MOSFETs

(photo courtesy of

Silicon Power

Corporation).

involved. This n-channel JFET should have a breakdown

voltage of about 25 V and current handling capability of

about 30% to 50% of the rated GTO current, depending

upon how fast one needs to turn the GTO off.

Fig. 7 - A schematic cross section of the SiC asymmetrical

GTO cell. The thickness of the p- base layer is adjusted for a

given breakdown voltage.

Thermal Considerations

The performance of present-day silicon devices is limited by

the thermal characteristics of devices and their packaging.

Silicon devices are limited by the junction temperature limit

of 150 °C (where typically a 125 °C temperature drop is

experienced between the junction and the ambient). On the

other hand, silicon carbide devices operating at 500 °C can

more than quadruple the power density of power control

modules. The removable heat flux can be increased by
18 Naval Research Reviews

450% simply by allowing the junction temperature to rise to

500 °C. Furthermore, significant reduction in the thermal

impedance of the packaged parts is possible as thermal

conductivity of SiC itself is three times better than Si at

offering lateral heat spreading.

The performance of any power device is ultimately limited

by the power dissipation within its structure, which determines the temperature rise:

∆T = TJ – TA = PDRθ

where PD is the power dissipation, TJ is the junction

temperature, TA is the ambient temperature (25 °C), and Rθ

is the thermal resistance. Typically, Rθ is dominated by the

package and for an advanced, double-sided, passively

cooled, bonded package aided by lateral spreading of heat in

SiC, we can expect an Rθ of 0.15 °C/W. Assuming a

junction temperature of 500 °C, we can dissipate about

3000 W/cm2. If the device is operated at 500 A/cm2 at 4 V

forward drop, the conduction losses will be only about

2000 W/cm2. Therefore, the switching frequency for a

given application will be determined by the switching losses

of the switch which, in turn, depend upon the turn-off time

of the switch. A higher switching frequency is useful in

reducing the size of the magnetic components in a typical

power conversion system. A switching frequency of 50 to

100 kHz is desirable but generally not achievable in silicon

devices.

Asymmetrical vs Symmetrical GTO

Figure 7 shows a schematic cross section of a 1-kV asymmetrical GTO cell [2,3]. The substrate is n+ 4H-SiC doped

with nitrogen to above 1 × 1019 cm-3. The p-type layer

consists of an ~1-μm p+ buffer layer doped with Al to about

1018 cm-3 and a 7-μm p- drift layer doped with Al at less

than 1 × 1015 cm-3. Next, there is a base layer of n-SiC, ~1

μm thick, doped with nitrogen at about 1018 cm-3. Finally,

the topmost layer is a very highly doped p+ layer, doped

with Al and about 0.7 μm thick. The width of the anode

stripe is in the 15 to 50 μm range. The reverse blocking

capability is minimal due to the presence of the p+ buffer

layer.

It turns out that the typical pulse width modulating (PWM)

inverter systems need only asymmetrical switches. It is

much easier to fabricate such devices because only one

blocking junction (J1) needs to be properly terminated. Our

design uses floating n+ guard rings to terminate this junction. This enables use of the multichip approach on the same

SiC wafer because now small cells can be individually

terminated. Symmetrical devices, however, will have two

blocking junctions — J1 and J2 — one on each side of the

drift layer, and no p-buffer layer. This means that symmetrical devices will have to be cut out of the wafer and beveled

on the edges to achieve edge termination. This makes the

multicell approach on the same host wafer almost impossible to implement as the cells tend to be small (~ 1 mm in

diameter). In addition, it is much easier to grow the lightly

doped, thick, p-type epitaxial layers required for asymmetrical devices than the precisely doped, half-as-thick p-type

layers required for symmetrical devices.

Experimental Results

Figures 8 through 10 show the results on various devices.

The “involute” structure has a diameter of 680 μm and

carries a forward current of 4 A (J~1100 A/cm2). The

breakdown voltage of this device is 600 V.

Fig. 8 - The “involute” design of SiC GTO

shows 600 V forward blocking voltage and 4 A

forward current at room temperature. The

blue glow indicates that the device is fully

turned on.

Figure 9 shows the forward current density vs forward drop

of SiC GTO at different temperatures. The measurement

was done with a gate current of 50 mA. The forward

voltage drop is higher than expected due to poor contact

resistance (~10-3 ohm • cm2) of Ti/Al alloy to p+ anode and

the resistance of the bulk n+ wafer (0.03 ohm • cm, 300 μm

thick). The forward current density of 500 A/cm2 is considered practical. Any higher value will result in excessive

conduction losses. The forward drop reduces at elevated

temperatures as expected. The turn-off transient was

measured on six interdigitated devices in parallel (Fig. 10).

The packaged device was turned off with a unit current gain

from 3 A forward conduction. The measurements indicate a

storage time of 130 ns and a fall time of 55 ns. These
Volume 51, Number 1 19

results clearly demonstrate the potential for a higher switching speed up to 50 to 100 kHz. The scale-up strategy, shown

in Fig. 11, entirely depends upon the expected improvements

in the materials technology. For now, smaller cells can be

connected in parallel to show large current handling capability. It is expected that within 2 to 3 years, the defect density

in the material will reduce to below 1 cm-2, thus enabling a

200 A single-cell-part.

Fig. 9 - Forward current density vs forward drop of SiC

GTO at different temperatures. The measurement was

done with a gate current of 50 mA.

Fig. 10 - Turn-off transient measured on six interdigitated

devices in parallel. The packaged device was turned off with

a unit current gain from 3 A forward conduction. The

results indicate a storage time of 130 ns and a fall time of

55 ns (courtesy of Silicon Power Corporation).

UMOSFET. MOSFET with

U-shaped gates.

1.8 mm, 10 A

4 mm, 50 A

5.6 mm, 100 A

8 mm, 200 A

JFET

GTO

Current

Fig. 11 - The GTO size projections with improvements in materials quality indicates that a 200-A

part may be possible within 2 to 3 years. The

current density has been assumed to be 500 A/cm2.

Power MOSFET Devices

In the area of power switching devices, vertical power

MOSFETs (Fig. 12) and IGBTs have been demonstrated [4-

10], but with more limited success because of the high

interface trap density (~5 × 1011 cm-2 • eV-1) and high fixedoxide charge (~1012 cm-2) for thermally grown gate oxides.

We have fabricated vertical UMOSFETs on 4H-SiC with

breakdown voltages up to between 1100 and 1400 V (Fig.

13).

Due to the smaller barrier heights for electron injection from

the SiC conduction band into the silicon dioxide conduction

band as compared to Si, the

time dependent dielectric

breakdown (TDDB) and hot

electron injection are

expected to be serious reliability issues. This is especially so

if high temperature (greater than 200 °C) operation is

considered [11]. This raises serious concerns about the

Fig. 12 - A schematic cross section of a unit cell of the SiC

UMOSFET showing high field points under on or off

conditions.

0 V 0 A

3 A 350 V

Time, 200 ns/div

6 interdigitated

devices in parallel

Storage time

130 ns

Fall time

~ 55 ns

T

T
20 Naval Research Reviews

reliability of SiC power MOS devices operating at elevated

temperatures. Long-term TDDB measurements are needed

to determine a maximum limit on the electric field across

the gate oxide in a 4H-SiC MOS system; however, it is clear

that this maximum electric field should be much lower than

that of silicon. It should be even lower for operation at

higher junction temperatures. This would adversely impact

the on-resistance of the inversion layer. Detailed FowlerNordheim measurements [11] indicate that at a given

temperature and electric field, the current density in 4H-SiC

MOS systems is about five times higher than that of 6H-SiC

systems because of the 4H-SiC MOS system’s smaller

effective barrier height. These reliability concerns make the

usefulness of power MOSFETs and IGBTs somewhat

questionable for high temperature operation.

DS

-1x10 -3

0x100

1x10 -3

2x10 -3

3x10 -3

4x10 -3

5x10 -3

0 200 400 600 800 1000 1200

IDS

 (A)

V (V)

VG = 60 V

VG = 50 V

VG = 55 V

VG = 0 V

4H-SiC UMOS

P13 ORTHO 20

January 21, 1997 

1100 V

Fig. 13 - Experimental or I-V characteristics

at 300 °C of a 4H-SiC UMOSFET indicating

1400 V breakdown voltage.

The SiC MOSFET devices have been found to have a very

low inversion layer electron mobility (1 to 10 cm2/Vs) as

evidenced by the I-V curves of Fig.13. This has generally

been attributed to high interface trap density near the

conduction band edge. Recently, we have discovered that if

ion-implanted regions are annealed above 1400 °C, the

surface of the SiC gets very rough due to the phenomenon

called “step-bunching.” The rms value of this roughness can

be a few hundred Å as shown in Fig. 14. This surface

roughness is responsible for the poor inversion layer

mobility of electrons in these devices. We have recently

made lateral devices on 6H-SiC with inversion layer

electron mobility of about 55 cm2/Vs [12] by avoiding stepbunching, as shown in the I-V plots of Fig. 15.

Fig. 14 - An atomic force microscopy (AFM) picture

of a SiC surface annealed at 1600 °C showing rms

roughness of about 100 Å.

SiC MOSFET (W/L=1E4/2 um) T=250C 

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

2

0 5 10 15 20 25 30

Vdrain(V)

Ids(A)

0V

+4V

+8V

+12V

+16V

+20V

Fig. 15 - The I-V characteristics of lateral

MOSFETs made on 6H-SiC. Oxide thickness

~340 Å, W/L = 1 cm/2 μm.

Conclusions

We have presented an overview of SiC power devices and

concluded that the JFET Controlled Thyristor (JCT) is the

most promising SiC switching device for high-power, hightemperature applications such as the All Electric Combat

Vehicle due to its ease of turn-off, potential for 500 °C

operation, and resulting reduction in cooling requirements.

In order to take advantage of SiC power devices, hightemperature packages and components with double-sided

contact need to be developed concurrently. Devices such as

MOSFETs and IGBTs will be unreliable at high temperatures due to fundamental issues such as charge injection and

trapping in the insulating layers. 6H-SiC is better in this

respect because it offers a higher barrier height for electron
Volume 51, Number 1 21

injection from conduction band into oxide and higher

inversion layer electron mobility. The phenomenon of

“step-bunching” explains the previously obtained low

surface electron mobilities. Finally, for future scaling up of

these devices, it is important to reduce not only micropipes

but the dislocation density from the low current levels of 104

cm-2 by at least an order of magnitude as well.

Acknowledgments

The authors thank the staff of the Central Processing

Laboratory at the Northrop Grumman STC, Northrop

Grumman Corporation, and several U.S. government

agencies (DARPA, ONR, and the Air Force) for their

support of this work.

References

1. A.K. Agarwal et al., “SiC Electronics,” IEDM Tech. Dig.,

1996.

2. A.K. Agarwal, J.B. Casady, L.B. Rowland, S. Seshadri, R.R.

Siergiej, W.F. Valek, and C.D. Brandt, “700-V Asymmetrical

4H-SiC Gate Turn-Off Thyristors (GTO’s),” IEEE Electron.

Dev. Lett. 18(11), 518-520 (1997).

3. S. Seshadri, B. Casady, A.K. Agarwal, R.R. Siergiej, L.B.

Rowland, P.A. Sanger, C.D. Brandt, J. Barrow, D. Piccone, R.

Rodrigues, and T. Hansen, “Turn-Off Characteristics of 1000

V SiC Gate Turn-Off Thyristors,” presented at the 10th

International Symposium on Power Semiconductor Devices

and ICs (ISPD’98), Kyoto, Japan, June 3-6, 1998.

4. A.K. Agarwal et al., “Critical Materials, Device Design,

Performance and Reliability Issues in 4H-SiC Power

UMOSFET Structures,” Mat. Res. Symp. Proc. 423, 87-92

(1996).

5. W. Palmour, J.A. Edmond, H.S. Kong, and C.H. Carter, Jr.,

“Vertical Power Devices in Silicon Carbide,” Silicon Carbide

and Related Materials, Institute of Physics Conf. Ser. 137,

499-502 (1994).

6. J.N. Shenoy, M.R. Melloch, and J.A. Cooper, Jr., “HighVoltage, Double-implanted Power MOSFETs in 6H-SiC,”

Device Research Conference, late paper, 1996.

7. N. Ramungul, T.P. Chow, M. Ghezzo, J. Kretchmer, and W.

Hennessy, “A Fully Planarized, 6H-SiC UMOS InsulatedGate Bipolar Transistor,” Device Research Conference, pp.

56-57, 1996.

8. J. B. Casady, A.K. Agarwal, L.B. Rowland, W.F. Valek, and

C.D. Brandt, “900 V DMOS and 1100 V UMOS 4H-SiC

Power FETs,” presented at the Device Research Conference,

Colorado State University, Fort Collins, Colorado, June 23-

25, 1997.

9. A.K. Agarwal, J.B. Casady, L.B. Rowland, W.F. Valek, M.H.

White, and C.D. Brandt, “1.1 kV 4H-SiC Power

UMOSFET’s,” IEEE Electron. Dev. Lett. 18(12), 586-588

(1997).

10. A.K. Agarwal, J.B. Casady, L.B. Rowland, W.F. Valek, and

C.D. Brandt, “1400 V 4H-SiC Power MOSFETs,” Proceedings of the International Conference on SiC, III-Nitrides and

Related Materials, Stockholm, Sweden, August 1997.

11. A.K. Agarwal, S. Seshadri, and L.B. Rowland, “Temperature

Dependence of Fowler-Nordheim Current in 6H- and 4H-SiC

MOS Capacitors,” IEEE Electron. Dev. Lett. 18(12), 592-594

(1997).

12. N.S. Saks, A.K. Agarwal, and S.S. Mani, “Characterization of

Electron Mobility in 6H-SiC MOSFETs,” presented at the 29th

IEEE Semiconductor Interface Specialists Conference, San

Diego, California, December 3-5, 1998.
22 Naval Research Reviews

The Authors

Anant K. Agarwal received a B.E. degree in electrical engineering from M.N.R. Engineering College,

Allahabad, India, in 1978, an M.S.E.E. degree from the University of Tennessee in 1981, and a Ph.D. in electrical

engineering from Lehigh University, Bethlehem, Pennsylvania, in 1984. In 1984, he joined AT&T Bell Laboratories

in Murray Hill, New Jersey, where he worked on GaAs digital circuits for optical communication. From 1985 to

1990, he was an Assistant Professor at the M.N.R. Engineering College, Allahabad, India. From 1990 to 1999,

he was with Northrop Grumman Science & Technology Center (STC), Pittsburgh, Pennsylvania, where he was

engaged in the development of SiGe HBTs and SiC high-power devices. Since March of this year, he has been

with CREE Research, Durham, North Carolina, where he is developing SiC high-power, high-temperature

devices. Dr. Agarwal has coauthored more than 60 technical papers and conference presentations and holds

four U.S. patents.

Sita S. Mani is currently working at Sandia National Laboratories in the area of W-CVD as a senior

member of the technical staff in the Microsystems Science, Technology and Components division. Previously,

she worked at Northrop Grumman, Pittsburgh, Pennsylvania, fabricating SiC power and microwave devices

and assisting in epitaxial growth efforts. She received her Ph.D. from Rensselaer Polytechnic Institute in Troy,

New York, in 1994.

Suresh Seshadri received a B.S. degree in electrical engineering from Columbia University and a Ph.D.

in electrical engineering from Yale University. His Ph.D. dissertation examined mechanisms for compositional

disordering of AlGaAs quantum well heterostructures. In 1995, he became a senior engineer at Westinghouse

Science & Technology Center (W-STC) where he was part of a group responsible for developing the first

practical SiC neutron detectors for high radiation environments, an achievement for which the group received

a 1997 Westinghouse Signature Award for Excellence. Continuing work on the development of SiC/GaN diodes

for wideband RF amplification at W-STC after its sale to Northrop Grumman Corporation in 1997, he has since

been primarily involved in the development of SiC diodes, Gate Turn-Off (GTO) and MOS Turn-Off (MTO)

thyristors, and the Junction Controlled Thyristor (JCT) for high-temperature, high-frequency, and high-power

switching applications. This work has recently culminated in the first demonstration of an all-SiC Pulse Width

Modulated (PWM) inverter. He has authored/coauthored 22 publications.

Jeffrey B. Casady received his B.S.E.E. and his M.S.E.E. from the University of Missouri, Columbia, in

1990 and 1992, respectively, and his Ph.D. in electrical engineering from Auburn University in 1996. His Ph.D.

dissertation was performed on advanced SiC unit process development and device characterization/

implementation. After receiving his Ph.D., he worked for approximately 2-1/2 years at Northrop Grumman

Corporation, Electronic Sensors & Systems Sector, as a senior engineer. While at Northrop Grumman, he

researched advanced SiC power and microwave device design and processing. He designed and processed

SiC GTOs, UMOSFETs, DMOSFETs, and SITs for a variety of externally and internally sponsored programs.

He then joined Mississippi State University as an Assistant Professor of Electrical Engineering in January 1999

to continue his research into wide bandgap semiconductor devices. Over the past five years, principal publications

include more than 35 technical articles relating to SiC devices, including one book chapter. In addition to his

technical activities, Dr. Casady has been a member of IEEE for 11 years, and is a member of both the IEEE

Electron Device Society and IEEE Power Electronics Society, as well as a member of Eta Kappa Nu. In 1997,

he received a Northrop Grumman award for timely contributions in the development and fabrication of SiC

UMOSFETs.
Volume 51, Number 1 23

Phillip A. Sanger received his B.A. degree in physics from Saint Louis University in 1970. Following a

Fulbright year of study in nuclear engineering at Saclay, France, he resumed his graduate studies at the

University of Wisconsin, Madison, Wisconsin, where he received his M.S. in 1974 and, ultimately, his Ph.D. in

1977. He performed his thesis research at Fermilab in Batavia, Illinois, on the effects of 400 GeV protons on

metals. From 1977 to 1989, he pursued the development of superconductors and superconducting MRI systems

at Oxford Superconducting Technology, Carteret. Following a brief period with the Superconducting SuperCollider

Laboratory in Dallas, Texas, between 1989 and 1992, he has been engaged in managing the development of

new technologies at the Northrop Grumman Science and Technology Center in Pittsburgh, Pennsylvania.

Since 1996 he has pursued SiC power devices for military and commercial power conditioning applications.

Charles D. Brandt has been working in the semiconductor field for the last 19 years in three primary areas

of research: growth and characterization of GaAs and InP crystals, III-V integrated circuit fabrication, and wide

bandgap (SiC, GaN) materials and device development. His early work on bulk crystals was directed at the

behavior of the transition elements V, Ti, and Sc as dopants in GaAs, InP, and AlGaAs and involved the use of

crystal growth techniques such as liquid-encapsulated Czochralski and horizontal Bridgman in addition to

extensive electrical and optical characterization. As a staff member at Raytheon’s Research Division in the

early 1980s, he was involved in the formulation of plasma deposition and etching techniques suitable for GaAs

FET fabrication. In recent years, Dr. Brandt’s work at Northrop Grumman has centered around the development

of advanced SiC-based devices covering the range from DC to Ka band, where he has managed development

programs for SiC SITs, MESFETs, MOSFETs, IMPATTs, GTOs, and diodes. He has also been extensively

involved in the development of advanced processing and packaging techniques for monolithic microwave

integrated circuits utilizing wafer-scale integration concepts. Dr. Brandt has over 50 technical publications in

these areas. Currently, Dr. Brandt manages the Microelectronics Department at the Northrop Grumman Science

& Technology Center in Pittsburgh, Pennsylvania.

Nelson Saks received his B.A. degree in physics from Amherst College in 1968 and his M.S. degree in

physics from the University of Maryland in 1973. He has held the position of research physicist at the Electronics

Science and Technology Division of the Naval Research Laboratory, Washington, D.C., from 1968 to the present.

He has worked on amorphous semiconductors, development of charge-coupled devices, and MOS devices

with an emphasis on electrical characterization techniques and radiation effects. Mr. Saks has enjoyed sabbaticals

at IMEC, Leuven, Belgium, where he worked on hot carrier injection in VLSI devices (1984-1985) and at

Northrop Grumman Corporation, Pittsburgh, Pennsylvania, where he worked on SiC power devices. He has

been active in the radiation effects community and was technical program chairman of the 1992 IEEE Nuclear

and Space Radiation Effects Conference. He served as advisor to the Defense Nuclear Agency from 1988 to

1995 on the basic mechanisms of radiation damage in electronic devices. His present interests include hot

carrier injection phenomena in MOSFETs; techniques for characterizing MOS devices, especially charge

pumping; growth and characterization of MOS gate dielectrics; basic mechanisms of radiation damage in MOS

structures; MOS devices on SiC and GaN; and lifetime in MOS power devices. He has authored or coauthored

more than 100 publications on these topics. Mr. Saks was elected Fellow of the IEEE in January 1996 for

“contributions to the understanding of the basic mechanisms of radiation damage in metal-oxide-semiconductor

devices.”
24 Naval Research Reviews

rather, it is a modulation-doped field effect transistor, a kind of

metal-oxide semiconductor FET in which a thin layer of a nitride

semiconductor is deposited on an undoped, and therefore electrically highly

resistive, substrate. The alloy nitride confines the electrons to a quantum

mechanical well that enhances their ability to move about in the plane, thus

increasing the device performance. In MODFETs, GaN can go where no

semiconductor has gone before, and is therefore in increasing demand in high

power, high temperature applications. This is because of its large bandgap, high

breakdown voltage, high thermal conductivity, and high electron velocities. Nitride

semiconductors are also sought after for use in optoelectronic devices that

operate in the UV and visible wavelengths. When used as UV sensors, GaN

devices have shown tremendous promise in applications that range from

maximizing fuel efficiency to detecting threats aimed at aircraft. In MODFETs,
Volume 51, Number 1 25

 A field effect transistor (FET) is built by molecular beam epitaxy (MBE) from two layers of semiconducting material. The top layer is active (current can flow), and the bottom layer is semi-insulating.

The top layer is composed of a source (of electrons), a gate, and a drain. When voltage is applied to the

gate, this material allows electrons to flow from the source to drain, thus amplifying voltage.

GaN has demonstrated record power levels at high frequencies. The nitrides

(GaN and its heterostructures with InGaN and AlGaN) used as light-emitting

diodes (LEDs) and laser diodes (LDs) are already being applied to scores of

applications as diverse as computers and traffic signals; they are excellent candidates

for an even wider diversity of technologies that include printing, agriculture, and

health care.

In this article, Dr. Hadis Morkoç of Virginia Commonwealth University, one

of the most cited physicists in the world because of his groundbreaking research

in many semiconductor composites, including wide bandgap semiconductors,

delivers a state-of-the-art address on the latest research and development of

GaN-based MODFETs and UV detectors. Besides exploring their promising

potential, he explains the science behind how they do what they do.

– S.O.

GaN-based semiconductors can provide

RF power to shipboard radars.

They can also be used as UV detectors

 to protect aircraft against hostile missile threats.
26 Naval Research Reviews

Hadis Morkoç

Virginia Commonwealth University

Richmond, Virginia

GaN-Based Modulation-Doped

FETs and UV Detectors

Introduction

Semiconductor nitrides such as aluminum nitride (AlN),

gallium nitride (GaN), and indium nitride (InN) are very

promising materials for their potential use in optoelectronic

devices (both emitters and detectors) and high power/

temperature electronic devices as have been treated in length

and reviewed recently [1-5]. These materials and their

ternary and quaternary alloys cover an energy bandgap

range of 1.9 to 6.2 eV, suitable for band-to-band light

generation with colors ranging from red (potentially) to

ultraviolet (UV) wavelengths. Specifically, nitrides are

suitable for such applications as surface acoustic wave

devices [6], UV detectors [7,8] Bragg reflectors [9],

waveguides, UV and visible light emitting diodes (LEDs)

[10-12], and laser diodes (LDs) [13] for digital data readwrite applications. During the last several decades, lasers

and LEDs have expanded remarkably both in terms of the

range of emission wavelengths available and brightness. The

nitride semiconductor-based LEDs have proven to be

reliable in such applications as displays, lighting, indicator

lights, advertisement, and traffic signs/signals; possible

applications include use in agriculture as light sources for

accelerated photosynthesis, and in health care for diagnosis

and treatment. Lasers, as coherent sources, are crucial for

high-density optical read and write technologies. Because

the diffraction-limited optical storage density increases

approximately quadratically as the probe laser wavelength is

reduced, nitride-based coherent sources at wavelengths

down to UV are attracting a good deal of attention. Optical

storage would enable the storage and retrieval of inordinate

numbers of images and vast quantities of text with untold

efficiency. Other equally attractive applications envisioned

include printing and surgery.

When used as UV sensors in jet engines, automobiles, and

furnaces (boilers), the devices would allow optimal fuel

efficiency and control of effluents for a cleaner environment.

Moreover, UV sensors that operate in the solar blind region

(260 to 290 nm) would have high detectivity because the

ozone layer absorbs solar radiation at those wavelengths,

thus virtually eliminating radiation noise. Consequently,

these detectors are expected to play a pivotal role in threat

recognition aimed against aircraft [7,8,14]. GaN/AlGaN UV

p-i-n detectors have demonstrated sensitivities of about

0.12 A/W and response speed of a few nanoseconds in

250 μm dot sizes; both values represent a quantum leap over

previous capability.

GaN’s large bandgap, large dielectric breakdown field,

fortuitously good electron transport properties, and good

thermal conductivity are conducive for use in high power

electronic devices [15]. Sheppard et al. [16] have already

reported that 0.45 mm gate, high power modulation-doped

FETs (MODFETs) on SiC substrates exhibited a record

power density of 6.8 W/mm in a 125 μm-wide device and a

record total power of 4 W (with a power density of

2 W/mm) at 10 GHz. Other groups have also reported on the

superior performance of GaN-based MODFETs on SiC and

sapphire substrates with respect to competing materials,

particularly at X band and higher frequencies [17-19].

Applications include use in amplifiers operative at high

temperature and in unfriendly environments as well as in

low-cost compact amplifiers for earthbound and space

applications.

Nitride semiconductors have been deposited by vapor phase

epitaxy (i.e., both hydride VPE [HVPE] and organometallic

VPE [OVPE]), and in a vacuum by molecular beam epitaxy

(MBE). With its innate refined control of growth parameters, in-situ monitoring capability, and uniformity, MBE is

well suited for depositing heterostructures and gaining

insight to deposition/incorporation mechanisms. MBE’s

control over growth parameters is such that any structure

can be grown in any sequence. The structures for IR lasers

in CD players, surface emitting vertical cavity lasers, and

high-performance electronic devices such as pseudomorphic

MODFETs have all been produced very successfully, most
Volume 51, Number 1 27

WURTZITE. A three-dimensional

hexagonal crystalline structure (such

as the blade of a hex wrench).

of them commercially, by MBE. Nitride growth, however,

requires much higher temperatures than those used in

producing the conventional Group III-V semiconductors for

which the MBE systems were designed. In addition, it has

proved difficult to provide active N species at sufficiently

high rates for nitride growth. Despite these mechanical/

engineering limitations and its relatively late entry, MBE

has already played a key role on a number of fronts, such as

high performance GaN-based MODFETs and fast solar

blind detectors, which are discussed in this article.

Modulation-Doped Field Effect Transistors

With its reduced impurity scattering and unique gate

capacitance-voltage characteristics, the MODFET has

become the dominant high frequency device. Among the

MODFET’s most attractive attributes are close proximity of

the mobile charge to the gate electrode and high drain

efficiency. As in the case of emitters, the GaN-based

MODFETs have quickly demonstrated record power levels

at high frequencies.

In MODFETs, the carriers that form the channel in the

smaller bandgap material are donated by the larger bandgap

material, ohmic contacts, or both. Since the mobile carriers

and their parent donors are spatially separated, short-range

ion scattering is nearly eliminated, which leads to mobilities

that are characteristic of nearly pure semiconductors. A

Schottky barrier is then used to modulate the mobile charge

that in turn causes a change in the drain current. Because of

this heterolayer construction, the gate can be placed very

close to the conducting channel, resulting in large

transconductances [20]. Figure 1 presents a schematic

representation of a GaN/AlGaN MODFET structure in

which the carriers are provided by the donors in the wider

bandgap AlGaN. In a MODFET device, the carriers can also

be provided by the source contact.

Table 1 - Piezoelectric Constants and Spontaneous

Polarization Charge in Nitride Semiconductors

AIN

1.46

–0.60

–0.081

–0.86

e33* [C/m2**]

e31* [C/m2]

P0 [C/m2

]

[e31-[c31/c33***]e33]

* e31 and e33 are piezoelectric constants

** C/m2 is coulombs per square meter

*** c31 and c33 are elastic constants

GaN

0.73

–0.49

–0.029

–0.68

InN

0.97

–0.57

–0.032

–0.90

Fig. 1 - Conduction band structure of a modulationdoped structure.

x

E1

E0

Ec Ec

EF2 EF

x = x x = 0 2

ED

∆Ec

di

EFi

Polarization

Polarization charge, however, affects device design and

operation in all nitrides, particularly nitride-based

MODFETs. Polarization charge arises from two sources:

piezoelectric (PE) effects and the difference in spontaneous

polarization between AlGaN and GaN, even in the absence

of strain. Spontaneous polarization has only recently been

fully understood. As shown in Ref. 21, nitrides lack inversion symmetry and exhibit piezoelectric effects when

strained along the [0001] direction. Piezoelectric coefficients in nitrides are almost an order of magnitude larger

than in traditional Group III-V semiconductors [22]. In

addition, wurtzite GaN has

a unique axis, thus allowing

spontaneous polarization

(P0, whose values are given

in Table 1) even in the

absence of strain. This can

manifest itself as polarization charge at hetero-interfaces.

Reference 23 provides a review of polarization effects.

Let us compare the relative importance of spontaneous

polarization to piezoelectric polarization. For a biaxially

strained layer, the effective piezoelectric polarization is

given by

 Pz

piezo= [e31-(c31/c33)e33]ε⊥, (1)

where ε⊥ = εxx+ εyy is the in-plane strain and c31 and c33 are

elastic constants.

For AlxGa1-xN pseudomorphically strained on a relaxed

GaN substrate, the strain ε⊥ is expected to be proportional to

x and given by ε⊥ = 2x(aGaN-aAlGaN)/aAlGaN, which is 0.051x

and is tensile. The piezoelectric polarization is then

Ppiezo = –0.044x, i.e., pointing in the [0001] direction. The

corresponding difference in spontaneous polarization

between AlxGa1-xN and GaN is also expected to be proportional to x and is given by ∆Pspon = -0.049x. So, the two

work in the same direction and are comparable in magni-
28 Naval Research Reviews

tude. Note that these are all in C/ m2 and that 1C/m2 =

0.624×1015 electrons/cm2. So, for x of order 0.1, we are

dealing with carrier concentrations of order 1012-1013.

For an InxGa1-xN layer, the situation is different: now, the

difference in spontaneous polarization is much smaller,

∆Pspon = -0.003x. Furthermore, the InxGa1-xN layer on GaN

would be under compressive stress ε⊥ = -0.203x, and

Ppiezo = +0.183x. Here the piezoelectric polarization

dominates and is opposite in direction but even larger in

absolute magnitude.

In the AlGaN case, the sign of the polarization is such as to

produce a potential energy for electrons sloping down from

the Ga face towards the N face. Thus, for a structure in

which the Ga face is turned towards the surface, the potential will slope down from the AlGaN surface towards the

AlGaN/GaN interface and helps to drive carriers towards the

two-dimensional electron gas (2DEG) forming at this

interface. For example, if there is an ohmic metal contact on

the AlGaN surface, electrons will flow towards the 2DEG

below that layer.

The most favorable situation for enhancing sheet carrier

concentration would occur for an InGaN quantum well on

top of relaxed n-GaN and below a AlGaN barrier with the

whole structure having cation polarity towards the surface.

In that case, the field will slope down towards the InGaN/

AlGaN interface in the quantum well and will help localize

the carriers in the 2DEG. Note that the piezoelectric

polarizations estimated here are based on the theoretical

values for a perfectly insulating material. The field is

screened by the carriers present in each layer. For example,

if carriers flow from a metal contact towards the 2DEG, then

this will set up a counteracting field. The self-consistent

field is ultimately determined by the condition that the

chemical potential for electrons (i.e., the Fermi level) must

be constant throughout the structure and thus depends on the

doping and band bending in the substrate and, possibly, in

each of the layers. At the least, one may expect these fields

to be reduced by a factor corresponding to the macroscopic

dielectric constant, i.e., a factor of order 10 but possibly

larger if the layers acquire conductivity by free carriers. So,

a more realistic expectation for the effects on sheet carrier

concentration is of order 1011-1012 electrons/cm2.

The difference between these and traditional device structures without polarization effects is that for uniform dopant

concentrations, one obtains parabolically varying potentials

with distance, whereas here the linear terms come from

polarization on top of the parabolic terms. These linear

terms lead to variations of the potential over a shorter

distance scale determined by the thickness of the layers,

whereas the parabolic terms correspond to the space-charge

regions. Thus the linear terms may help to localize carriers

if the polarity of the structure is chosen properly.

Some further words of caution about the above estimates are

needed. If the AlGaN layers are not pseudomorphic but

partially relaxed (by misfit dislocations for example), then

the piezoelectric effect would be reduced but the spontaneous polarization would still be present. If the interfaces are

not atomically sharp but exhibit a certain degree of interdiffusion, the differences in spontaneous polarization would be

reduced as well. Finally, if domains with inverted polarity

exist, the overall polarization effects may be washed out.

Also note that in an inverted structure with nitrogen (N)

polarity towards the surface, it may be possible to create a

two-dimensional hole gas (2DHG) at the AlGaN/

substrateGaN interface, providing that free holes are

available. However, if an n-type GaN layer is placed on top,

a 2DEG may form on top of the AlGaN layer.

The immediate impact of this polarization is that the field

generated by this process must be considered together with

that induced by the applied voltage and charge redistribution. Moreover, as alluded to earlier, free carriers can also be

drawn from any shallow and weakly bound impurities and

metal in contact with the semiconductor. In any case, the

free carriers would tend to screen the piezoelectric-induced

polarization field. An additional complicating factor in

nitrides in relation to polarization is that the semiconductor

tends to twist and tilt in a columnar mode, in an effort to

minimize strain as shown in Fig. 2. These columns do not

have the same cation/anion ordering polarity as shown in

Fig. 2. In the presence of strain, Ga polarity domains and N

polarity domains would have opposite polarization, causing

increased scattering [24-26].

Figure 3 is a schematic representation of an ideal inversion

domain boundary formed in growth along the [0001]

direction. On the left of the boundary, the growth initiates

Fig. 2 - Domains in GaN, with N polarity (nitrogen surface

layer) on the right and Ga polarity (with Ga on surface)

on the left side under compressive residual strain. The

arrows show the direction of the E field in each of the

domains.

E E

Ga polarity N polarity

sapphire substrate
Volume 51, Number 1 29

with N, and on the right it begins with Ga. On the left side,

the bond along the [0001] direction is from Ga to N; this is

called Ga polarity. On the right side, the [0001] bonds are

from N to Ga; this is called N polarity. In N polarity and

under tensile strain, the PE field generated points toward the

surface, whereas that for the Ga polarity region points in

from the surface. When the strain is compressive, the

direction of the field changes.

Yet an additional complicating factor is the asymmetry in

the barrier discontinuities between GaN and its binary and

ternaries caused by spontaneous polarization [27]. The

spontaneous polarization arises simply because of the

ionicity of the bonds and the low symmetry in wurtzite. In

fact, Bernardini et al. [28] showed that the field that occurs

in quantum wells is almost completely determined by both

the difference in spontaneous polarization between the two

bulks, and their PE contribution. The field (i.e., the slope of

the potential) is quite independent of the offset (i.e., the

dipole discontinuity that occurs at the interface between the

two materials). In other words, the dipole at the interface

appears to be independent from the charge (monopole)

accumulated at the interface. Recent experimental results

obtained in the author’s laboratory with colleagues from the

University of Lecce support this theory in that the spontaneous polarization component is dominant.

MODFET Description

Simplified analytical descriptions of MODFET operation

have been developed which show quantitatively the effect of

charge stored at the hetero-interface on mobility and carrier

velocity. The analytical model of Drummond et al. [29]

describes MODFET operation well. The model provides

analytical expressions that relate the channel charge to the

gate and drain potentials and predict the current voltage

characteristics. In short, the MODFET provides large

mobile carrier concentrations in close proximity to the gate,

which was previously unachievable. A similar model [15,30]

has been developed for GaN/AlGaN MODFETs and used to

predict their performance. In short, the favorable velocity

field characteristics of GaN point to performance that is

competitive with GaAs even in terms of small signal

properties. In terms of power performance, GaAs devices

are supplanted by GaN devices with increasing margin as

the nitride technology evolves. There are, however, anomalies such as high- frequency current collapse/degradation

that must first be understood and circumvented before the

intrinsic properties of GaN can be fully exploited.

In addition to the small signal description, an accurate

modeling of device power is very useful. In high power

semiconductor devices, it is imperative that the effect of

temperature on device performance is accounted for

accurately. As in small signal modeling, the first step in

power modeling is to establish the basic device geometrical

factors that are needed to calculate the current voltage

characteristics. Once these are known, the output characteristics superimposed with the load line can be used to

estimate the power level that can be obtained from the

device provided that it is not limited by the input drive as

shown in Fig. 4. In Class A operation, the maximum power

that can be expected from the drain circuit of a device is

given by

 8

( ) max

dson Vb Vknee I P − = (2)

where Idson is the maximum drain current (this is the drain

current with a small positive voltage on the gate electrode),

Vb is the drain breakdown voltage, and Vknee is the knee

voltage as shown in Fig 4. The allowable positive gate

voltage (≈ 1V) will depend on the channel doping and the

work function of the gate metal. The positive gate voltage is

limited by the onset of forward Schottky-diode current. The

DC load line shown in Fig. 3 would be used in a Class A RF

amplifier with the maximum drain voltage Vd = Vb/2. The

slope of the load line is 1/RL where RL is the value of the

load resistance at the output of the FET. What can be

gleaned from Eq. (2) is that Vb and Idson must be made as

large as possible. The utility of wide bandgap semiconductors such as GaN in this juncture is that the drain breakdown

voltage is larger than that in conventional Group III-V

semiconductors. In general, the drain can be swung to

Fig. 3 - Schematic representation of an inversion

domain boundary formed in growth along the

[0001] direction.

Ga

N

Ga polarity Nitrogen polarity

Inversion domain boundary

Tensile

Tensile

Compressive

Compressive

E field direction E field direction

Substrate

c-direction

[0001] [0001]
30 Naval Research Reviews

voltages up to within 80% of the drain breakdown for a 20%

margin of safety. It should be pointed out that the maximum

drain current in nitride semiconductor-based MODFETs is

in the same ballpark as that of more conventional semiconductors. This implies that increased power handling capability is a direct result of large breakdown voltages and thermal

conductivity and the fact that higher junction temperatures

can be tolerated. Ability to increase drain bias increases the

load resistance and makes it easier to impedance match,

particularly in devices with large gate widths.

In power devices, power dissipation within the device

increases the junction temperature and alters the output

characteristics. On the one hand, higher junction temperatures with respect to the case temperature would enhance the

heat dissipation to the power of four of the temperature

differential, but along with it come reduced current and

increased series resistances, which in turn increase the heat

dissipation. Moreover, the thermal conductivity of the

semiconductor decreases with increased temperature,

exacerbating the situation. Consequently, the effect of

junction temperature on the output characteristics must be

taken into consideration. Temperature-dependent material

parameters, if known, can be used to calculate the output

characteristics with respect to temperature. However, a more

pragmatic approach, particularly when the aforementioned

parameters and or models required are not available, can be

taken in which one measures the output characteristics of

the device under consideration as a function of temperature.

The junction temperature is critically dependent on the

substrate thermal conductivity that is available for various

substrates including GaN [1]. The functional dependence of

thermal conductivity on temperature is

 χ (T) = χ (T0 ) (T/T0)

-r

 , (3)

where the coefficient r is 0.559, 0.443, 0.524, and 0.544 for

Si, GaAs, SiC, and sapphire, respectively [31]. Thermal

conductivity of sapphire, SiC, GaAs, and Si as a function of

temperature are shown in Fig. 5. In Fig. 5, χ (T0) has also

been appropriately reduced to account for the doping of the

substrate material.

Schottky Barriers for Gates

Any semiconductor device requires metal contacts and

MODFETs are no exception. These devices require ohmic

source and drain contacts as well as a rectifying Schottky

barrrier for controlling the charge in the channel. Schottky

barrier-related processes for GaN-based devices are nascent,

but rapid progress is being made. Until recently it has been

difficult to fabricate good quality single-crystal films on

which a Schottky metal could be deposited, and upon which

the properties of Schottky barriers could be studied.

However, considerable progress has been made with Pt-GaN

Schottky barriers [32,33] that have been successfully

implemented in GaN-based MODFETs [15,34-38].

Recent successes in growing good quality single-crystal

Group III-V GaN layers prompted the fundamental electrical property studies of metal-semiconductor barriers on

GaN. In order to determine the properties of only the metalsemiconductor junction, one must be able to model the

semiconductor. Semiconductors with large defect concentrations are notorious for exhibiting parasitic processes in

current-voltage and capacitance-voltage characteristics that

cloud the picture. Consequently, good epitaxial layers as

well as good metal-semiconductor interfaces are imperative.

During the evolutionary period, while the sample quality is

acceptable, temperature and frequency dependence of the

capacitance-voltage characteristics and temperaturedependent current-voltage characteristics are measured and

Fig. 4 - Pentode-like output I-V characteristics with the

load line. The loss due to the on resistance is also

indicated.

Slope = 1/Ron

I2 R Loss

DC Load line (slope = 1/RL Vknee )

Idss

Idson

Vd xVb

Drain Voltage x=0.8

Drain Current

Fig. 5 - Thermal conductivity vs temperature for

SiC and sapphire [31].

10

103 102 10

Thermal Conductivity (W/cmK)

Temperature (°C)

1.0

0.1

Sapphire

SiC

Si

GaAs
Volume 51, Number 1 31

analyzed for determining effective metal-semiconductor

barrier height. To get a large Schottky barrier height for

rectifying metal contacts on GaN, which is imperative for

low leakage, metals with large work functions such as

Au [39] and Pt [32] have been explored. Hacke et al. [39]

have studied Schottky barriers made of Au on unintentionally doped n-GaN grown by HVPE. The forward current

ideality factor was nidl ~1.03 and the reverse bias leakage

current was <10-10 A at a reverse bias of -10 V. While the

current-voltage measurement indicated the barrier height to

be 0.844 eV, the capacitance measurements led to a value of

0.94 eV.

Suzue et al. [32] have studied the Pt Schottky barriers on

unintentionally doped n-GaN. Temperature- dependent

current-voltage and capacitance-voltage characteristics in

the range of -195 °C to 42 °C were studied to gain insight

about the current conduction mechanism. Any excess

current observed is traditionally attributed to defects

(generation recombination centers) and surface leakage

current. The ensuing current is called the Shockley-ReadHall (SRH) recombination current resulting from the midgap states. If one neglects this excess current, a barrier

height of about 0.8 eV is deduced as opposed to about 1 eV

deduced from the CV measurements. Because of the effect

of excess current on the slope of the I-V curve, C-V

measurements in this particular case may represent the

metal barrier height. An examination of the C-V plots,

however, indicated that under reverse bias condition, the

capacitance depended insignificantly on the density of traps.

The curves corresponding to all temperatures were largely

linear, which yielded barrier heights ranging between 0.95

and 1.05 eV. Reduced capacitance with decreasing temperature is consistent with relatively deep donors. Binari et al.

[40] determined Ti Schottky barriers heights to be 0.58 and

0.59 eV from the current-voltage and capacitance measurements, respectively. The ideality factor nidl is approximately

1.28. The diode series resistance (Rs

) is 100 W.

The ternary AlxGa1-xN is an essential component of nitridebased MODFETs, which makes the investigation of metal

AlxGa1-xN contacts imperative. M. A. Khan et al. [41]

reported the fabrication of a Cr/Au Schottky barrier on

n-AlGaN. Moreover, M.R.H. Khan et al. [42] studied the

Schottky barrier characteristics of the Au-AlxGa1-xN system.

A typical current-voltage characteristic of an Al0.14Ga0.86N

Schottky diode had an ideality factor of 1.56 under reverse

bias and a threshold voltage of about 0.9 V at 0.1 A. The

reverse bias leakage current was recorded to be marginally

low (10-10 A) for a reverse bias of -10 V. By using the

current-voltage method, the barrier height and electron

affinity were determined to be 0.94 eV and 4.16 eV,

respectively. From the C-2 vs V plot, the same barrier height

and the electron affinity were deduced to be 1.3 ± 0.05 eV

and 3.8 eV, respectively. As the AlGaN quality increases,

periodic as well as more in-depth investigations will become

necessary to get an accurate picture of intrinsic parameters. In

short, the current conduction mechanism in metal-semiconductor structures is strongly affected by surface and bulk

states. Deviations from an ideal ideality factor, such as is the

case here, indicate such states. The situation gets more

complicated with AlGaN and gets worse as the AlN mole

fraction is increased. Likewise, capacitance-voltage measurements also are affected by states that are charged, either by

interface state or by bulk state. As is the case in many facets

of research and development, insights into the metal-nitride

contacts will be gained in an evolutionary manner hinging

upon the developments in nitride layers.

Contacts to GaN

Ohmic contacts in power devices are extremely important

because they affect their efficiency as well as heat dissipation.

Initial inferior results helped fuel concerns that GaN-based

electronic devices may not perform well. Early specific

contact resistivities on n-type GaN using Al and Au metallizations [43] were in the range of 10-4 and l0-3 Ω cm2. Major

improvements were realized by using Ti/Au [44] and TiAl

[45], in that specific contact resistivities in the high

10-6 Ω cm2 were obtained with the latter. Carrying the TiAl

contact work one step further, Wu et al. [46] confirmed that,

except at very high annealing temperatures, the ohmic contact

suggested by Lin et al. [45] functions very effectively. At very

high temperatures, Al of the metal contact melts and tends to

ball up, resulting in rough surfaces and increased ohmic

contact resistances as pointed out already by Lin et al. [45]. In

an attempt to circumvent this difficulty, Wu et al. [46]

designed a separate layer-metallization method where a

realignment and deposition of a second thin Ti layer, and a

2000 Å Au overlayer were carried out. Specific contact

resistivities were in the range of 3.0 × 10-6 Ω cm2 and

5.5 × 10-6 Ω cm2, depending on the doping concentration in

the semiconductor.

In an attempt to obtain improved ohmic contacts, Fan et al.

[47] have designed a multilayer ohmic contact method. By

using a composite metal layer of Ti/Al/Ni/Au (150Å/2200Å/

400Å/500Å), they obtained very low contact resistivities.

Specifically, for n-GaN with doping levels between

2 × 1017 cm-3 and 4 × 1017 cm-3, they obtained specific

contact resistivities in the range of ρs

=1.19 × 10-7 Ω cm2 and

8.9 × 10-8 W cm2, respectively. Calculation of the contact

resistivity was based on the assumption that the semiconductor sheet resistance underneath the contacts remains un-
32 Naval Research Reviews

changed, which is not true for nonalloyed contacts. As for

the current conduction mechanism in these ohmic contacts,

the large metal-semiconductor barriers diminish the possibility of thermionic-emission-governed ohmic contacts to

GaN. The alternative mechanism is some form of tunneling

that may take place if GaN is so heavily doped as to cause a

very thin depletion region. Tunneling is possible if, due to

annealing (for example, at 900 °C for 30 s), Al and Ti along

with Ni undergo substantial interaction with each other and

GaN. Investigations showed that Ti receives N from GaN,

forming a metallic layer, while the lack of N on GaN

provides the desired benefit of increased electron concentration through N vacancy formation [48]. Aluminum passivates the surface and also possibly reacts with Ti to form

TiAl. Reference 2 provides further details.

AlGaN/GaN MODFETs

To reiterate, MODFET’s performance is due to the conduction channel that allows large sheet carrier concentrations to

be maintained and its unique capacitance-voltage relationship [49,50]. Moreover, spatial separation of scattering

centers (such as ionized donors) from the electrons leads to

low field transport void of ionized impurity scattering. What

is somewhat unique to GaN and its alloys is its spontaneous

polarization and the strain-induced piezoelectric effect

[21,22,27,51-53] that causes redistribution of mobile and

weakly bound charge and charge collected from metal

contacts.

Stengel et al. [30] have recently modeled electronic properties of modulation-doped structures based on the III-nitride

semiconductor system. The structure considered was a

wurtzitic AlxGa1-xN/GaN normal MODFET. Because of the

conduction band discontinuity, the electrons diffusing from

the larger bandgap AlGaN into the smaller bandgap GaN

form a triangular quantum well at the AlxGa1-xN/GaN

interface, which is the hallmark of MODFETs. The source

and drain contacts were assumed to penetrate down to the

GaN layer that hosts the 2DEG.

To illustrate this, Fig. 6 shows the energy band diagram for a

normally on (N-ON) MODFET calculated by Stengel et al.

[30] for an AlN mole fraction of x = 0.25, a donor concentration in AlGaN, Nd = 1019 cm-3, and an undoped AlGaN

layer thickness (spacer layer), Wsp = 20 Å. Also shown is the

electron gas concentration at the heterointerface. For the

simulated N-ON MODFET, a gate bias of VG = 0.04 V, and

doped AlGaN layer thickness of d = 200 Å were used. The

2DEG does not extend to the AlGaN region because of a

high AlxGa1-xN/GaN conduction band discontinuity (more

than 500 meV as compared to 142 meV for AlxGa1-xAs/

GaAs at x = 0.3). Because of this, and the fact that the

probability of electron wave functions extending to the

AlxGa1-xN is very low, a thinner spacer would be sufficient

to achieve the optimal mobility in the 2DEG due to a lower

alloy scattering [54]. For the MODFET shown in Fig. 6, the

quasi-Fermi level in GaN is far above the lowest energy

level, and the peak volume concentration of electrons in the

2DEG is 1019 cm-3. Also, some of the donor atoms (for z

between -100 Å and -50 Å) are now neutralized, and some

electrons start to appear in the AlxGa1-xN region. Because of

these, a further rise of the gate bias causes not only an

increase in the donor neutralization, but also an increase in

the electron concentration in AlxGa1-xN. However, the

2DEG concentration remains unaltered.

Calculations employing typical parameters [30] for GaN and

AlGaN indicate that the peak value of 2DEG concentration

for AlxGa1-xN /GaN MODFETs is around 2-5 × 1012 cm-2.

However, much larger values of n2D have been measured,

which is most likely due to further ionization and redistribution of shallow charge by the piezoelectric effect that may

also call for charge collection from metal contacts. Neglecting the piezoelectric effect, Fig. 7 shows n2D vs VG plots for

various spacer thicknesses. As long as the unintentional

doping level in the spacer layer is low, the effects of varying

the AlN mole fraction, x, and Wsp on n2D are essentially

equivalent, because at the end of the spacer layer, the value

of Ec-EF is very close to Ec(interface)-Wsp (dEc/dz) (interFig. 6 - Band diagram for a normally on MODFET. The

origin of energy for these band diagrams is the Fermi level.

The left side of the z = 0 line corresponds to the AlGaN

region and the right side to the GaN. The donor level in

AlGaN is represented by Ed , and the quantum energy

levels in GaN are represented by E0 and E1. The term Ec

represents the conduction band edge in AlGaN and GaN

layers.

1200

1000

800

11019

0

600

400

200

0

–200

–200 –150 –100

z(Ang)

Energy (meV)

n3D (cm–3)

–50 0 50 100 150 200

E1

E0

Ed

EC

n3D
Volume 51, Number 1 33

face). Experimental data describing the effect of the spacer

layer thickness on the mobility of electrons in the 2DEG are

needed in order to evaluate optimized values for this spacer

layer. However, one may predict that these values for the

AlxGa1-xN /GaN system would be smaller than those for the

AlxGa1-x As/GaAs system, because the AlxGa1-xN /GaN

system provides a deeper confinement.

Stengel et al. [30] calculated the variation of transconductance with gate bias for various channel lengths as shown in

Fig. 8. When the channel length was decreased from 1 to

0.2 mm, the peak transconductance increased from 420 to

above 900 mS/mm. These very large transconductance

values are much higher than those experimentally observed.

They should change with improvements in layer quality and

are made possible when the values of both VG and VD are

chosen to be small to avert the velocity saturation of the

carriers. This is indeed very encouraging considering that,

for all practical purposes, MODFETs with reduced channel

lengths are very desirable and have lower leakage currents at

the operating point. Notably, the peak transconductance of

these MODFETs is obtained for lower gate biases.

Experimental Performance of GaN MODFETs

Initial GaN MODFETs utilized the background donors in

the AlGaN layer, the density of which is not controllable, to

say the least, and any other free and weakly bound electrons

drawn to the interface. Congruent with the early stages of

development and the defect-laden nature of the early GaN

and AlGaN layers, the MODFETs exhibited very low

transconductances and a low-resistance and a high-resistance state before and after the application of a high drain

voltage (20 V). As in the case of GaAs/AlGaAs MODFETs,

hot electron trapping in the larger bandgap material at the

drain side of the gate is primarily responsible for the current

collapse. The negative electron charge accumulated because

of this trapping causes a significant depletion of the channel

layer, more probably a pinch-off, leading to a drastic

reduction of the channel conductance and the decrease of

the drain current. This continues to be effective until the

drain-source bias is substantially increased, leading to a

space-charge injection and giving rise to an increased drainsource current.

With improvements in the materials quality available, the

transconductance, current capacity, and drain breakdown

voltage are all increased to the point that GaN-based

MODFETs are now strong contenders in the arena of high

power devices/amplifiers, particularly at X band and higher

frequencies. As is the case for FET device structure,

improved and high resistivity buffer layers have once again

played a pivotal role. For chronological purposes, a brief

review of the latest class of MODFETs with high

transconductances and current levels is given later in this

article.

MODFETs with a gate length of 2 μm, gate width of 40 μm,

and the drain-source separation of 4 μm exhibited drain

currents of approximately 500 mA/mm and extrinsic

Fig. 7 - Plots of the 2DEG concentration as a

function of the gate-source bias VG for various values

of the spacer layer thickness in AlxGa1-xN.

4

3.5

3

2.5

2

1.5

1

0.5

–1 0 1

VG(V)

n2D(1012 cm–2)

2 3

0

Wsp = 0 Ang

Wsp = 10 Ang

Wsp = 20 Ang

Wsp = 40 Ang

Fig. 8 - Variation of the transconductance gm of MODFETs

as a function of the gate-source bias VG, at the optimal

value of the drain-source bias VD . The optimum value of VD

is defined to be the value at which the transconductance

peak reaches its maximum value. The parameter is the

channel length, L, with values of 0.2, 0.5, 1, and 3 mm. The

parameters used for the 2DEG are d = 130 Å, Nd = 1019

cm-3,Wsp = 20 Å, x = 0.25, and Ed = 45 meV.

–2

400

300

200

100

0

–1 0

VG(V)

gm(mS/mm)

123

L = 0.2 μm

L = 0.5 μm

L = 1.0 μm

L = 3.0 μm
34 Naval Research Reviews

transconductances of approximately gem = 185 mS/mm. The

drain breakdown voltage for 1 μm gate to drain spacing was

approximately 100 V, the exact value depending on the layer

design and quality of the layered structure. Although terms

such as piezoelectric doping are being used rather liberally

to increase the sheet carrier concentration, ultimately

regardless of the source of the carriers, the strength of the

electric field that can be tolerated under the gate without

excessive leakage will set an upper limit on the number of

carriers that can accumulate at the interface and be modulated. Use of multi-2DEG structures is one obvious method

to increase the current capability of MODFETs, and they

have been employed. In those cases, the GaN layer is

straddled by two doped AlGaN molecules that donate

electrons to the channel, thus increasing the number of

electrons available for current conduction. By Hall effect

measurement, the mobility and sheet carrier densities in the

2DEG were about 304 cm2/Vs and 3.7 × 1013 cm-2, respectively, at room temperature. The sheet carrier concentration

may have been affected by piezoelectric effect. A number of

double heterochannel MODFETs (DHCMODFETs) with

gate lengths of 1.5 to 1.75 μm and a gate width of 40 μm

have been reported.

The maximum drain saturation current IDS corresponding to

a drain-source voltage VDS = 7 V, and gate source voltage

VGS = 3.5 V in a DHCMODFET is about 1100 mA/mm,

which is important because in high power devices, the input

is momentarily forward biased. The DHCMODFET has a

room temperature extrinsic transconductance of gm = 270

mS/mm. The value of the total resistance RT extracted from

the linear region of the I-V curves is 4 Ω/mm. Near pinchoff, the drain breakdown voltage is about 80 V, indicating

excellent power potential of the device. These measurements

were made in a nitrogen- pressurized container to avoid

possible oxidation of the contacts and probes. The maximum drain-source current and extrinsic transconductance of

the DHCMODFET are 500 mA/mm and 120 mS/mm,

respectively. These devices maintain reasonable output

characteristics at temperatures as high as 500 °C with

maximum drain current and extrinsic transconductance

values of 380 mA/mm and 70 mS/mm, respectively. Cooling

to room temperature restored the characteristics, which

demonstrates the robustness of this material system and of

the metallization employed. It should be noted, however,

that high power operation requires large drain breakdown

voltages with the added benefit of having large output

resistances, which ameliorates impedance matching.

MODFETs have progressed to a point where microwave

measurements have been performed on a variety of devices

with gate lengths as wide as 2 μm and as narrow as about

0.2 μm. A typical MODFET structure with 2 μm gate

lengths has been tested for small-signal S-parameters

performed at bias conditions used for the power measurements (i.e., 15 V, -2.5 V, and 20 mA for the drain voltage,

gate voltage, and drain current, respectively). The unity

current gain cutoff frequency (f

t

) and maximum frequency

of oscillation (f

max) were 6 GHz and 11 GHz, respectively, at

both 15 and 30 V bias. Values of f

t

 and f

max in excess of 50

GHz and 100 GHz have been reported for short channel

(about 0.2 μm) devices, respectively. The power measurements for 2 μm devices were taken at 4 GHz with the input

power swept from 5 dBm to 18 dBm in 14 steps. The input

and output matches, which were used during the power

sweep, were determined by iterating between source and

load pulls. The output match was selected to optimize the

output power, and the input match was selected to maximize

the delivered power. The devices were biased to VDS = 15 V

and VGS = -2.5 V. The IDS at this bias was approximately 20

mA, which corresponds to 260 mA/mm. Devices exhibited 6

dB gain for various input levels. The maximum output

power was 20.6 dBm, and the peak power added efficiency

(PAE) was 17.5%. This corresponds to a normalized output

power density of 1.5 W/mm. In general, devices on sapphire

substrates suffer from the low thermal conductivity of

sapphire substrates and exhibit negative differential resistance in the output characteristics. Remedies include better

heat sinking by flip-chip mounting and the use of high

resistivity 4H-SiC substrates, which provide good thermal

conductivity but are hard to obtain.

GaN MODFET devices that have been grown in the author’s

laboratory on conducting 6H-SiC substrates exhibited

output characteristics that lacked the negative resistance

(i.e., they exhibited good heat sinking). There have subsequently been a few reports of MODFET power devices on

high resistivity SiC [16,19,55] and p-type SiC [18]substrates with phenomenal improvement in power handling

capability notwithstanding the rapid progress on sapphire

substrates. On sapphire, recent 0.7-μm gate-length

Al0.5Ga0.5N/ GaN MODFETs exhibited a current density of

1 A/mm, three-terminal breakdown voltages up to 200 V,

and CW power densities of 2.84 and 2.57 W/mm at 8 and

10 GHz, respectively, representing a marked performance

improvement for GaN-based FETs.

Unprecedented power levels are being achieved with nearhalf-micron gate lengths. With 0.7-μm gate length devices

on SiC substrates, where the gate-source spacing and gatedrain spacing were 0.5 and 0.8 μm, respectively, a total

output power of 2.3 W in a device with a 1.28 mm gate

periphery has been obtained [19]. The power gain at the

2.3 W output power point was 3.6 dB with a PAE of 13.3%.
Volume 51, Number 1 35

The power measurement was conducted at a drain voltage of

33 V. The current and power gain cutoff frequencies were 15

and 42 GHz. The contact resistance, though not the best,

was between 2.6 and 3.5 Ωmm. The maximum normalized

trans-conductance was 270 mS/mm, and the drain current

was 293 mA/mm.

In power devices, the thermal limitation can never be

eliminated completely as is the case in nitride devices,

particularly when fabricated on sapphire substrates with a

thermal conductivity of only approximately 0.3 W/cmK.

Inclusion of thermal limitation leads to the results shown in

Fig. 9 for devices that compete in the high-power device

arena [15,31]. Since new device developments do in general

compete with existing and alternative technologies, a brief

account of competing technologies for power arena is given

below. The Si metal semiconductor FET (MESFET)

analytical curve, modeled for its simplicity, is slightly above

the SiC analytical curve and indicates a maximum power

density of 0.35 W/mm at VdS = 7V, which is slightly lower

than 0.39 W/mm. Since Si RF MESFETs are unavailable,

commercial Si RF metal-oxide semiconductor FET

(MOSFET) results were used for comparison instead. At

low voltages, the Si MOSFET data parallel the analytical

curve suggesting the validity of the functional dependence

of power density on drain voltage. Also shown are two

higher power density data points 0.4 W/mm, VdS = 28V and

0.87 W/mm, VdS = 48 V. These higher power densities were

obtained with specially designed RF power MOSFETs that

incorporate a lightly doped drain and field plates that

significantly increase the breakdown voltage.

The GaAs analytical curve shows the highest power density

of all of the devices at the lowest voltages primarily because

of the higher electron mobility of GaAs. However, the low

breakdown field limits the GaAs MESFET’s drain voltage to

about 8 V and power density to 0.63 W/mm including thermal

effects. Typical commercially available GaAs MESFET power

densities are below 1 W/mm. However, high-performance

GaAs FETs with more complex device cross sections have

achieved power densities as high as 1.4 W/mm at 18 V. At 100

V, the SiC MESFET has calculated maximum power densities

of 7.96 W/mm with thermal effects and 9.7 W/mm without

thermal effects. The highest demonstrated CW power density

3.3 W/mm (Vds = 50 V) for a SiC MESFET [56] is also shown

for comparison. Additional SiC data again illustrate the

functional dependence of power density on drain voltage. The

GaN analytical results are highly dependent on the thermal

conductivity of the substrate. With a sapphire substrate, the

device is severely thermally limited to 2.24 W/mm at 30 V

with a resulting channel temperature of over 400 °C. With a

SiC substrate, however, the analysis predicts that a GaN

MODFET could achieve 15.5 W/mm at 100 V with a channel

temperature of about 300 °C. To date, the highest power

density achieved for a 0.45 μm × 125 μm GaN MODFET is

6.8 W/mm at 10 GHz with a PAE of 52.4% and an associated

gain of 10.65 dB. At 16 GHz, the same figures in the same

order are 4.4 W/mm, 27.3%, and 6.9 dB [16]. We should

caution that while the power density figure can be used during

the evolution process, eventually the total power figure must

prevail. Normalized power density measurements, though

frequently reported (a trap the present author also fell into),

are often misleading because smaller gate widths naturally

lead to larger power densities [16]. This experimental datum

point is actually slightly higher than the simulated result,

possibly because of the very small size of the experimental

device (100 μm width).

The GaN results of analytical models are highly dependent on

the thermal conductivity of the substrate. With a sapphire

substrate, the device is severely thermally limited to

2.24 W/mm at 30 V with a resulting channel temperature over

400 °C. However with a SiC substrate, the analysis predicts

that a GaN MODFET could achieve 15.5 W/mm at 100 V

while keeping the channel temperature at about

300 °C [31]. The key to further improvements lies with our

ability to control the polarity of the films, to prepare domain

free material, and to reduce defects. If the past few years are

any indication, substantial progress is in the wings.

Ultraviolet Detectors

The alloy AlGaN with about 50-60 % AlN absorbs in the

region of the spectrum where the Sun’s radiation is absorbed

by the Earth’s ozone layer, thereby providing a dark background. A detector operating in this region sees a dark Fig. 9 - Simulated and experimental RF power density

data for Si, GaAs, SiC, and GaN FETs [31].

1 10 100

Drain Voltage (V)

RF Power Density (W/mm)

0.01

0.1

1.0

10

LDMOSFET

SiC SIT 

Si

GaN MODFET/Al2O3

Model 

Model 

Model 

Data

Data

Data

GaN MODFET/SiC

GaAs MESFET

4H-SiC MESFET
36 Naval Research Reviews

background that results in reduced radiation noise and

allows the detection of minute emission from certain threats.

UV (solar blind) detectors are imperative for the military in

that they will enable airborne, sea borne, and ground

equipment to detect and warn against ground-to-air, air-toair, air-to-ground, and ground-to-ground missile threats. The

system, at the heart of which is a UV detector, must be able

to detect and track extremely weak signals from rapidly

moving threats, which necessitates operation in the solarblind region of the spectrum, 260 to 290 nm.

The continuing proliferation and increasing lethality of

surface-to-air and air-to-air missiles pose serious threats to

Air Force aircraft. Countering these threats successfully

requires early detection of the missile launch. One promising method of detecting incoming missiles is sensing the

UV emissions from the rocket plume. Particularly at

altitudes lower than 20,000 ft, the solar-blind region is ideal

for this purpose because the solar background radiation is

almost entirely absorbed by atmospheric ozone. Thus, the

~3000 K blackbody emission from a missile plume stands

out prominently against a dark background, thereby making

false alarm rates very low.

Such a missile threat warning system, therefore, requires

two-dimensional arrays of highly sensitive photodetectors to

image the solar-blind UV spectral range. The current

solution to this problem uses photomultiplier-type detectors

and optical filters to block out all light except that between

260 and 290 nm. The problems with this approach are the

limited sensitivity due to low transmission through the

filters and low quantum efficiency of the photocathodes; the

detection of wavelengths outside the solar-blind region

caused by filter roll-off at wavelengths longer than 290 nm;

and the large size and weight of the photomultiplier assembly and associated high-voltage power supply. By comparison, solid state detectors offer the advantages of being

compact and rugged, of having near-unity quantum efficiency, and of having better rejection of long wavelength

radiation due to their sharp UV band edge. Currently,

however, no solid state UV photodetectors have the sensitivity required for this application.

Numerous papers concerning GaN and AlxGa1-xN UV

detectors have appeared in the literature. The detector

structures reported include photoconductors (PCs), metal

semiconductor-metals (MSMs), and photovoltaics (PVs,

including Schottkys, p-n junctions, and p-i-ns). These

detector structures were typically grown on sapphire

substrates by variations of MBE and metal-organic chemical

vapor deposition (MOCVD). Improvements in III-nitride

materials growth and detector processing have resulted in

progressively higher responsivity PV and PC ultraviolet

detectors in the last few years [7,8,57-63]. In the past two

years, AlxGa1-xN/GaN-based detector response times have

decreased dramatically from the ms to the ns regime.

Ultraviolet detectors of the PV [8,61] and PC [58,64]

varieties and based on nitride semiconductors prepared by

MBE have been reported. Ultraviolet detectors fabricated in

all MBE-grown materials have shown good quantum

efficiencies, noise equivalent power, and speed of response.

We only make brief mention of photoconducting detectors

in this article because they suffer from large dark-current

and memory effects. The long delay in turn-off response has

been attributed by many to majority carrier traps that are

relatively shallow, small electron capture cross-section

coupled with large hole capture rate [63,65] and unique

potential barriers [66]. Munoz et al. [66] argued that the

high gain in PC detectors is caused by a modulation

mechanism of the conductive volume in the layer. Because

carriers are photo-generated, they are spatially separated by

potential barriers generated by band-bending associated

with surface and bulk dislocations. Carrier recombination

and capture are controlled by such potential barriers and an

intrinsic, nonexponential recovery process, leading to long

delay times and large gains.

Surface-illuminated GaN and AlGaN/GaN p-i-n UV photodetectors prepared by reactive molecular beam epitaxy

(RMBE) on sapphire substrates have exhibited excellent

performance. The p-i-n structures are generally grown on Cplane sapphire substrates by plasma-enhanced MBE or

RMBE. Figure 10 shows the typical structure along with a

schematic cross sectional diagram of the device structure. In

the configuration shown, light is coupled from the top. The

thickness of the GaN layer can be adjusted for the light to be

absorbed entirely in the top p-type GaN layer or the lightly

and unintentionally n-doped GaN layer (referred to as the ilayer) or both. AlGaN (AlGaN homojunction), GaN (GaN

homojunction), and GaN with AlGaN window varieties have

been explored. The GaN homojunction structure reported in

Fig. 10 - Cross-sectional diagram of a typical

p-n junction UV detector.

p+

nn+

Sapphire

hν

contact

contact contact

contact
Volume 51, Number 1 37

Ref. 8 consists of the following layers in growth order: 1) a

thin AlN buffer layer, 2) a 3-μm Si-doped GaN n-layer, 3) a

1-μm GaN i-layer, 4) a 0.2-μm Mg-doped GaN p-layer, and

5) a 100 Å AlGaN cap layer. The AlGaN(p)-GaN(i)-GaN(n)

(AlGaN/GaN ) heterojunction structure consists of the same

set of layers as above except that the 0.2-μm p-GaN layer

was replaced with a 0.2-μm p-Al0.12Ga0.88N layer (bandgap

~3.7 eV). Thus, the p-layer is transparent to light at the band

edge of the GaN i-layer. In the device fabrication that

followed, the samples were etched down to the n-layer,

forming 250-μm diameter mesas by a reactive ion etching

(RIE) process. The n-type contact was made by depositing

Ti/Al/Ti/Au multilayers and alloying at 900 oC for 30 s

using a rapid thermal annealing furnace. The Ni/Au ring

contacts to the top p-layer were then deposited and annealed

in a tube furnace at 650 oC for 90 s in nitrogen.

 Figure 11 shows the typical I-V characteristic curves of

both the GaN and AlGaN/GaN p-i-n devices. It should be

noted that the reverse current near-zero bias is limited by the

measurement setup. Preliminary measurements indicate the

actual current to be several orders of magnitude smaller. The

homojunction has a turn-on voltage of 4 V and a reverse

breakdown voltage of ~ -100 V, while the heterojunction has

turn-on and breakdown voltages of 5.7 V and -50 V, respectively. The increased leakage current in the AlGaN/GaN

p-i-n device is attributed to the lower quality of the pAlGaN layer as opposed to that of the p-GaN layer.

Figure 12 shows the spectral response of a GaN UV detector

with an Al0.12Ga0.88N window layer as measured by top

illumination at normal incidence for zero bias and -10 V

bias. At 364 nm, the spectral bandwidth of the monochromator was approximately 1 nm, and the light intensity on the

detectors was approximately 0.06 mW/cm2. A mechanical

chopper modulated the incident light, and a lock-in amplifier recorded the photo current from the p-i-n photo diode. A

calibrated Si photodetector was used to measure the spectral

intensity of the light source and convert the measured

detector current to absolute responsivity. As can be seen, the

responsivity drops by more than three orders of magnitude as

the incident light’s wavelength increases from 360 nm to 390

nm. Moreover, reverse bias enhances the response

(0.15 A/W at -10 V) while maintaining the sharp cutoff edge

and low noise characteristics of the zero-bias responsivity.

Under -10 V bias and at 364 nm, the external quantum

efficiency is ~51% for the AlGaN/GaN detector and the

internal quantum efficiency is estimated to be over 80%

considering the surface and interface reflection.

The rms noise current at 400 nm is approximately 1 pA

(limited by the measurement setup), corresponding to an NEP

of approximately 8.3 pW. The AlGaN/GaN p-i-n device also

exhibits higher responsivity in the 260 to 300 nm spectral

region of importance. To a first approximation, we observe

that the curved shape of the responsivity of the AlGaN/GaN

Fig. 11 - Typical I-V characteristics of GaN and AlGaN/

GaN p-i-n devices.

10–12

10–7

10–8

10–9

10–10

10–11

0

10

20

30

40

50

-30 -25 -20 -15 -10 -5 0 5 10

GaN(p)/GaN(i)

AlGaN(p)/GaN(i)

Reverse Bias (A)

Bias Voltage (V)

Forward Bias (mA)

Fig. 12 - Spectral responsivity for an AlGaN-p/GaN p-i-n

photodetector under zero bias (0 V) and -10 V reverse bias.

.

240 280 320 360 400

Responsivity (A/W)

Wavelength (nm)

–10 V

0 V

100

10–5

10–4

10–3

10–2

10–1

p-i-n in the 260 to 300 nm spectral region is similar to that of

the GaN p-i-n but blue-shifted to the AlGaN bandgap. When

the defect concentrations in layers are reduced, biased

operation and designs that allow absorption in the i-layer

would increase the sensitivity and the speed of response

somewhat. The AlGaN p-layer is much desired since it is

transparent to the light at the GaN bandgap. Ideally what this

means is that photons with energies below that of the AlGaN

bandgap would be absorbed in the i-layer, where they would

generate carriers, that in turn, would be swept by the electric

field. For photons with energy greater than that of the

bandgap of the p-layer, excessive absorption in the top layer

is still a problem that is exacerbated by the small diffusion

coefficient (about 0.1 μm).
38 Naval Research Reviews

An equally important figure of merit for a photodetector is its

speed. The time response for the devices incorporating

several design features such as all GaN, all AlGaN, and

heterojunction types at zero bias were measured at 355 nm

using a pulsed-nitrogen laser and by monitoring the decay of

the photocurrent. The exponential decay times of the photocurrent of the GaN, Al0.05Ga0.95N/GaN, and Al0.1Ga0.9N

p-i-n were found to be about 29, 22, and 12 ns, respectively,

as shown in Fig. 13. Fall times as low as 6 ns have been

measured [67]. Generally, detectors exhibit two regimes in

the photocurrent time decay. At high excitation intensities,

trap states in the GaN bandgap presumably become saturated

and the detectors are then capable of resolving the 8 ns

excitation pulse. At lower excitation intensities, a longer time

response tail dominates the decay. In this regime, the detectors have a rise time of less than 5 ns and a fall time (1/e) of

approximately 31 ns.

Noise characteristics of the three types UV detectors (GaN

homojunction, Al0.03Ga0.97N homojunction, and Al0.1Ga0.9N /

GaN heterojunction devices) measured by the group of Prof.

H. Temkin at Texas Technical University are plotted in

Fig. 14, while the noise parameters are tabulated in Table 2.

Detector noise for all three devices for small bias values was

below the noise floor of the measurement setup. Hence, the

noise characteristics were measured at a reverse bias of 28 V.

Also tabulated is the corner frequency f

c (i.e., the frequency

where the 1/f noise density is equal to the shot noise density). The NEP is then found as usual, by calculating the total

rms noise current and dividing it by the responsivity of the

device.

Despite the remarkable developments detailed above, true

solar-blind detectors with very low noise levels and signal

amplification have not yet been reported. The ternary AlGaN

with large AlN mole fractions would have to be prepared

with high quality. If p-n junction varieties are to be used,

high p-doping in such a high mole fraction AlGaN is no easy

task and none has been demonstrated. Schottky barrier

varieties may remove the requirement of p-doping. However,

if avalanche photo diodes are pursued for their much-needed

gain, as required by systems applications, the Schottky

Table 2 - Noise Data for GaN Homojunction, Al0.03Ga0.97N Homojunction,

and Al0.1Ga0.9N/GaN Heterojunction Devices

Parameter GaN Diode

5 nA

7 pA

158 Hz

1.18 × 10–11 W

2.14 × 10–14 W/Hz1/2

7 nA

7 pA

1638 Hz

9.1 × 10–12 W

1.25 × 10–14 W/Hz1/2

7 nA

7 pA

1638 Hz

2.06 × 10–11 W

2.49 × 10–14 W/Hz1/2

Id @ –28 V

Id @ –2 V

fc

NEP (total)

NEP @ 10 kHz

Notes: 1) Bandwidth was assumed 30 MHz for GaN and 50 MHz for AlGaN diodes

2) Responsivity was assumed 0.07 A/W for GaN and 0.12 A/W for AlGaN diodes

3) I/f noise power density at 1 Hz is 8 × 10–18 A2

/Hz for the Al0.1 Ga0.9 N/GaN diode, 3.8 × 10–21 A2/Hz 

 for the Al0.03Ga0.97N diode, and 1.8 × 10–22 A2

/Hz for the GaN diode

AlGaN Diode AlGaN/GaN Diode

Fig. 13 - The time decay of the photo current for GaN,

Al0.1Ga0.9N/GaN, Al0.1Ga0.9N p-i-n detectors [67].

Al0.05Ga0.95N(p)-GaN(i)-GaN(n), τ = 22 ns

Al0.1Ga0.9N p-i-n, τ = 12 ns

GaN p-i-n, τ = 29 ns

Time (ns)

Relative Response 0 10 20 30 40 50

0

1

0 200 400 600 800 1000

100

10–3

10–2

10–1

Relative Response

Time (ns)

Fig. 14 - Noise power density measured at -28 V for three

different diodes, measured by the group of Prof. H. Temkin

at Texas Tech University.

-28 V Bias

Al0.03Ga0.97N Diode

Al0.1GaN0.9/GaN Diode

GaN Diode

Sn

Id

2

s0

f =

10–18

10–20

10–21

10–22

10–23

10–24

10–25

101 104 103 102

10–19

Frequency (Hz)

Noise Power Density (A2/Hz)
Volume 51, Number 1 39

barrier varieties may not be as successful as the

p-n junction varieties have been in avalanche photo diodes

based on conventional compound semiconductors. Moreover, very little is really known about the ionization coefficients of electrons and holes in GaN and its alloys. The

calculations are frustrated by the lack of reliable data on the

properties (such as complete band structure) of these

materials. These difficulties should, however, be viewed as

challenges rather than obstacles as has been the case for just

about any advancement. A reasonable midway set of

performance goals would be to achieve near 80% to 90%

internal efficiency, 10-14 W NEP, higher than 100 MHz

frequency response, 260 to 290 nm wavelength response,

and 10-4 of peak response relative response at 300 nm for

AlGaN detectors.

Table 3 presents a no-doubt incomplete collage of various

UV detectors for comparing and appraising the state-of-theart. What can be discerned is that the recent progress is

extraordinary and that MBE-prepared detectors are very

competitive with MOCVD-fabricated detectors despite the

scant activity. The last row in Table 3 (i.e., for the PMTbased imager) indicates the possibilities of the hybrid

approach used by the Lockheed-Martin Company.

The single-pixel detector development is beginning to give

way to arrays for imaging. By combining the spectral

signature and image of the source, one can more definitively

identify the origin. For example, an 8 × 8 GaN Schottky

barrier photo diode array for UV imaging having pixel sizes

of 200 μm × 200 μm, and responsivity of 0.06 A/W has been

reported [76]. A purportedly RC time-limited speed of

response of 50 ns was reported. In another effort, a 1 × 18

Table 3 - Comparison of the Performance (Responsivity, R, Internal

Gain-Quantum Efficiency Product, G·η, NEP, and Photocurrent Decay Time, τ)

of some State-of-the-Art III-Nitride Photodetectors

Detector R (A/W)

18-300

0.035-2.8 × 10–6

125

0.3

0.18

0.14

0.11

0.15

~670

64-1500

600

0.5

0.48

0.63

~3000

57 pA at 10V

~7 × 10–12

7 × 10–14

8 × 10–12*

~10–15

AIx Ga1–xN PC

AIx Ga1–xN PC

GaN PC

GaN MSM

n-GaN Schottky PD

GaN p-n PD

AlGaN/GaN p-n PD

GaN p-i-n PD

AlGaN/GaN p-i-n PD

GaN p-i-n PD

PMT-based imager

*Limited by measurement setup. The total NEP value at a reverse bias of 28 V is 2.06 × 10–11 W.

 PC = photoconductor, MSM = Metal Semiconductor-Metal, PD = photodiode, PMT = photomultiplier.

G·η (internal)

1-2 × 10–3

0.13-0.36 × 10–3

>20 × 10–9

120 × 10–9

1.7 × 10–8

10

8.2 × 10–3

1.2 × 10–8

9 × 10–11

<10–3

NEP (W) τ (s)

68

69

70, 71

72

51

57

73

62

8

74

75

Ref.

GaN MSM linear array with a responsivity of 3250 A/W at a

bias of 10 V with a response time of 0.5 ± 0.2 ms was also

reported [77].

Acknowledgments

This work was supported by the Office of Naval Research

and the Air Force Office of Scientific Research and monitored by C.E. Wood, Y.S. Park, M. Yoder, and G. L. Witt.

The author appreciates nudging by C. E. Wood to volunteer

and finish the project, and fruitful discussions with and

assistance of M. Estes, P.J. Schreiber, F. Hamdani, G. Smith,

C. W. Litton, D.C. Reynolds, D.C. Look, H. Jiang, B. Gil,

H. Temkin, J. Schetzina, A. Salvador, B. Goldenberg, W.

Yang, and S. Krishnankutty. Last but not least, the author

thanks S.N. Mohammad and F. Stengel for extensive use of

their contributions, and S. Oresky of the Naval Research

Laboratory for editorial contributions.

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22. F. Bernardini, V. Fiorentini, and D. Vanderbilt, “Spontaneous

Polarization and Piezoelectric Constants in III-V Nitrides,”

Phys. Rev. B 56, R10024 (1997).

23. For a review see, R. Resta, “Macroscopic Polarization in

Crystalline Dielectrics: the Geometric Phase Approach,” Rev.

Mod. Phys. 66, 899 (1994) and references therein.

24. B. Daudin, J.L. Rouviere, and M. Arley, “Polarity Determination of GaN Films by Ion Channeling and Convergent Beam

Electron Diffraction,” Appl. Phys. Lett. 69, 2480 (1996).

25. P. Vermaut, P. Ruterana, G. Nouet, A. Salvador, and H.

Morkoç, “Structural Defects Due to Interface Steps and

Polytypism in III-V Semiconducting Materials: a Case Study

Using High Resolution Electron Microscopy of the 2H-AlN/

6H-SiC Interface,” Philos. Mag. A 75, 239-259 (1997).

26. V. Potin, P. Ruterana, G. Nouet, A. Salvador, and H. Morkoç,

“The Atomic Structure of the {10-10} Inversion Domains in

GaN/Sapphire Layers,” Proc. of Gallium Nitride and Related

Materials II Symposium, pp. 323-328, C.R. Abernathy, H.

Amano, and J.C. Zolper, eds., Vol. 468, pp. 323-328 (1997),

Mater. Res. Soc; Pittsburgh, PA, USA, San Francisco, CA,

USA; 1-4 April.

27. G.A. Martin, A. Botchkarev, A. Agarwal, A. Rockett, and H.

Morkoç, “Valence-band Discontinuities of Wurtzite GaN,

AlN, and InN Heterojunctions Measured by X-ray Photoemission Spectroscopy,” Appl. Phys. Lett. 68, 2541 (1996).

28. F. Bernardini and V. Fiorentini, “Macroscopic Polarization

and Band Offsets at the Nitride Heterojunctions,” Phys. Rev.

B. 57(16), 1-4 (1998).

29. T.J. Drummond, H. Morkoç, K. Lee, and M. Shur, “Model for

Modulation Doped AlxGa1-xAs/GaAs Field Effect Transistors,” IEEE Electron. Dev. Lett. EDL-3, 338-341 (1982).
Volume 51, Number 1 41

30. F. Stengel, S.N. Mohammad, and H. Morkoç, “Theoretical

Characteristics of AlGaN/GaN MODFETs,” J. Appl. Phys.

80(5), 3031-3042 (1996).

31. C. Weitzel, L. Pond, K. Moore, and M. Bhatnagar, “Effect of

Device Temperature on RF FET Power Density,” Proc. of

Silicon Carbide, III-Nitrides and Related Materials, ICSI,

August 1997, Stockholm, Sweden, Trans Publications, Ltd.,

Materials Science Forum, Vols. 264-268, pp. 907-912 (1998).

32. K. Suzue, S.N. Mohammad, Z.F. Fan, W. Kim, Ö. Aktas, A.E.

Botchkarev, and H. Morkoç, “Electrical Conduction in

Platinum-Gallium Nitride Schottky Diodes,” J. Appl. Phys.

80(6), 4467-4478 (1996).

33. S.N. Mohammad and H. Morkoç, “Progress and Prospects of

Group III-V Nitride Semiconductors,” Progress in Quantum

Electronics 20 (5 and 6), 361-525 (1996) (a monograph).

34. Ö. Aktas, W. Kim, Z. Fan, S.N. Mohammad, A. Botchkarev,

A. Salvador, B. Sverdlov, and H. Morkoç, “High Transconductance-Normally-Off GaN MODFETs,” Electron. Lett.

31(16), 1389-1390 (1995).

35. Z. Fan, S.N. Mohammad, Ö. Aktas, A. Botchkarev, A.

Salvador, and H. Morkoç, “Very Low Resistance Multi-layer

Ohmic Contact to n-GaN,” Appl. Phys. Lett. 69(9), 1229-1231

(1996).

36. S.N. Mohammad, Z.-F. Fan, A. Salvador, Ö. Aktas, A.E.

Botchkarev, W. Kim, and H. Morkoç, “Near Ideal PlatinumGaN Schottky Diodes,” Appl. Phys. Lett. 69(10), 1420-1422

(1996).

37. Z. Fan, C. Lu, A. Botchkarev, H. Tang, A. Salvador, Ö. Aktas,

W. Kim, and H. Morkoç, “AlGaN/GaN Double

Heterostructure Channel Modulation Doped Field Effect

Transistors (MODFETs),” Electron. Lett. 33, 814-815 (1997).

38. Ö. Aktas, Z. Fan, A. Botchkarev, M. Roth, T. Jenkins, L.T.

Kehias, and H. Morkoç, “Microwave Performance of AlGaN/

GaN Inverted MODFETs,” IEEE Electron. Dev. Lett. 18, 293-

295 (1997).

39. P. Hacke, T. Detchprohm, K. Hiramatsu, and N. Sawaki,

“Schottky Barrier on n-type GaN Grown by Hydride Vapor

Phase Epitaxy,” Appl. Phys. Lett. 63, 2676-2678 (1993).

40. S.C. Binari, H.B. Dietrich, G. Kelner, L.B. Roland, K.

Doverspike, and D.K. Gaskill, “Electrical Characterisation of

Ti Schottky Barriers on n-type GaN,” Electron. Lett. 30, 909-

910 (1994).

41. M.A. Khan, J.M. Van Hove, J.N. Kuznia, and D.T. Olson,

“High Electron Mobility GaN/AlxGa1-xN Heterostructures

Grown by Low-pressure Metalorganic Chemical Vapor

Deposition,” Appl. Phys. Lett. 58, 2408-2410 (1991).

42. M.R.H. Khan, H. Nakayama, T. Detchprohm, K. Hiramatsu,

and N. Sawaki, “A Study on Barrier Height of Au-AlxGa1-xN

Schottky Diodes in the Range of 0<x<0.20,” in Topical

Workshop on III-V Nitrides Proc., Nagoya, Japan, 1995, Solid

State Electronics 41(2), 259-266 (1997).

43. J.S. Foresi and T.D. Moustakas, “Metal Contacts to Gallium

Nitride,” Appl. Phys. Lett. 62, 2859-2861 (1993).

44. M.A. Khan, J.N. Kuznia, A.R. Bhattarai, and D.T. Olson,

“Metal Semiconductor Field-Effect Transistor Based on

Single-Crystal GaN,” Appl. Phys. Lett. 62, 1786-1787 (1993).

45. M.E. Lin, Ma, F.Y. Huang, Z. Fan, L. Allen, and H. Morkoç,

“Low Resistance Ohmic Contacts on Wide Band-Gap GaN,”

Appl. Phys. Lett. 64, 1003-1005 (1994).

46. Y. Wu, W. Jiang, B. Keller, S. Keller, D. Kapolnek, S.

Denbaars, and U. Mishra, “Low Resistance Ohmic Contacts to

n-GaN with a Separate Layer Method,” Proc. in Solid State

Electronics 41(2), 75-78 (1997).

47. Z. Fan, S.N. Mohammad, W. Kim, Ö. Aktas, A.E. Botchkarev,

and H. Morkoç, “Very Low Resistance Multi-layer Ohmic

Contact to n-GaN,” Appl. Phys. Lett. 68, 1672-1674 (1996).

48. S. Ruvimov, Z. Liliental-Weber, J. Washburn, K.J. Duxstad,

E.E. Haller, S.N. Mohammad, Z. Fan, and H. Morkoç,

“Microstructure of Ti/Al and Ti/Al/Ni/Au Ohmic Contacts for

n-GaN,” Appl. Phys. Lett. 69(11), 1556-1558 (1996).

49. M. Moloney, F. Ponse, and H. Morkoç, “Gate Capacitance

Voltage Characteristics of MODFETs: Its Effect on Transconductance,” IEEE Trans. Electron. Dev. ED-32(9), 1675-1684

(1985).

50. H. Morkoç, H. Ünlü, and G. Ji, Fundamentals and Technology

of MODFETs, Vol. I (John Wiley and Sons, New York, 1991).

51. L. Smith, “Strain-generated Electric Fields in (111) Growth

Axis Strained-layer Superlattices,” Solid State Communications

57, 919 (1986).

52. A. Bykhovski, B. Gelmont, and M. Shur, “The Influence of the

Strain-induced Electric Field on the Charge Distribution in

GaN-AlN-GaN Structure,” J. Appl. Phys. 74, 6734 (1993).

53. G.A. Martin, S.T. Strite, A. Botchkarev, A. Agarwal, A.

Rockett, W.R.L. Lambrecht, B. Segall, and H. Morkoç,

“Valence Band Discontinuity Between GaN and AlN Measured

by X-Ray Photemission Spectroscopy,” Appl. Phys. Lett. 65(5),

610-612 (1994); G.A. Martin, A. Botchkarev, A. Rockett, and

H. Morkoç, “Valence-band Discontinuities of Wurtzite GaN,

AlN, and InN Heterojunctions Measured by X-ray Photoemission Spectroscopy,” Appl. Phys. Lett. 68(10), 2541-2543

(1996); G.A. Martin, Ph.D. Thesis, Department of Physics,

University of Illinois (1996).

54. C. Weisbuch and B. Vinter, Quantum Semiconductor Structures

(Academic Press, San Diego, 1991).

55. S. Binari, J.M. Redwing, G. Kelner, and W. Kruppa, “AlGaN/

GaN HEMTs Grown on SiC Substrates,” Electron. Lett. 33(3),

242-243 (1997).
42 Naval Research Reviews

56. K.E. Moore, C.E. Weitzel, K.J. Nordquist, L.L. Pond, III, J.W.

Palmour, S. Allen, and C.H. Carter, Jr., IEEE Electron. Dev.

Lett. 18(2), 69-70 (1997).

57. Q. Chen, M. Asif Khan, C.J. Sun, and J.W. Yang,

“Visible-Blind Ultraviolet Photodetectors Based on GaN p-n

Junctions,” Electron. Lett. 31, 1781-1782 (1995).

58. D. Walker, X. Zhang, P. Kung, A. Saxler, S. Javadpour, J. Xu,

and M. Razeghi, “AlGaN Ultraviolet Photoconductors Grown

on Sapphire,” Appl. Phys. Lett. 68, 2100-2101 (1996).

59. K.S. Stevens, M. Kinniburgh, and R. Beresford, “Photoconductive Ultraviolet Sensor Using Mg-Doped GaN on Si(111),”

Appl. Phys. Lett. 66, 3518-3520 (1995).

60. D. Wickenden, Z. Huang, D.B. Mott, and P.K. Shu, “Development of Gallium Nitride Photoconductive Detectors,” Johns

Hopkins APL Tech. Digest 1S, 217-225 (1997).

61. Q. Chen, J.W. Yang, A. Osinsky, S. Gangopadhyay, B. Lim,

M.Z. Anwar, M. Asif Khan, D. Kuksenkov, and H. Temkin,

“Schottky Barrier Detectors on GaN for Visible-Blind

Ultraviolet Detection,” Appl. Phys. Lett. 70, 2277-2279

(1997).

62. J.M. Van Hove, R. Hickman, J.J. Klaassen, and P.P. Chow,

“Ultraviolet-Sensitive, Visible-Blind GaN Photodiodes

Fabricated by Molecular Beam Epitaxy,” Appl. Phys. Lett. 70,

2282-2284 (1997).

63. A. Osinsky, S. Gangopahyay, R. Gaska, B. Williams, M.A.

Khan, D. Kuksenkov, and H. Temkin, “Low Noise p-p-n GaN

Ultraviolet Photodetectors,” Appl. Phys. Lett. 71, 2334-2336

(1997).

64. F. Binet, J.Y. Duboz, E. Rosencher, F. Scholz, and V. Härle,

“Mechanism of Recombination in GaN Photodetectors,” Appl.

Phys. Lett. 69, 1202-1204 (1996).

65. F. Binet, J.Y. Duboz, E. Rosencher, O. Briot, and R.L.

Aulombard, “Properties of a Photovoltaic Detector Based on

an n-type GaN Schottky Barrier,” J. Appl. Phys. Lett. 81,

6449-6454 (1997).

66. E. Munoz, E. Monroy, J.A. Carrido, I. Izpura, F.J. Sanchez,

M.A. Garcia, E. Calleja, B. Beaumont, and P. Gibart,

“Photoconductor Gain Mechanism in GaN Ultraviolet

Detectors,” Appl. Phys. Lett. 71, 870-872 (1997).

67. G. Smith, M.J. Estes, D. Tang, A. Salvador, Z. Fan, G. Xu, A.

Botchkarev, and H. Morkoç, “Megahertz Bandwidth AlxGa1-

xN/GaN-Based p-i-n Detector,” presented at the SPIE

meeting, San Jose CA, January 25-28 (1998).

68. B.W. Lim, Q.C. Chen, J.Y. Yang, and M. Asif Khan, “High

Responsivity Intrinsic Photoconductors Based on AlxGa1-

xN,” Appl. Phys. Lett. 68, 3761 (1996).

69. D. Walker, X. Zhang, P. Kung, A. Saxler, S. Javadpour, J. Xu,

and M. Razeghi, “AlxGa1-xN (0≤x≤1) Ultraviolet

Photoconductors Grown on Sapphire by Metal-Organic

Chemical Vapor Deposition,” Appl. Phys. Lett. 70, 949-951

(1997).

70. M. Misra, T.D. Moustakas, R.P. Vaudo, R. Singh, and K.S.

Shah, “Photoconducting Ultraviolet Detectors Based on GaN

Films Grown by Electron Cyclotron Resonance Molecular

Beam Epitaxy,” Proc. of SPIE Conference on X-ray and UV

Sensors and Applications 2519, 78 (1995).

71. M. Misra, D. Korakakis, R. Singh, A. Sampath, and T.D.

Moustakas, “Photoconducting Properties of Ultraviolet

Detectors Based on GAN and AlxGa1-xN Grown by ECRassisted MBE,” F.A. Ponce, T.D. Moustakas, I. Akasaki, B.A.

Monemar, eds., MRS Proc. 449, 597 (1997).

72. J.C. Carrano, P.A. Grudowski, C.J. Elting, R.D. Dupuis, and

J.C. Campbell, “Very Low Dark Current Metal-Semiconductor-Metal Ultraviolet Photodetectors Fabricated on SingleCrystal GaN Epitaxial Layers,” Appl. Phys. Lett. 70, 1992-

1994 (1997).

73. Honeywell, Inc., unpublished results, private communications

with D. Zook and B. Goldenberg.

74. J.C. Carrano, T. Li, D.L. Brown, P.A. Grudowski, C.J. Eiting,

R.D. Dupuis, and J.C. Campbell, “High-speed p-i-n Ultraviolet Photodetectors Fabricated on GaN,” Electron. Lett. 34(18),

1779-1781 (1998).

75. Lockheed-Martin, PMT-based imager; private communication

with Dr. Mike Estes of Air Force Research Laboratory (1997).

76. W. Lim, S. Gangopadhyay, J.W. Yang, A. Osinsky, Q. Chen,

M.Z. Anwar, and M.A. Khan, “8×8 GaN Schottky Barrier

Photodiode Array for Visible-Blind Imaging,” Electron. Lett.

33(7), 633-634 (1997).

77. Z.C. Huang, J.C. Chen, D.B. Mott, and P.K. Shu, “High

Performance GaN Linear Array,” Electron. Lett. 32(14), 1324-

1325 (1996).
Volume 51, Number 1 43

The Author

Hadis Morkoç received his B.S.E.E and M.S.E.E.

degrees from Istanbul Technical University, Turkey, in 1968

and 1969, respectively. He received his Ph.D. degree in

electrical engineering in 1975 from Cornell University,

Ithaca, New York. He was employed at Varian Associates,

Palo Alto, California from 1976 to 1978, working on various

novel FET structures and optical emitters based on thennew semiconductor structures, and at the University of

Illinois at Urbana-Champaign from 1978 to 1997, where

he conducted research on a number of heterostructures

and devices. He held visiting positions at AT&T Bell

Laboratories (1978-1979), the California Institute of

Technology, and the Jet Propulsion Laboratory (1987-

1988). In 1997, he joined the Virginia Commonwealth

University in Richmond, Virginia. His recent research

interests have focused on wide bandgap Group III-nitrides

and devices. He has coauthored a two-volume book on

MODFETs. He has also authored a book on wide bandgap

nitrides and devices; some 30 book chapters; 40 tutorialreview-popular articles dealing with heterostructures,

electronic, and optoelectronic devices; some 850 technical

journal articles; and 100 conference papers. He holds

several patents on various FETs, including the basic

pseudomorphic MODFET patent, and serves or has

served as a consultant to some 20 major industrial

laboratories.
– S.O.

How do they DO that? How do they DO that?

?

Although their potential has been recognized since the 1960s, the Group

III-nitrides have trailed way behind the easier-to-grow Si and GaAs semiconductors on the development curve. Ideal substrates upon which to

grow the nitrides could not be found, and experiments on existing substrates yielded materials with too high concentrations of defects to afford the desired electronic properties.

Nearly 30 years later, however, Si and GaAs have been pushed to their

theoretical limits while GaN and the other wide bandgap semiconductors are just beginning to show their stuff. Why the late blooming? The

recent and startling advances in the application of the electronic capabilities of the GaN semiconductors are the direct result of several developments: greater precision in growth techniques; higher purity of chemicals

used in film deposition; the development of new techniques for selective

deposition of the impurities through which the nitrides get their remarkable electronic properties; and realization of immobile electrical impurities under high applied voltages. One of the most significant insights was

that older methods of growing Si and GaAs semiconductors could be

applied to producing these materials as well. By using new variations on

these already mature technologies, researchers are working to produce

GaN thin films on GaN substrates as well as on sapphire and silicon carbide (SiC). They are graduating a class of semiconductors whose performance is expected to drive microelectronic and optoelectronic device

capabilities well into the next century. In this article, North Carolina State

University materials science and engineering researchers brief us on recent developments in the growth of these materials and deliver an accessible tutorial on the methods, the materials science, and the challenges

remaining in producing Group III-nitride thin film semiconductors.

44 Naval Research Reviews
Photograph courtesy of Aixtron

Volume 51, Number 1 45
46 Naval Research Reviews

Kevin Linthicum, Drew Hanser, Tsvetanka Zheleva, Ok-Hyun Nam, and Robert F. Davis

Department of Materials Science and Engineering, North Carolina State University

Raleigh, North Carolina

Recent Progress in the Growth

and the Control of Defects and

Electronic Properties in

III-Nitride Thin Films

III-Nitride Material Properties

Gallium nitride (GaN), aluminum nitride (AlN), and indium

nitride (InN) form a group of wide bandgap semiconductor

materials called the Group III-nitrides. These materials have

long been recognized for their potential for optoelectronic

and microelectronic semiconductor device applications.

Substantial research on III-nitride growth was initiated in

the early 1960s; however, progress slowed because ideal

substrates could not be found and the consequential growth

of GaN thin films contained substantial concentrations of

defects and lacked the desired electronic properties. These

investigations essentially ended in the early 1970s while

research on silicon (Si) and gallium arsenide (GaAs)

materials advanced rapidly. However, the 1990s have

brought significant advances in the sophistication of growth

techniques, the purity of the chemicals used for film

deposition, the controlled introduction and activation of

selected impurities that give the desired electronic properties

to these films, and the immobility of the aforementioned

defects under an applied electrical voltage. These advances

have allowed rapid progress to be made in the III-nitrides.

The III-nitrides possess considerable physical hardness,

good thermal conductivity, and high melting temperatures.

Ternary alloys of selected fractions of GaN, AlN or InN, and

device structures made from layers of these compounds and

their alloys can be produced to emit light in colors that span

virtually the entire visible spectrum and expand into the

ultraviolet (UV) spectra. Green, blue, and UV light emitting

devices, UV light detectors, and microelectronic devices for

high temperature, high power, and high frequency are

representative applications. Recently, III-nitride light

emitting diodes (LEDs) have been commercially realized.

The range of green-to-UV wavelengths represents a commercially important region of the electromagnetic spectrum.

Market projections show that blue and green LEDs will

represent the majority of the estimated $3 billion per year

GaN-based device market by 2006. The combination of

GaN-based blue and green LEDs with gallium arsenidebased red LEDs already forms the basis for large scale, fullcolor video displays. GaN-based solid-state white light

sources with long lifetimes (i.e., more than 100,000 hours)

and significantly reduced power consumption requirements

offer advantages over incandescent or fluorescent light

sources. The demonstration of continuous operation of

GaN-based laser diodes (LDs) at room temperature with a

projected lifetime of 10,000 hours has been reported [1].

The use of violet or UV LDs in optical data storage applications offers the potential for a fourfold increase over current

data storage densities.

Significant progress has also been made regarding III-nitride

microelectronic devices; however, continued research is

required to achieve commercial viability. High electron

mobility transistors (HEMTs) currently comprise the bulk of

microelectronic research. These devices take advantage of

the thermal, electronic, and piezoelectronic properties of

GaN and AlGaN and perform in environments not covered

by current technology. Further advances, including novel

material fabrication techniques discussed herein, have

already increased the capabilities of select III-nitride

optoelectronic devices; these techniques will likely have the

same positive effect on the microelectronic devices.

During the next decade, the transition of III-nitride materials

and related device technologies from laboratory experiments

to commercial products will be fueled by recent advances in
Volume 51, Number 1 47

the process routes for growth of these materials and the

recent realization of commercially viable III-nitride devices.

Although III-nitrides are grown using different environmental conditions (e.g., pressures, temperatures, and chemical

precursors) than those used for the more developed silicon

and gallium arsenide material systems, the techniques used

to grow the III-nitride thin films are essentially the same.

These scientifically mature growth and deposition techniques as well as the innovative approaches material

scientists have recently employed and adapted to overcome

many of the difficult obstacles to III-nitride growth are

discussed below.

Growth of III-Nitride Thin Films

One of the challenges that faced researchers in the evolution

of III-nitride semiconductors was the development of

processes to grow the thin films needed for the different

types of devices. Single-crystal thin semiconductor films are

normally the building blocks for various devices and are

achieved via “epitaxial growth.” The word epitaxy is derived

from the Greek word epi meaning “on” and the Latin word

taxis meaning “arrangement” or “order.” Epitaxial growth

is a process through which a crystalline substance with a

particular arrangement and orientation of atoms is grown

upon a crystalline substrate with the same atomic arrangement and orientation. There are two different types of

epitaxial growth: homoepitaxy and heteroepitaxy.

Homoepitaxy refers to epitaxial growth where the thin film

and substrate are the same material. A common example is

the growth of a silicon epitaxial layer on a silicon substrate

or wafer. Heteroepitaxy refers to the epitaxial growth of

dissimilar materials. Gallium nitride growth on sapphire or

silicon carbide is an example of this type of growth.

Epitaxially grown sequences or multiple layers of thin films,

sometimes called heterostructures, form the basis of many

different types of semiconductor devices.

Many epitaxial thin film growth processes have been

developed, including molecular beam epitaxy (MBE),

hydride vapor phase epitaxy (HVPE) and metalorganic

vapor phase epitaxy (MOVPE), and derivatives of these

methods. The MOVPE technique is commonly used for the

heteroepitaxial deposition of compound semiconductors,

including III-nitride thin films. Characteristics of this

method include the use of high purity chemical sources, a

high degree of compositional control and uniformity, high

growth rates, and the ability to grow abrupt junctions.

Growth of GaN thin films using MOVPE is achieved at

temperatures ranging from 900 to 1100 °C (1655 to 2015 °F)

via the introduction of gases into a specially designed

reaction chamber at specific rates and ratios. The reaction

chamber is usually maintained at or below a pressure of one

atmosphere during the growth process. The gases decompose on the hot substrate and the resulting products of this

reaction combine to form the thin film. This chemical

reaction must be carefully controlled to obtain optimum film

properties. During growth, the deposition conditions can be

altered to manipulate the compositions and electrical

properties of the film to meet the application requirements

for the material and desired device.

Two of the more important crystal properties that must be

closely matched between the thin film and substrate are the

distances of separation among the atoms at the thin film/

substrate interface and the amount of expansion and

contraction that occurs during heating and cooling. Any

mismatch in these properties can result in imperfections and

defects in the thin film. These defects directly affect the

success of the fabrication and performance of the final

optoelectronic or microelectronic device.

The ideal candidate substrate for the growth of GaN thin

films is a GaN wafer. These wafers allow homoepitaxial

growth and provide exact matches of the atomic spacing and

of the coefficients of thermal expansion between the

substrate and thin film. Several research groups are investigating the growth of bulk GaN crystals and very thick films

through various techniques; however, commercially available large area GaN wafers appear to be several years away.

The wide bandgap semiconductor community is therefore

challenged with growth of heteroepitaxial films.

Heteroepitaxial growth of GaN films was initially performed

directly on sapphire and silicon carbide substrates. Relatively large mismatches in atom-atom separation between

GaN and those substrates made nucleation and growth of

thin films difficult. For the growth of the GaN thin film to

proceed, the gallium and nitrogen atoms near the substrate

interface must shift their positions to accommodate the

misalignment and assume the spacing of the substrate. This

shift in atomic position results in the formation of crystalline

defects that exhibit the characteristic of threading vertically

from the substrate interface through the newly deposited

GaN thin film. These “threading defects” are detrimental to

the performance of optical and microelectronic GaN-based

devices.

To reduce the influence of these mismatches in atomic

spacing, researchers now employ an intermediate “buffer”

layer which is grown on the substrate prior to the growth of

the GaN or other III-nitride film. Although the buffer layer

has reduced the effects of the atom-atom strain created by
48 Naval Research Reviews

the differences in atomic spacing, the densities of the

threading defects in these thin films are on the order of one

million times higher than in other semiconductor systems.

The dark, vertically oriented stripes visible in the GaN film

in the transmission electron microscopy (TEM) image

shown in Fig. 1 are these threading defects. Consequently,

until recently, the film quality relative to other semiconductors remained poor. In 1997, a method to further improve the

quality of the heteroepitaxially grown GaN by a marked

reduction in defect density was employed for GaN films, as

discussed in Ref. 2 and in the following paragraphs.

Fig. 1 - Threading defects are seen on this TEM image of

GaN grown using an AIN buffer layer on SiC.

Analysis of TEM micrographs of GaN material deposited at

NCSU using a technique known as selective area growth

revealed that certain regions of the film contained a much

lower dislocation density (less than 104 cm-2) relative to

other regions where the dislocation density was approximately 109 cm-2 (typical of levels found in most GaN films).

Selective area growth is not new to the semiconductor

industry, having been used often for the growth of silicon

and gallium arsenide thin films. In the case of GaN, one

approach for this technique employs covering a layer of

gallium nitride with a layer (“mask”) of silicon dioxide

(glass). Small circular or rectangular holes or “windows” are

then etched in the mask through which the underlying GaN

is exposed. A GaN film is subsequently regrown under

conditions such that growth occurs epitaxially only in the

windows and not on the silicon oxide mask. Figure 2 shows

a scanning electron micrograph (SEM) of the resulting

growth within circular holes. These are hexagonal pyramids

that have an ideal shape for a microelectronic device known

as a field emitter.

Fig. 2 - SEM image of GaN hexagonal pyramid array.

If the regrowth of GaN is allowed to continue, eventually

lateral growth over the mask occurs. This is shown schematically in Fig. 3; a TEM image of typical LEO growth

(defined below) is shown in Fig. 4. Analysis of these films

revealed that the threading defects were only present in the

window regions of the regrown film. Very few threading

dislocations were visible in the regrown GaN that extended

laterally over the mask.

By using a selected orientation of the rectangular windows,

growth temperature, and gallium precursor flow rate, a

continuous, coalesced layer of GaN (shown schematically in

Fig. 5) containing a drastic reduction in threading dislocations compared to conventional growth can be obtained, as

seen in Fig. 6. This technique of selective area growth is

Fig. 3 - Schematic of selective growth.

200 nm

AIN

SiC

GaN
Volume 51, Number 1 49

known as lateral epitaxial overgrowth (LEO). Figure 7 is a

color cathodoluminescent image of uncoalesced LEO. The

window regions show a yellow-green color due to emission

from defects. The overgrown regions show blue color due to

an absence of defects. Shortly after these and other observations regarding defect reduction in LEO material were

reported, Nichia Chemicals employed the LEO growth

technique for their blue laser diodes and announced an

increase in device lifetime from a few hundred hours to the

estimated and previously noted 10,000 hours [1]. As a

result, the LEO technique has received considerable interest

from the III-nitride research community.

Fig. 4 - Cross-sectional TEM of typical lateral epitaxial

overgrowth.

Fig. 5 - Schematic of coalesced lateral epitaxial overgrowth.

Although the advantage of the LEO growth technique is

reduced defect densities, the area of low defect density GaN

is confined to strips in the coalesced films over the mask

regions. The film over the window regions still contains an

unacceptably high level of the threading defects. The result

is the disadvantage that the fabrication of semiconductor

devices is confined to select areas on the surface of the film.

One possible solution to this problem is a reiteration of the

LEO process with the second layer

translated. Figure 8 is an image of this

process obtained using scanning

electron microscopy.

In addition to finding new methods to

reduce the defect densities stemming

from heteroepitaxial growth on

sapphire and silicon carbide, researchers have recently renewed their

efforts to obtain device-quality, large

area GaN films on silicon substrates

because of their larger area and lower

cost. The disadvantage of using

silicon for GaN growth is that the atomic spacing mismatch

between GaN and silicon is much greater than that between

GaN and sapphire or silicon carbide. Additionally, silicon is

not as thermally stable as GaN, sapphire, or silicon carbide,

and thus any device application for GaN using silicon as a

substrate has limited use. A possible solution meeting all

these concerns is to use the LEO technique to grow devicequality GaN on large-area silicon substrates and then

remove the silicon substrate from the GaN using established

silicon processing techniques.

Fig. 6 - Example of coalesced lateral epitaxial overgrowth.

Introducing Impurities to Control

Electronic Properties

In addition to growing GaN as well as other semiconductor

films with low defect densities, another key requirement for

fabricating optoelectronic and microelectronic devices is the

ability to precisely control the electrical properties of the

thin films. By “using” and “controlling” electrons, we can

“communicate” with semiconducting materials and exploit

their unique crystalline responses to different “stimuli” (e.g.,

light, pressure, heat, and electrical voltages and currents).

Most materials can be broadly classified using the modern

band theory of solids as metals, semiconductors, or insula-
50 Naval Research Reviews

tors. In this theory, the bands consist of electrons that share

similar energy levels. In metals, electrons move in the

partially filled valence (bonding) band upon the application

of a voltage potential and are the charged entities by which

an electrical current is achieved and measured. This is

referred to as current conduction. In insulators and semiconductors, the valence band is completely filled and, in order

for conduction to occur, enough energy must be provided to

cause the electrons to jump into the conduction band. The

size of the energy gap separating the bands distinguishes

insulators from semiconductors: Insulators exhibit very wide

energy gaps, whereas semiconductors have energy gaps that

span from narrow to wide.

Another form of electrical conduction is found in materials.

When an electron moves from the valence band, it leaves

behind a vacancy. An electron from an adjoining atom can

fill that vacancy, leaving behind a vacancy in the adjoining

atom. This vacancy is called a hole and acts as another

entity that undergoes electrical conduction. However, unlike

electrons, holes do not actually carry a charge.

A semiconductor with electrons as the main current conduction mechanism is called an “n-type” semiconductor. When

a large number of holes are created, holes become the main

conduction mechanism in the semiconductor. In this case,

Fig. 8 - SEM image of double layer LEO.

the semiconductor is called “p-type.” The introduction of certain elements into a semiconductor

leads to current conduction through either electrons or holes. The careful introduction of these

elements into the semiconductor during growth is

called doping, and allows for control of both the

conduction mechanism and conduction level.

Many key semiconductor devices, including LEDs

and LDs, are based on the positioning of adjacent

n- and p-type semiconductor layers, p-n junctions,

and the resultant electrical properties of the

interface between the layers. By controlling the

interaction of electrons and holes in the semiconductor, either optoelectronic or microelectronic

phenomena result. By controlling the flow of current

conduction (by holes, electrons, or both), microelectronic

devices such as transistors are obtained. By controlling

either the recombination of holes and electrons resulting in

light emission or the absorption of light resulting in the

production of electrons and holes, optoelectronic devices

such as LEDs and light detectors, respectively, are obtained.

The current level of progress in the development of GaN

materials and commercially viable devices, namely GaNbased LEDs and LDs, has been the direct result of the

relatively recent achievement of p-type conduction in GaN.

Efficient n-type doping of GaN through the use of silicon

proved relatively easy to achieve. Doping with magnesium

is the most common method of achieving p-type conduction

in GaN. Initially, MOVPE-grown GaN doped with Mg was

not p-type, but was highly nonconductive. It was later

determined [3] that an interaction between hydrogen and

Mg occurred during the growth of the GaN thin films, which

created a neutral complex that prevented the formation of

holes in GaN. Hydrogen is a gas that is typically present in

large quantities in the MOVPE chamber during growth of

GaN. Low resistivity, p-type Mg-doped GaN was first

achieved via post-growth, low-energy electron-beam

irradiation, and later through post-growth thermal annealing.

These processes dissociated the magnesium and hydrogen in

the GaN and allowed p-type conduction in the semiconductor.

Summary

Recognized for their stability in harsh environments and

their unique electronic properties, the III-nitrides have

become one of the few select classes of materials being

engineered into optoelectronic and microelectronic devices

demanded by today’s technologically advancing society.

Although early progress was slowed by the lack of ideal

substrates and p-type doping ability, the pace of successful

Fig. 7 - Color CL image of lateral epitaxial overgrowth of GaN.
Volume 51, Number 1 51

III-nitride device fabrication has accelerated within the last

few years. Faced with the challenges resulting from

heteroepitaxial growth, materials scientists have found

innovative ways to improve the crystallinity and to reduce

the levels of undesirable and electronically active impurities

in GaN and other III-nitride materials. Considerable

attention is currently being given to using LEO techniques

to drastically reduce threading defect densities that directly

affect the operation and efficiency of optoelectronic and

microelectronic devices. Having achieved reduced defect

densities and p-type conductivity, the III-nitride community

has finally demonstrated commercially viable GaN-based

devices. The first commercially available LED was announced in 1993 [4]. This was followed by the laboratory

development of the laser diode in 1996 [1]. An improved

version of the LD was fabricated using a LEO base layer in

1997 [5]. Several other investigators have also recently

reported successes in fabricating high electron mobility field

effect transistors [6-9]. Research regarding the III-nitride

materials is now focused on the optimization of these and

other devices; thus, methods to improve heteroepitaxial

growth and to attain further reduction of the threading

defects are once again at the forefront.

Acknowledgments

The authors thank T. Weeks, Jr. and M. Bremser for their

assistance in the growth of the materials described in this

article and for helpful discussion. We also thank Cree

Research, Inc. for the silicon carbide wafers. The NCSU

research discussed in this article was conducted under ONR

contracts monitored by C. Wood and J. Zolper, respectively.

R. Davis was supported in part by the Kobe Steel, Ltd.

University Professorship.

References

1. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada,

T. Matsushita, H. Kiyoku, Y. Sugimoto, T. Kozaki, H.

Umemoto, M. Sano, and K. Chocho, “InGaN/GaN/AlGaNbased Laser Diodes with Modulation-Doped Strained-layer

Superlattices Grown on an Epitaxially Laterally Overgrown

GaN Substrate,” Appl. Phys. Lett. 72, 211 (1998).

2. O. Nam, T. Zheleva, M. Bremser, and R. Davis, “Lateral

Epitaxy of Low Defect Density GaN Layers via Organometallic Vapor Phase Epitaxy,” 39th Electronic Materials Conference, Fort Collins, Colorado, June 25-27 (1997).

3. H. Amano, M. Kitoh, K. Hiramatsu, N. Sawaki, and I.

Akasaki, “P-Type Conduction in Mg-Doped GaN Treated

with Low-Energy Electron Beam Irradiation (LEEBI),” Jpn. J.

Appl. Phys. 28, L2112 (1989).

4. S. Nakamura, T. Mukai, and M. Senoh, “Candella-Class

High-Brightness InGaN/AlGaN Double-Heterostructure BlueLight-Emitting Diodes,” Appl. Phys. Lett. 64, 1687 (1994).

5. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada,

T. Matsushita, H. Kiyoku, Y. Sugimoto, T. Kozaki, H.

Umemoto, M. Sano, and K. Chocho, “Continuous-Wave

Operation of InGaN/GaN/AlGaN-based Laser Diodes Grown

on GaN Substrates,” Appl. Phys. Lett. 72, 2014 (1998).

6. L. Eastman, “GaN Materials for High Power Microwave

Amplifiers,” in Mater. Res. Soc. Symp. Proc. Vol. 512: WideBandgap Semiconductors for High Power, High Frequency,

and High Temperature, Spring 1998.

7. S. Sriram, “RF Performance of AlGaN/GaN MODFETs on

High Resistivity SiC Substrates,” Mat. Res. Soc. Symp Proc.,

Fall 1997.

8. U. Mishra, “GaN/AlGaN Materials for Microwave HEMTs,”

in Mat. Res. Soc. Symp. Proc. Vol. 512: Wide-Bandgap

Semiconductors for High Power, High Frequency, and High

Temperature, Spring 1998.

9. M. Khan, “Recent Progress in Doped Channel GaN-AlGaN

HFETs for High Power Applications,” Mat. Res. Soc. Symp.

Proc. Vol. 512: Wide-Bandgap Semiconductors for High

Power, High Frequency, and High Temperature, Spring 1998.
52 Naval Research Reviews

The Authors

Kevin Linthicum received his B.S. degree from Virginia Polytechnic

and State University. He is currently pursuing his Ph.D. degree at North

Carolina State University. His major interests are the conventional, lateral,

and pendeo-epitaxial growth of III-nitrides using molecular beam and

metalorganic vapor phase epitaxy (MOVPE). He is currently researching the

growth of GaN on silicon substrates. He is a veteran of the U.S. Navy, where

he spent eight years in the nuclear submarine force. He is a plank owner of

the USS Georgia (SSBN 729) to which he was assigned for six years.

Drew Hanser received his B.S. in ceramic engineering from the

University of Missouri at Rolla in 1992. In 1998, he completed his Ph.D. in

materials science and engineering at North Carolina State University, where

his research focused on the MOVPE growth and characterization of GaN

and InGaN materials. He is currently working at NC State in the Materials

Research Center as a postdoctoral researcher, focusing on the large area

growth of low defect density GaN layers using lateral- and pendeo-growth

techniques for microelectronic devices.

Tsvetanka Zheleva is an NRC/NAS Research Associate with the U.S.

Army Research Laboratory at Adelphi, Maryland and an Adjunct Assistant

Professor at the Department of Materials Science and Engineering of North

Carolina State University. She received her B.S. degree in physics in 1980

and her M.S. degree in optics and spectroscopy in 1981 from Sofia University,

Bulgaria. After that, she worked as a junior scientist in The Central Laboratory

of Photo Processes at the Bulgarian Academy of Sciences, studying various

surface characterization techniques such as AES, XPS, LEED, and RBS.

From 1986 to 1990, she worked as a scientist at the Laboratory of Electron

Microscopy in the Applied Mineralogy Institute at the Bulgarian Academy of

Sciences, where she focused her research on structural characterization of

high-temperature superconductors by transmission electron microscopy and

X-ray diffraction. She received her Ph.D. from NCSU in 1995 on atomic scale

characterization of semiconductor heterostructures by TEM. Since 1995 she

has focused her research on wide bandgap materials characterization,

including structural characterization of heterostructures by TEM and X-ray

diffraction, linear elasticity theory, modeling of interfaces, and lateral growth

of GaN.
Volume 51, Number 1 53

Ok-Hyun Nam received his B.S., M.S, and Ph.D. in the Department of

Metallurgical Engineering at Yonsei University, Korea, in 1986,1989, and

1994, respectively. He was a Research Associate in the Korean Institute of

Science & Technology from 1994 to 1995, and in the Materials Research

Center at North Carolina State University from 1995 to 1998. His research

has focused on the MOVPE growth of III-nitrides for optoelectronic and

microelectronic devices, and selective area growth for fabricating field

emission displays and very low dislocation density GaN epilayers. He is

also a member of the research staff at Samsung Advanced Institute of

Technology in Korea, where he is conducting research on the MOVPE growth

of III-nitrides for short wavelength optoelectronic devices including blue/

green light emitting diodes and violet laser diodes.

Robert F. Davis is Kobe Steel, Ltd. Distinguished University Professor

of Materials Science and Engineering and Alumni Graduate Teaching

Professor in the Department of Materials Science and Engineering at North

Carolina State University. He received his Ph.D. in materials science and

engineering from the University of California, Berkeley. He is a Fellow of the

American Ceramic Society. He is the recipient of the ALCOA Distinguished

Research Award, the Alumni Research Award, the ORNL Excellence in

Publications Award, and the Richard M. Fulrath Memorial Award. His research

interests are focused on areas of growth, characterization, and device-related

activities involving SiC, GaN, AlN, and nitride alloy thin films. He has edited

or coedited six books, authored or coauthored approximately 200 chapters

in edited proceedings or in books, published more than 200 peer-reviewed

papers in archival journals, and given more than 150 invited presentations.
aaaaaaaaaaaaaaaaaaaaaaa a

54 Naval Research Reviews

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Volume 51, Number 1 55

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grown GaN-based device, which makes me darn near

perfect!” might be the boast of a new generation of

gallium nitride semiconductor films now being

produced and studied in many research programs

under the sponsorship of ONR. One such research

program is that undertaken at the University of

California at Santa Barbara (UCSB) by Dr. Umesh

Mishra and his colleagues. The almost dislocationfree GaN thin films they produce by lateral epitaxial

overgrowth (LEO) and other techniques hope to

become the bases for new devices critical to DoDrequired communications and sensing applications.

Secure optical communications, wide bandwidth/high

power amplifiers, compact power supplies, and solarblind detectors for, among other uses, missile

detection should all benefit in the near term from

this research. By detailing the activities underway at

UCSB, this article helps to present the story behind

the story of the R&D of this promising class of wide

bandgap semiconductors.

– S.O.
56 Naval Research Reviews

Umesh K. Mishra and Stacia Keller

University of California at Santa Barbara

Santa Barbara, California

Gallium Nitride-Based Solutions for

Communications and Sensing

The Department of Defense has an urgent need to reduce the

size and increase the sensitivity and efficiency of its

communications and sensing systems, and to enhance the

system lifetime in hostile environments. The alloy system

aluminum-gallium-nitride (AlGaN) offers a unique opportunity to provide a quantum leap in achieving these goals. At

UCSB we have had a concerted effort, largely through ONR

support, in the areas of

• Blue/UV light emitters (for secure optical communications)

• AlGaN/GaN high electron mobility transistors

(HEMTs) (for (a) wide bandwidth/high power

amplifiers for reduced radar aperture, and

(b) compact power supplies)

• AlGaN/GaN heterojunction bipolar transistors

(HBTs; same applications as above), and

• AlGaN solar-blind detectors (for missile detection).

Success in these efforts has in the past relied prominently on

our understanding of the growth and doping of these alloys

and their heterojunctions, and we anticipate that this will

remain fundamental for a few years to come. Continued

success of ONR-funded efforts here is based on the unique

and close collaboration of faculty from materials growth to

circuits and is hopefully apparent in the report.

To explore the great potential offered by the semiconductor

material, crystal growth efforts at UCSB were directed

towards the synthesis of high quality GaN epitaxial layers

and AlGaN/GaN heterojunctions. Besides metalorganic

chemical vapor deposition (MOCVD), molecular beam

epitaxy (MBE) has also been applied recently as a growth

technique on MOCVD-grown GaN templates. In the

MOCVD process, GaN layers were grown on c-plane

sapphire, using the precursors trimethylgallium,

thrimethylaluminum, and ammonia. Under optimized

conditions, GaN films with dislocation densities of 4 × 108

cm-2 were obtained, a value which is among the lowest

reported for GaN films grown on sapphire. The 300 K

electron mobility in 4-μm-thick GaN layers reached values

of 820 cm2/Vs (close to the best reported value of 900 cm2/

Vs by S. Nakamura et al. [1 ]) and carrier concentrations of

n = 2 × 1016 cm-3. The photoluminescence of such films is

dominated by band edge emission, even at excitation levels

as low as 2.2 mW/cm2. The GaN surface is atomically

smooth and characterized by steps that are one GaN bilayer

high.

The MOCVD growth of 20- to 30-nm-thick AlGaN layers

on top of semi-insulating GaN films had been optimized

with respect to the structural quality of the AlGaN layers

and the mobility (μ) and the sheet charge (ns) of the twodimensional electron gas (2DEG) forming at the GaN/

AlGaN interface. The structural quality of the AlGaN layers

depended strongly on the growth conditions and the

aluminum composition of the films. Coherently strained

AlGaN layers were obtained under conditions ensuring a

high surface mobility of adsorbed metal species. Thus, stepflow growth of AlxGa1-xN up to xAl = 1 could be achieved at

low ammonia flows during growth (see Fig. 1). Besides the

tendency to grow in an island growth mode, AlGaN layers

with xAl > 0.2 developed stacking fault-related defects,

which typically started to form at dislocations. Thus, the

formation of these defects could be suppressed by growing

on almost dislocation-free laterally epitaxial overgrown

(LEO) GaN. The characteristics of the 2DEG were evaluated by Hall measurements after van der Pauw. The mobility

of the 2DEG decreased with increasing aluminum mole

fraction in the AlGaN layer and increasing surface roughness of the underlying GaN, in agreement with theoretical

considerations that suggest a strong confinement of the

2DEG at the interface [2]. For Al0.33Ga0.67N/GaN heterojunctions, electron mobilities of 1690 cm2

/Vs at 300 K and

4400 cm2/Vs at 15 K were achieved. The sheet charge of the
Volume 51, Number 1 57

2DEG was most strongly influenced by the aluminum

composition in the AlGaN layer and increased with increasing Al mole fraction due to spontaneous polarization and

piezoelectric effects.

AlGaN/GaN HEMTs were fabricated from a variety of

wafers. In terms of power density, GaN HEMTs have

demonstrated superior microwave power performance

compared to gallium arsenide (GaAs) p-HEMTs and other

conventional III-V devices. UCSB’s GaN HEMTs on

sapphire substrates have achieved 3.3 W/mm at 18 GHz [3],

and just recently the power performance could be improved

to 4.2 W/mm at 6 GHz (Fig. 2), which are the highest

reported values for GaN HEMTs on sapphire. An even

higher power value of 6.8 W/mm at 10 GHz had been

recently achieved for AlGaN/GaN HEMTs grown on semiFig. 1 - Atomic force microscopy images of 18-nm-thick Al0.35Ga0.65N layers grown on GaN by MOCVD with an ammonia

flow of (a) 1.5 l/min, (b) 3 l/min, and (c) 6 l/min (grayscale 3 nm). Island growth and defect formation in the AlGaN layer are

suppressed at low NH3 flows during growth.

Fig. 2 - Power performance of a GaN HEMT

grown on sapphire.

insulating SiC substrates [4] that, though attractive, are

currently in limited supply. To reduce the thermal limitations

caused by the sapphire substrate, we chose at UCSB a flipchip bonding approach of GaN HEMTs grown on sapphire

onto AlN substrates [5]. With this approach, 3.2 W output

power from a 1-mm wide GaN HEMT was achieved, which

is close to the best reported power density for large devices

on SiC substrates.

These latest developments in the GaN HEMT device

technology enabled the fabrication of the first GaN-based

broadband power amplifier with performance superior to

that of its GaAs-based counterparts. The modified capacitive-division traveling wave power amplifier (TWPA) circuit

topology was chosen in this work (Fig. 3). The amplifier had

about 7 dB of small-signal gain with a 3-dB bandwidth of

1 to 8 GHz; both input and output return losses were less

than 15 dB. Output power levels of 3.6 W when biased at

18 V and of 4.5 W when biased at 22 V were obtained at

mid-band (4 GHz) (Fig. 4). To our knowledge, this is the

highest output power for a TWPA using solid-state field

effect transistors. The circuit could not be reliably biased

above 20 V, however, because of poor wafer uniformity as

the device technology is still relatively immature. Much

improved performance of the broadband power amplifier is

expected with further developments in (Al,Ga)N material

growth.

To study the effects of an improved structural quality of the

epitaxial layers, AlGaN/GaN HEMTs had been grown on

LEO GaN. The LEO technique results in virtually dislocation-free GaN in the overgrown areas (Fig. 5) [6 ]. While the

forward I-V characteristics of devices on regular and LEO

GaN were similar, the reduced dislocation density in the

LEO GaN had a major impact on the gate leakage: HFETs

0

5

10

15

20

25

30

35

40

0

5

10

15

20

25

30

35

40

-5 0 5 10 15 20 25

Pout

P (dBm) out

PAE (%)

P (dBm) in

 = 4.2 W/mm

PAE = 34%

Biased at Vds = 25 V, Id = 200 mA/μm, f = 6 GHz

all data scaled from Lg = 1 μm, W = 100 mm device

(a) (b) (c)
58 Naval Research Reviews

Fig. 3 - First demonstration of a GaN-based broadband

power amplifier. Photograph of the fabricated GaN

modified TWPA and close-up view of the flip-chip bonded

GaN HEMT (the device is visible through the transparent

sapphire substrate).

Fig. 4 - Device performance of the amplifier:

(a) Power vs frequency when biased at

Vds = 18 V, class AB mode (b) Power sweep

at 4 GHz when biased at Vds = 22V, class AB mode.

on LEO GaN consistently showed much lower (up to two

orders of magnitude) gate leakage current than HFETs on

standard dislocated GaN. Current efforts are directed

towards optimizing the growth and coalescence of insulating

LEO GaN layers to study the impact of a reduced dislocation density on HFET forward current and transconductance

characteristics.

 A further area of research has been the development of

GaN-based HBTs. In May 1998, UCSB demonstrated the

first GaN bipolar transistor (Fig. 6). This device had a DC

current gain of three, achieved with a 200-nm base and

regrown extrinsic base contacts. The HBT program at UCSB

consists of two main objectives: kV range-switching

transistors and high-power microwave devices for amplifier

applications. Recent devices have been tested with an

operating voltage of greater than 100 V. Primary concerns

for the progress of the AlGaN/GaN HBT are improving base

contact quality and device current gain as well as assessing

AC performance. Advances in processing technology have

led to an improvement in regrown base contact quality. Also,

device structures designed to assess the breakdown voltage

of bipolar transistors as a function of collector thickness for

use in kV switching devices are being fabricated, as are

transistors compatible with AC characterization equipment.

Besides the efforts in the field of electronic devices, UCSB

has also demonstrated the first solar-blind (Al,Ga)N-based

high-speed UV detector (with a cutoff wavelength of 290

nm or less) that uses a photodiode structure. A key issue in

the development of high quality photodiodes is the reduction

of defects in the material, particularly threading dislocations. These defects can result in a high leakage current for

the devices, which in turn reduces the available signal-tonoise ratio and, hence, the detectivity of the detectors. To

reduce the high dislocation density usually present in GaN

films, again the LEO technique, with its ability to produce

Fig. 5 - Schematic of the AlGaN/GaN HEMT grown on

lateral epitaxially overgrown (LEO) GaN. The laterally

overgrown regions are virtually free of dislocations.

Dislocations

propagate

vertically

HFET on

LEO GaN

LEO is

dislocation 

HFET on free

regular GaN

Sapphire

AlGaN

SiO2

(a)

0

5

10

15

20

25

30

35

40

-5

0

5

10

15

20

25

30

35

23456789

Freq (GHz)

Pout

Pin (dBm) 

PAE (%)

PAE

0

5

10

15

20

25

30

35

40

0

5

10

15

20

10 15 20 25 30 35

Gain

Pout

Pin (dBm) 

Pout & Gain 

PAE (%)

PAE

(b)
Volume 51, Number 1 59

virtually dislocation-free overgrown regions, was used. An

AlGaN/GaN p-i-n structure was deposited on LEO GaN.

Photodiodes were then fabricated with the active regions

(defined by a 10 μm × 1000 μm mesa) placed so that they

were only over the almost dislocation-free LEO wing

regions and not the dislocated GaN seed regions. Response

times for these diodes were measured to be as low as 4.5 ns

for 90%-to-10% decay time (see Fig. 7). Current-voltage

measurements of the diodes exhibited dark current densities

as low as 10 nA/cm2 at –5 V on many devices. These current

densities were many orders of magnitude lower than those

obtained from the same p-i-n structure deposited on regular

(dislocated) GaN. Spectral response measurements revealed

peak responsivities of 0.05 A/W at a wavelength of 285 nm;

these were of the same order as the peak responsivities on

the equivalent dislocated structure. This indicates a muchimproved signal-to-noise ratio for the devices on the LEO

GaN. The dislocated devices also exhibited a slower

response time and a less sharp cutoff, with significantly

more sub-bandgap response.

Recently, also, first experiments to optimize the electrical

properties of AlGaN/GaN heterostructures by using MBE as

a growth method were undertaken at UCSB. To eliminate

problems related to the sapphire nucleation process and the

resultant high dislocation and extended defect densities

often observed for direct MBE growth of GaN on sapphire,

efforts concentrated on homoepitaxial growth on high

quality GaN on c-plane sapphire templates grown by

MOCVD at UCSB. The studies were performed in a Varian

Gen II MBE system. Active nitrogen was supplied by a

water-cooled EPI Unibulb Nitrogen Plasma source. Conventional effusion cells were used to provide the Group III

elements. On top of an n-type GaN MOCVD template, an

approximately 250 nm thick unintentionally doped MBEGaN layer was deposited, followed by ~50 nm of MBEAl0.07Ga0.93N. Temperature-dependent Hall measurements

(van der Pauw configuration) revealed mobility values of

14,500 cm2/Vs at 77 K, and 20,000 cm2/Vs at 12 K (see Fig.

8), which are the highest reported values for an AlGaN/GaN

2DEG (recently verified at AFRL). Since the sample was

grown on an n-type GaN template, the room temperature

measurements were significantly affected by parallel

conduction of low-mobility carriers in the template. Using a

Fig. 7 - First solar-blind high speed AlGaN/GaN photo-diode.

(a) Typical pulsed response at 263 nm for diodes fabricated

on the virtually dislocation-free wing region of LEO GaN in

comparison to diodes fabricated on dislocated GaN. (b)

Spectral response curves for diodes fabricated on the wing

region of LEO GaN and on dislocated GaN.

Fig. 6 - First GaN-based HBT. Common emitter characteristics of a 1 μm × 20 μm HBT. The base current was

stepped from 0 to 350 μA in 50 μA steps.

0 10 0

400 10 -6

800 10 -6

1.2 10 -3

0 5 10 15 20 25 30 35

Applied Bias (Common Emitter) [V] 

Collector Current [A] 

β = 2/div

10–4

10–3

10–2

10–1

10–5

200 250 300 350 400

Wavelength, λ (nm)

Responsivity, R (A/W)

dislocated

wing

detection limits of

experimental setup

0

0

0.2

0.4

0.6

0.8

1.0

1.2

20 40 60 80 100

t (ns)

dislocated

Photocurrent (arbitrary units)

wing

(b)

(a)
60 Naval Research Reviews

two-layer conduction model, with one layer being the

template and the second layer the 2DEG, a room-temperature mobility of 1860 cm2/Vs and a carrier concentration of

4.8 × 1012 cm-2 were extracted. Further MBE growth studies

on semi-insulating GaN on sapphire wafers for device

applications are under way.

In summary, UCSB’s research team of the authors and

Professors S. DenBaars, J. Bowers, E. Hu, M. Rodwell, and

R. York have undertaken a wide ranging effort to utilize the

potential of the alloy system (Al,Ga)N for semiconductor

devices, both in the fields of crystal growth and device

processing and application of new device concepts. As a

result, several AlGaN/GaN-based devices have been

demonstrated for the first time. Although the device performance is still limited by the relative immaturity of the

semiconductor material, problems have been analyzed and

will be addressed in future work.

References

1. S. Nakamura, T. Mukai, and M. Senoh, “In-situ Monitoring

and Hall Measurements of GaN Grown with GaN Buffer

Layers,” J. Appl. Phys. 71, 5543 (1992).

2. Y. Zhang and J. Singh, “Charge Control and Mobility Studies

for an AlGaN/GaN High Electron Mobility Transistor,” J.

Appl. Phys. 85, 587 (1999).

3. Y.-F. Wu, B.P. Keller, P. Pusl, M. Le, N.X. Nguyen,

C. Nguyen, D. Widman, S. Keller, S.P. DenBaars, and

U.K. Mishra, “Short-Channel Al0.5Ga0.5N/GaN MODFETs

with Power Density > 3 W/mm at 18 GHz,” submitted to

Electron. Lett.

Fig. 8 - World-record 2DEG electron mobility for an

Al0.07Ga0.93N/GaN heterojunction achieved by MBE

growth. Temperature dependence of the electron

mobility and the carrier density.

103

104

2

1

0.9

0.8

0.7

0.6

0.5

0.4

Sheet Density NS,1013 cm–2

0 50 100 150 200 250 300

Temperature, K

Mobility 

μ, cm2/V s

4. S.T. Sheppard, K. Doverspike, W.L. Pribble, S.T. Allen, and

J.W. Palmour, “High Power Microwave GaN/AlGaN HEMTs

on Silicon Carbide,” 56th Annual Device Research Conference, June 1998.

5. B.J. Thibeault, B.P. Keller, Y.-F. Wu, P. Fini, U.K. Mishra,

C. Nguyen, N.X. Nguyen, and M. Le, “High Performance

and Large Area Flip-Chip Bonded AlGaN/GaN MODFETs,”

56th Annual Device Research Conference, June 1998.

6. D. Kapolnek, S. Keller, R. Vetury, R.D. Underwood,

P. Kozodoy, S.P. DenBaars, and U.K. Mishra, “Anisotropic

Epitaxial Lateral Growth in GaN Selective Area Epitaxy,”

Appl. Phys. Lett. 71, 1204 (1997); H. Marchand, X.H. Wu,

J.P. Ibbetson, P.T. Fini, P. Kozodoy, S. Keller, J.S. Speck,

S.P. DenBaars, and U.K. Mishra, “Microstructure of GaN

Laterally Overgrown by Metalorganic Chemical Vapor

Deposition,” Appl. Phys. Lett. 73, 747 (1998).

The Authors

Umesh K. Mishra received his Ph.D. from Cornell

University, Ithaca, New York, in 1984. He has worked in

various laboratory and academic institutions, including

Hughes Research Laboratories in Malibu, California,

University of Michigan at Ann Arbor, and General Electric

in Syracuse, New York, where he has made major

contributions to the development of AlInAs-GaInAs HEMTs

and HBTs. He is currently a Professor in the Department

of Electrical and Computer Engineering, University of

California at Santa Barbara. His current research interests

are in oxide-based III-V electronics and III-V nitride

electronics and optoelectronics. He has authored or

coauthored more than 300 papers in technical journals

and conferences and holds six patents. Dr. Mishra received

the Hyland Patent Award given by Hughes Aircraft and

the Scientist of the Year Award presented at the

International Symposium on GaAs and Related

Compounds.

Stacia Keller received the Diploma and Ph.D. degree

in chemistry from the University of Leipzig, Germany, in

1983 and 1986, respectively. As a scientific assistant, she

continued her experimental and theoretical work in the

field of epitaxial growth of GaAs- and InP-based

semiconductors. In 1994, she joined the Electrical and

Computer Engineering Department, University of

California at Santa Barbara. Her current research interests

are the crystal growth and characterization of Group IIInitrides for electronic and optoelectronic applications. She

has authored or coauthored more than 100 technical

publications and conference presentations and holds one

patent.
Volume 51, Number 1 61

You can give the Naval Research Laboratory a green-thumbs

up for its recent harvests of GaN films grown on sapphire

substrates. From this remarkable crop, NRL researchers have

produced high electron mobility transistor (HEMT) devices

that will feed the hungry mouths of DoD’s high-end

consumers of low-frequency power switching devices, highfrequency microwave radar amplification, and electronic

warfare. By cultivating these controllably

doped GaN

films, NRL is

helping to plant

the seeds of

victory for

tomorrow’s

combatant.

In this article, Dr. Daniel Koleske and other NRL researchers

discuss the GaN materials requirements and present a model

that shows the required relationship between temperature

and V/III ratio for achieving near-stoichiometric GaN growth.

They also describe how pressure influences GaN

decomposition rate: understanding that is key to weeding

out defects and increasing grain size during growth, thus

yielding a crop of GaN thin films that have better electronic

properties. The authors also give the formula for avoiding

those pesky trapping-type defects in GaN MODFET devices.

By reducing trapping-type defects in GaN-based devices and

otherwise improving GaN thin films, NRL materials scientists

are helping the Navy and all of DoD to reap the benefits of

its research seed money.

– S.O.
62 Naval Research Reviews

Daniel D. Koleske, Alma E. Wickenden, Richard L. Henry, Mark E. Twigg, Steven C. Binari,

Paul B. Klein, Jaime A. Freitas, Jr., James C. Culbertson, and Mohammad Fatemi

Naval Research Laboratory

Washington, D.C.

GaN-Based Materials Growth

for Microwave Electronic

Device Applications

Introduction

GaN-based electronic devices have demonstrated excellent

potential for meeting future Navy needs in microwave radar,

electronic warfare, communications, and power switching.

Recently, Sheppard et al. [1] have demonstrated working

devices with an output power density of 6.8 W/mm and

larger devices with a total power output of 9.1 W (3 W/mm).

Despite this success, a number of growth and fabricationrelated issues remain that must be solved before these

devices are inserted in the Fleet.

Many of the growth-related issues emanate from the lack of

a suitable substrate for GaN growth. GaN thin films are

typically grown on sapphire or silicon carbide (SiC), which

do not lattice match GaN well. Because of the mismatch,

GaN has a large number of defects and dislocations.

However, in working GaN-based LEDs, the dislocation

density can number as high as 1010 cm-2 [2], demonstrating

that usable devices can be fabricated. Until GaN substrates

are available (which is not likely in the next 5 years), most

GaN growth efforts will be aimed at further reduction of

defects to improve device performance.

The link between material quality and device performance

has been demonstrated in the work of Shuji Nakamura and

coworkers at Nichia Chemical Industries, Ltd. Nakamura

began work on GaN growth in 1989, which led to the

invention in 1990 of a reactor specifically for nitride growth

[2]. After his invention, Nakamura quickly demonstrated

world-record, 300 K electron mobilities of 930 cm2V-1s-1. In

1994, Nakamura demonstrated a 2-candella blue LED [2],

and in 1998 a commercial blue laser [3]. It was the early

success in high-quality GaN growth that led to these device

successes.

At NRL, we have a focused program in the growth, characterization, and fabrication of GaN-based microwave devices.

This in-house capability allows accelerated development of

our materials and devices to meet emerging Navy device

requirements. In this article, we highlight some of our recent

work that has focused on the MOVPE growth of highquality GaN films and methods to reduce defects in GaN.

Types of Devices and Materials Requirements

For microwave power operation, the primary emphasis is on

the development of the high electron mobility transistor

(HEMT) structure, as shown in Fig. 1. The active region of

this device is a thin conducting GaN layer, which is created

by placing a thin AlGaN (20 to 30 nm) layer on the insulating GaN layer. The AlGaN layer creates a 2-dimensional

electron gas (2DEG) in the GaN region (denoted by the

dashed line) near the AlGaN/GaN interface. In optimized

Fig. 1 - In this AlGaN/GaN HEMT device cross section,

the yellow area represents the proton-implanted regions.
Volume 51, Number 1 63

AlGaN/GaN films, we have recently achieved electron

mobilities as high as 1530 cm2V-1s-1 at sheet carrier concentrations of 1.2 × 1013 cm-2.

The most stringent requirement for the HEMT structure is

the growth of high quality, insulating GaN. In addition to

being insulating, these layers must have high electron

mobility when the AlGaN layer is placed on top. Defects

located in the insulating GaN layer have been shown to

restrict electron conduction through the device layer.

Because the quality of the insulating GaN layer has been

shown to play a role in microwave device performance, we

are focusing on the growth of high quality, insulating GaN.

In 1995, NRL researchers showed how trapping-type defects

influence the characteristics of GaN-based devices [4]. An

example of trapping-type defects limiting the power

performance is shown in Fig. 2 [5]. Each of the HEMT

devices in Fig. 2 has similar DC and small-signal characteristics; however, for each device, the drain current responds

differently when short (< 1 μs) gate-voltage pulses are

applied. Due to trapping effects, devices that have a poor

pulse response characteristic (i.e., low Ipulse/Idc) also have

lower power output, as shown in Fig. 2. The lower current

measured under pulse conditions can be attributed to the

time dependence of the trapping mechanism. Under microwave test conditions, this trapping mechanism prevents the

drain current from responding to the applied gate signal,

resulting in lower power output.

Traps also reduce drain current at low voltages leading to

current collapse in the device. This is shown in Fig. 3, where

the drain current is plotted vs voltage in the dark (dashed

lines) and under white light illumination (solid lines) [6].

The low voltage current collapse (in the dark) is consistent

with electron trapping in the insulating GaN layer. The

trapped electrons form a space charge region at the bottom

of the conductive channel, pinching off the device and

limiting the current flow through the device. With the light

on, the electrons are excited out of these trap states and the

space charge region does not develop. As a result, a normal

set of drain characteristics is measured. From these examples, it is evident that trapping-type defects play a role in

limiting the performance of GaN-based microwave devices.

GaN Growth on Sapphire

Prior to the 1990s, there was little interest in GaN as a

viable semiconductor, because of the large number of n-type

defects, low electron mobilities, and rough surface morphology. This changed by 1990 because of the discovery of

improved GaN growth on low-temperature AlN nucleation

layers (NL) deposited between sapphire and GaN [7]. With

the NL innovation, smooth GaN films with improved

structural and electrical properties were produced. Using

MOVPE growth, GaN films can be grown n-type with high

mobility, which is ideally suited for LEDs and lasers [2], or

insulating, which is ideally suited for microwave FET

devices.

Fig. 2 - The figure shows the correlation of microwave

power output at 2 GHz with drain current response to

gate voltage pulses shorter than 1 μs. The gate voltage

is pulsed from a level less than the threshold voltage to

0 V. The drain voltage is maintained at a fixed value.

Fig. 3 - Drain current is plotted as a function of the

drain voltage with white light illumination (solid

lines) and in the dark (dashed lines). In the dark, the

drain current collapses for drain voltages less than

50 V are evident after the first curve is measured.
64 Naval Research Reviews

(a) (b)

In typical GaN films, the electron mobility drops as the

n-type carrier concentration drops below 1017 cm-3. The

classical explanation for this observation is that the material

is compensated with acceptor-type point defects. Another

proposed explanation is that the electron mobility drops

because the electrons scatter at the grain boundaries [8].

Both descriptions probably apply to GaN because GaN

films are composed of individual grains and contain deep

donor and acceptor levels. The GaN grain structure is shown

in Fig. 4. The grains can be tilted with respect to the

sapphire normal (c-axis) and the grains can also be rotated

about the c-axis. Typically, this degree of disorder in

semiconductors leads to atomic miscoordination in the

lattice, which results in energy levels in the bandgap.

Because of the presence of grain boundaries, a similarity has

been suggested in the electrical conduction between GaN

films and polycrystalline Si (poly-Si) [8]. In models

developed to explain hole conduction in poly-Si, it was

shown that poly-Si behaves like single crystal Si when the

doping level is substantially larger than the trap density [9].

However, when the doping and trap densities are similar, the

hole mobility drops and the resistivity increases by orders of

magnitude. In a similar way, deep level acceptor concentrations in GaN in excess of the shallow donor concentration

could cause the resistivity to increase and the mobility to

drop. At a fixed trap density in poly-Si, the transition from

high mobility to high resistivity shifts from higher to lower

dopant concentration as the grain size increases. For poly-Si

it has been shown that increases in grain size directly

correlate with improved hole mobility. Similar correlations

between increased grain size and improved electron mobility have been observed in other polycrystalline compound

semiconductors [10].

Toward Defect and Trap Identification

and Reduction in GaN

To improve the GaN films, our objectives, presented in the

following subsections, include developing growth methods

to reduce off-stoichiometric-type defects in GaN, reducing

misorientation in our GaN films, and increasing the GaN

grain size. In addition, we are working to reduce trappingtype defects, to identify these defects in fabricated GaN

devices, and to design vertical GaN-based devices.

Kinetic Model to Explain

Stoichiometric GaN Growth

In this section, we show how two growth

parameters, growth temperature (TG) and

the V/III ratio (the ratio of N to Ga), are

linked to each other and to the resultant

GaN quality. Figure 5(a) is a view of one

of our MOVPE reactors. The reactor has a

showerhead for the separate delivery of the

trimethylgallium (TMGa) and ammonia

(NH3) reactants. Since reactor designs vary

widely for GaN growth, a more fundamental understanding of the GaN growth

chemistry is required. This understanding

would facilitate the GaN process development independent of reactor geometry.

Figure 5(b) shows five reaction steps

important to growth. These include

1) adsorption and decomposition of the

reactants to the growth species, 2) desorption, 3) surface

diffusion, 4) incorporation into the solid, and 5) decomposition of the solid. Each of these reaction steps and their

relevance to GaN growth are detailed in a recent article [11].

The primary finding in this article is that near-stoichiometric

GaN growth can be achieved when the reactant flux ratio of

NH3 to TMGa (i.e., FNH3/FTMGa or V/III ratio) is equal to the

N to Ga desorption ratio (i.e., kN/kGa). Since the N and Ga

desorption rates depend exponentially (Arrhenius behavior)

on temperature, their ratio (i.e., kN/kGa) also depends

exponentially on temperature. When FNH3/FTMGa ≈ kN/kGa is

used for GaN growth, the N and Ga surface concentrations

are close to stoichiometric [11].

Fig. 4 - This GaN microstructure shows the ordered polycrystalline structure

of the individual GaN grains. The side view (a) shows the relative tilt of each

grain, and the top view (b) shows the twist of the grains with respect to one

another. Edge-type dislocations mainly originate at the grain boundaries, while

screw-type dislocations are generally found throughout the entire GaN film

(modeled after Hersee [7]).
Volume 51, Number 1 65

(b)

(a)

This model explains many GaN growth properties. For

example, it explains the transition from smooth to rough

morphology. Figure 6(a) presents Normarski pictures of

both smooth and rough GaN surface morphology. The solid

line in Fig. 6(a) is the N to Ga desorption ratio, kN/kGa. The

V/III ratio and TG from four growth runs are plotted in Fig.

6(a). For these runs, the V/III ratio (i.e., FNH3/FTMGa) was

varied at TG = 1040 °C. Three of the growth runs had a

smooth surface morphology (blue circles), while one of the

runs had a rough surface morphology (the red square). Note

that growth conditions (i.e., V/III ratio and TG) that produced the smooth morphology films are all above the kN/kGa

line, while the growth condition for the rough morphology

film falls below the kN/kGa line. The surface morphology

becomes rough because not enough N is on the surface and

Ga droplets form.

When the growth conditions used by other researchers are

compared with the kN/kGa line in Fig. 6(b), the line properly

describes the transition from smooth to rough morphology.

As in Fig. 6(a), the red squares denote growth conditions

resulting in rough morphology, while the blue circles denote

growth conditions resulting in either smooth surface

morphology or optimized film growth. Note that all but

three of the smooth morphology growth conditions (blue

circles) are above the kN/kGa line, while all of the rough

morphology growth conditions (red squares) fall below the

kN/kGa line.

Figure 7 compares the optimized V/III ratio and TG from

five different groups to the kN/kGa line (blue circles). For

each of these points, the V/III ratio and TG were independently varied to determine the optimal growth conditions.

As shown in Fig. 7, the optimized values of V/III ratio and

TG correlate well with the kN/kGa line, demonstrating that

FNH3/FTMGa ≈ kN/kGa for optimized growth. Since different

MOVPE reactors were used for each of five data points, the

GaN growth model developed at NRL satisfies one of our

goals, namely to show how the GaN growth parameters

Fig. 5 - (a) Side view of the close-spaced showerhead reactor.

This reactor is one of the two MOVPE reactors currently

used at NRL for the growth of GaN. (b) The individual

reaction steps considered for the kinetic model [11].

Fig. 6 - (a) The values of the V/III ratio and growth temperature are shown with pictures of the resulting surface

morphology. The filled blue circles denote smooth morphology, and the red squares denote rough morphology. The

solid line is the kN/kGa desorption ratio line. (b) Similar to

plot (a), except that the smooth morphology or optimized

growth conditions (blue circles) and rough morphology film

(red squares) from over 43 literature references are plotted

[10].
66 Naval Research Reviews

must be selected with respect to one another independent of

reactor design.

Study of GaN Decomposition

The MOVPE growth of GaN is conducted at temperatures

that typically range from 900 to 1100 °C. High temperatures

are necessary to dissociate NH3 and form atomic N at the

growth surface. This temperature range is substantially higher

than the onset temperature for GaN decomposition, which

has been reported to be as low as 600 °C. In early GaN

growth studies, GaN decomposition was identified as an

important chemical reaction detrimental to growth, because it

decreases the GaN growth rate.

Recently, we have studied GaN decomposition in flowing H2,

N2, and NH3 at MOVPE growth temperatures and pressures

[12]. For GaN films annealed in H2, we found an increase in

the GaN decomposition rate by increasing the chamber

pressure. Figure 8 plots the GaN decomposition rate (blue

diamonds) as a function of total H2 pressure for 10-min

anneals at a temperature of 992 °C. The GaN decomposition

rate was found to increase by a factor of 3 when the pressure

was increased from 10 torr to 700 torr. At lower temperature,

the increase in the decomposition rate vs pressure is even

more dramatic. Also plotted in Fig. 8 are the measured rates

for Ga desorption (red squares) and liquid Ga droplet buildup

(green circles). The droplets increase in size as the pressure is

increased, as shown on top of Fig. 8. The increase in the GaN

decomposition rate at higher pressures is due to an increased

dissociation of H2, which then combines with surface N to

reform NH3 [13]. At higher pressures, more H2 is dissociated,

which leads to enhanced GaN decomposition.

While detrimental to growth, enhancing the GaN decomposition rate at high pressures brings the growth closer to

equilibrium. The enhanced decomposition helps to remove

weakly bonded Ga and N atoms, thereby increasing order in

the GaN crystal structure. Also, the higher decomposition

rate may prevent smaller GaN grains from forming on the

NL. As a result, the GaN grain size is increased. The

increase in grain size and subsequent improvement in

electrical properties at higher pressures has been observed in

films grown at NRL (see below) and reported elsewhere

[14].

When NH3 is added to the H2 flow, the GaN decomposition

rate was reduced to 1 × 1017 cm-2s-1. This is highly surprising, because GaN growth rates are typically 1 × 1015 cm-2s-1

or 1 μm/hour. Put another way, the GaN decomposition rate

is 100 times greater than the GaN growth rate. This means

that GaN growth is under kinetic control, and this agrees

with the initial assertion of the kinetic model described

earlier.

GaN Orientation on A-Plane Sapphire

A systematic study was also performed at NRL to optimize

the GaN NL in a new close-spaced showerhead reactor

geometry [15]. In this study, the sapphire nitridation and low

temperature GaN NL were varied. We found that the NH3

exposure for optimal nitridation depended on the temperature, NH3 flow, and time. We also found that the nucleation

Fig. 7 - The values of the V/III ratio and growth temperature for optimized GaN film growth from five different

groups are compared to the kN/kGa desorption ratio line.

Fig. 8 - The GaN decomposition rate (blue diamonds), Ga

desorption rate (red squares), and liquid Ga accumulation

rate (green circles) for 10-min anneals in H2 at 992 oC are

plotted as a function of reactor pressure. Also shown are

Normarski phase contrast images of the surface morphology after annealing with arrows indicating the anneal

pressure.
Volume 51, Number 1 67

site density is strongly influenced by growing the NL at a

reduced V/III ratio. AFM and TEM studies of the initial

stages of GaN growth on the NL demonstrated improvements in grain orientation as the nitridation and nucleation

layer process parameters were varied. For GaN film grown

at 76 torr, the electron mobility was 300 cm2V-1s-1 for Sidoped films.

Other work has focused on improving grain alignment by

growth on slightly vicinal a-plane sapphire [16]. Hall

mobility and X-ray rocking curve data showed a nearly

twofold improvement for simultaneous GaN growths on

slightly misoriented sapphire compared to on-axis sapphire.

TEM revealed a reduction in the edge dislocation density,

suggesting better grain alignment on the misoriented

substrates, while photoluminescence measurements demonstrated a pronounced decrease in the yellow band emission

between the vicinal and on-axis a-plane sapphire substrates.

These studies highlight the importance of grain alignment

for the growth of high quality GaN.

Development of Large Grain GaN

at High Pressure

After optimizing sapphire nitridation and NL growth, GaN

growth was conducted at pressures greater than 76 torr [15].

When the NL ramp and GaN film growth were conducted at

150 torr, the electron mobilities were nearly doubled

compared to similar growths at 76 torr. A TEM cross section

of the GaN film grown at 150 torr is shown in Fig. 9. This

film had large grains (4 μm in Fig. 9) and a dislocation

density on the order of 108 cm-2.

By controlling the sapphire nitridation, NL growth, and

growth pressure, we increased the GaN grain. Since the

article by Wickenden and coworkers at NRL [15], mobilities

exceeding 700 cm2V-1s-1 for Si doping levels of 1.4 × 1017

cm-3 have been achieved in the showerhead reactor. In

addition, insulating GaN suitable for HEMT development

was achieved on AlN nucleation layers. Initial AlGaN/GaN

HEMT structures on the insulating GaN have measured 300

K mobilities exceeding 1400 cm2V-1s-1 and sheet carrier

concentration of 1.2 × 1013 cm-2 over the 2-in. wafer. TEM

cross section measurements of this AlGaN/GaN HEMT

structure reveal large grains in the insulating, underlying

GaN. Initial devices fabricated on the HEMT structure show

no significant trapping effects. In agreement with the

models of electrical conduction in polycrystalline semiconductors, increasing the GaN grain size has a large impact on

improving electronic properties [17], especially the electron

mobility.

Growth Pressure Influence

on GaN Electronic Properties

Typically, to achieve the desired insulating GaN for

HEMTs, growth pressures of 76 torr or less are used. If the

insulating GaN has a high mobility when it is Si doped or a

AlGaN layer is grown, the films are deemed suitable for

HEMT devices. Recently, we have observed that the GaN

films change from insulating at 49 torr to conductive (ntype) when the growth pressure is increased above 100 torr.

The electron mobility reached a maximum at 250 torr where

mobilities of 600 cm2V-1s-1 with 9 × 1016 cm-3 n-type

carriers were measured. We have tentatively assumed that

the loss of the insulating GaN quality is due to a reduced

Fig. 9 - Shown is a TEM cross

section of optimized GaN

growth on sapphire. Individual

GaN grains can be imaged by

slightly tilting the GaN sample

in the TEM.
68 Naval Research Reviews

deep acceptor (or trap) concentration. We are currently

investigating the influence of growth pressure on GaN film

quality and possible trapping-type defects, which are

introduced at lower growth pressures.

Toward Identification of Trapping-Type Defects

Research is also underway at NRL to identify trapping-type

defects. These defects have been observed to influence

maximum operational frequency, power output, leakage

current, current collapse, and transconductance dispersion

[6]. Persistent photoconductivity (PPC) studies in a number

of GaN device structures suggest a distribution of trap levels

in the GaN material [18]. Previous device [6] and PPC [18]

studies suggest that the traps are located in the insulating

GaN.

At NRL, a spectroscopic method has been developed to

detect and measure trapping-type influences in GaN FET

device structures. Binari et al. previously observed that the

drain current was restored in GaN MESFETs when light

was applied [6]. Shorter wavelength light was found to be

more efficient in restoring the drain current (i.e., decreasing

the current collapse), suggesting the emission of trapped

electrons through a photoionization process. The increase in

the drain current in light vs dark (∆I) normalized to the drain

current in the dark (ID) and the incident photon flux (Φ(hν))

is plotted in Fig. 10(a) as a function of photon energy for

GaN MESFET devices fabricated at NRL [19]. Two distinct

trap levels are obtained from the fits. The photoionization

spectrum is energy broadened because of strong coupling

between the trap levels and the GaN vibrational levels as

shown in Fig. 10(b). This vibronic coupling explains why a

distribution of trap levels was initially proposed in the PPC

studies. Further investigations of trapping-type defects are

currently underway using photoluminescence, magnetic

resonance, and current collapse spectroscopy (described

above), which is unique because it directly measures the trap

characteristics in working device structures.

Vertical Power Device Structures

If the assertion that grain size limits electron mobility in

GaN films is true, device designs where the electrons do not

cross grain boundaries should be considered. Electrons

conducted vertically through the GaN will not scatter from

grain boundaries. We are currently designing a vertical

power transistor to test the benefits of this type of structure.

Future Research Focus

In this article, we have highlighted the current efforts in

GaN materials growth, characterization, and device fabrication. Our research on GaN growth kinetics [11], GaN

decomposition [12], and surface nitridation and nucleation

layer growth [15] has been essential to achieving the high

quality, high mobility GaN films described earlier. Our

observation from TEM data that GaN grain size correlates

with electron mobility [17] drove our selection of which

growth parameters to vary. These insights allowed us to

optimize GaN film growth in a new reactor design and

achieve record mobilities for this reactor.

Although many GaN-based devices with outstanding

properties have been fabricated, the scientific understanding

necessary to routinely reproduce these results and advance

device performance is currently lacking in many critical

areas. This is especially true for many of the individual

MOVPE growth steps that influence grain orientation and

grain size. Further reduction in trap density is a high priority

at NRL because such reduction should directly lead to

improvements in device performance. This research is

critical in order for GaN to reach its full potential for

microwave power electronics.

Fig. 10 - (a) The difference in drain

current in light vs dark (∆I) is

plotted vs the photon energy. The

value of ∆I is normalized to the

drain current in the dark and to the

photon energy. (b) The trap energy

levels are plotted as a function of

the lattice distortion. Vertical

transitions between the trap energy

levels and the conduction band

(CB) lead to vibrational excitation

of the lattice, which gives the

appearance of a broadened

distribution of trap energy levels.
Volume 51, Number 1 69

Acknowledgments

The authors acknowledge support from the NRL base

program and the Office of Naval Research, including partial

support from the RF Solid State Technology program and

the Power Electronic Building Block (PEBB) program. The

authors acknowledge the following individuals for their

contributions to the GaN effort: D.K. Wickenden, J.M.

Koleske, R.J. Gorman, L.T. Ardis, W.J. DeSisto, K.

Doverspike, L.B. Rowland, D.K. Gaskill, J.E. Butler, G.

Beadie, W.S. Rabinovich, V. Bermudez, M. Peckerar, H.B.

Dietrich, W. Kruppa, D. Park, R. Bass, D. McCarthy, W.E.

Carlos, G. Keller, K. Ikossi-Anastasiou, E.R. Glaser, T.A.

Kennedy, B.V. Shanabrook, G. Jernigan, J.M. Killiany, P.E.

Thompson, and G.M Borsuk.

References

1. S.T. Sheppard, K. Doverspike, W.L. Pribble, S.T. Allen, J.W.

Palmour, L.T. Kehias, and T.J. Jenkins, “High Power GaN/

AlGaN HEMTs on Silicon Carbide,” to be published, IEEE

Elect. Dev. Lett., April 1999.

2. S. Nakamura and G. Fasol, The Blue Laser Diode - GalliumNitride Based Light Emitters and Laser (Springer-Verlag,

New York, 1997).

3. S. Nakamura, “InGaN/GaN/AlGaN-Based Laser Diodes with

an Estimated Lifetime of Longer than 10000 Hours,” Mat.

Res. Soc. Symp. Proc. 482, 1145-1156 (1998).

4. W. Kruppa, S.C. Binari, and K. Doverspike, “Low-Frequency

Dispersion Characteristics of GaN HFETs,” Electron. Letts.

31, 1951-1952 (1995).

5. S.C. Binari, K. Ikossi-Anastasiou, W. Kruppa, H.B. Dietrich,

R.L. Henry, D.D. Koleske, and A.E. Wickenden, “Correlation

of Drain Current Pulsed Response with Microwave Power

Output in AlGaN/GaN HEMTs,” submitted to the Spring

1999 Materials Research Society.

6. S.C. Binari, W. Kruppa, H.B. Dietrich, G. Kelner, A.E.

Wickenden, and J.A. Freitas, Jr., “Fabrication and Characterization of GaN FETs,” Solid-State Electron. 41, 1549-1554

(1997).

7. H. Amano, N. Sawaki, I. Akasaki, and Y. Toyoda,

“Metalorganic Vapor Phase Epitaxial Growth of a High

Quality GaN Film Using an AlN Buffer Layer,” Appl. Phys.

Lett. 48, 353-355 (1986).

8. S.D. Hersee, J.C. Ramer, and K.J. Malloy, “The Microstructure of Metalorganic-Chemical-Vapor-Deposition GaN on

Sapphire,” MRS Bull. 22, 45-51 (1997).

9. D.M. Kim, A.N. Khondker, S.S. Ahmed, and R.R. Shah,

“Theory of Conduction in Polysilicon: Drift-Diffusion

Approach in Crystalline-Amorphous-Crystalline Semiconductor System – Part I: Small Signal Theory,” IEEE Trans. Elect.

Dev. ED-31, 480-493 (1984).

10. L.L. Kazmerski, Electrical Properties of Polycrystalline

Semiconductor Thin Film (Academic Press, New York, 1980)

pp. 59-133.

11. D.D. Koleske, A.E. Wickenden, R.L. Henry, W.J. DeSisto, and

R.J. Gorman, “Growth Model for GaN with Comparison to

Structural, Optical, and Electronic Properties,” J. Appl. Phys.

84, 1998-2010 (1998).

12. D.D. Koleske, A.E. Wickenden, R.L. Henry, M.E. Twigg, J.C.

Culbertson, and R.L. Gorman, “Enhanced GaN Decomposition in H2 Near Atmospheric Pressures,” Appl. Phys. Lett. 73,

2018-2020 (1998).

13. C.D. Thurmond and R.A.Logan, “The Equilibrium Pressure

of N2 over GaN,” J. Electrochem. Soc. 119, 622-626 (1972).

14. J. Han, T.-B. Ng, R.M. Biefeld, M.H. Crawford, and D.M.

Follstaedt, “The Effect of H2 on Morphology Evolution

During GaN Metalorganic Chemical Vapor Deposition,” Appl.

Phys. Lett. 71, 3114-16 (1997).

15. A.E. Wickenden, D.D. Koleske, R.L. Henry, R.J. Gorman,

J.C. Culbertson, and M.E. Twigg, “The Impact of Nitridation

and Nucleation Layer Process Conditions on Morphology and

Electron Transport in GaN Epitaxial Films,” J. Elect. Mat. 28,

299-305 (1999).

16. M. Fatemi, A.E. Wickenden, D.D. Koleske, M.E. Twigg, J.A.

Freitas, Jr., R.L. Henry, and R.J. Gorman, “Enhancement of

Electrical and Structural Properties of GaN Layers Grown on

Vicinal-Cut, A-Plane Sapphire Substrates,” Appl. Phys. Lett.

73, 608-610 (1998).

17. M.E. Twigg, R.L. Henry, A.E. Wickenden, D.D. Koleske, and

J.C. Culbertson, “Nucleation Layer Microstructure, Grain

Size, and Electrical Properties in GaN Grown on A-Plane

Sapphire,” submitted to Appl. Phys. Lett.

18. G. Beadie, W.S. Rabinovich, A.E. Wickenden, D.D. Koleske,

S.C. Binari, and J.A. Freitas, Jr., “Persistent Photoconductivity in n-type GaN,” Appl. Phys. Lett. 71, 1092-1094 (1997).

19. P.B. Klein, J.A. Freitas, Jr., and S.C. Binari, “Photoionization

Spectra of Traps Responsible for Current Collapse in GaN

MESFETs,” presented at the American Physical Meeting,

March 1999.
70 Naval Research Reviews

The Authors

Daniel D. Koleske is a research chemist at NRL conducting MOVPE growth of GaN.

Dr. Koleske specializes in kinetic descriptions of semiconductor growth. Prior to NRL,

Dr. Koleske held a National Research Council (NRC) postdoctoral fellowship at NRL.

Before the NRC fellowship, Dr. Koleske was a postdoctoral researcher at IBM in Yorktown

Heights, NY, working on Group IV growth kinetics. Dr. Koleske obtained a B.S. in

chemistry at the University of Wisconsin, Madison, and a Ph.D. in physical chemistry at

the University of Chicago.

Alma E. Wickenden is a materials research engineer at NRL involved in the crystal growth of

GaN films. Dr. Wickenden received her B.S. in physical science and her M.S. and

Ph.D. in materials science and engineering from Johns Hopkins University. She has

experience in VPE GaAs growth at Westinghouse Electric Corporation, and lithographic

and device process development at Bendix Corporation.

Richard L. Henry is currently the section head of the Thin Film Deposition Section at NRL and is

involved in the MOVPE growth of GaN. Dr. Henry received his B.A. in chemistry at

Millersville State University in 1972 and his Ph.D. in chemistry at Brown University in

1976.

Mark E. Twigg is a materials research engineer at NRL, conducting TEM studies of GaN thin film

structure. Dr. Twigg received his B.S. in physics at the University of Maryland and his

M.S. and Ph.D. in metallurgical engineering at the University of Illinois at Urbana/

Champaign. Before coming to NRL, Dr. Twigg held postdoctoral positions at the National

Bureau of Standards and at AT&T Bell Laboratories.

Steven C. Binari received the B.S. degree in physics from the College of William and Mary

in 1979 and the Master of Engineering Physics degree from the University of Virginia

in 1980. In 1981, he joined NRL. He has worked on a wide range of semiconductor

materials and devices, specializing in GaN-based electronic devices.
Volume 51, Number 1 71

Paul B. Klein received a B.S. in physics from the Polytechnic Institute of Brooklyn in 1967 and an

M.S. from Yale University in 1968. He earned a Ph.D. from Yale University in 1975.

In 1975, he accepted an NRC postdoctoral associate, and in 1977 was hired permanently

at NRL. He specializes in optical characterization of semiconductor materials.

Jaime A. Freitas, Jr. is currently using a combination of defect-sensitive techniques to study

defect incorporation in GaN. Dr. Freitas received his M.Sc. and Ph.D. degrees from the

State University of Campinas (Sao Paulo, Brazil). He was an Associate Professor of

physics at the Federal University of Ceara, Brazil before coming to the NRL as a visiting

scientist in 1982. Dr. Freitas was a senior research physicist for Sachs/Freeman

Associates, Inc. from 1984 to 1996.

James C. Culbertson is a research scientist in the Electronic Materials Branch at NRL.

Dr. Culbertson received a B.S. in engineering physics from the University of California

at Berkeley in 1975 and a Ph.D. in physics in 1983. He was an NRC-NRL postdoctoral

associate in 1983 and was hired as a staff member in 1985.

Mohammad Fatemi received his M.Sc. degree in theoretical physics in 1965 and his Ph.D.

degree in physics in 1970 from the George Washington University. Since 1970, he has

been active at NRL in the X-ray and neutron diffraction and topographic studies of strainrelated phenomena in structural and electronic materials.
Dr. W.J. Choyke

Department of Physics

University of Pittsburgh

Pittsburgh, Pennsylvania

After receiving his Ph.D. in 1952 from Ohio State University, Jim Choyke joined

Westinghouse Research Laboratories in Pittsburgh, Pennsylvania. He became

an Adjunct Professor in the Physics Department of the University of Pittsburgh

in 1974, and in 1988, after early retirement from Westinghouse, became a

Research Professor of Physics. Currently, with his colleague Prof. Robert Devaty, he runs

the Large Bandgap Laboratory of the Department.

In 1955 he met Lyle Patrick who had just come from the University of Chicago, and this

led to a scientific collaboration on SiC that only ended with Patrick’s early retirement in the

mid-1970s. The first objective of this collaboration was to verify that all the common

polytypes of SiC were indirect semiconductors. Indeed, this work went on concurrently

with the famous study on the indirect transitions of Si at the Royal Radar Establishment in

Malvern, England.

The next effort was stimulated by the theoretical work of Murray Lampert, then at Berkeley,

on possible four-particle and three-particle exciton complexes in solids. Choyke and Patrick

used low temperature photoluminescence to yield wonderfully sharp spectral lines for SiC,

which led to a much better understanding of shallow impurity centers in SiC polytypes. The

Profiles in Science
Cover: Photograph of blue and blue/green LEDs courtesy of Cree Research, Inc.

Colin Wood

Office of Naval Research

From the Guest Editor . . .

aster, cheaper, smaller” is now

the standing order for satellite

and probe development at the

National Aeronautics and Space “F Administration.

development require electronics performance that can

only be delivered by wide bandgap semiconductors.

What is it about these newer semiconductors that

makes them so necessary and attractive? Large

bandgaps mean that leakage and avalanche onset by

generation-recombination currents are insignificant

at normal temperatures. With three-to-tenfold higher

thermal conductivities, heat dissipation is more

efficient, and dielectric constants are approximately

30% lower, thereby reducing parasitic capacitances.

Most importantly, however, is that breakdown voltages

are an order of magnitude higher, allowing higher

voltage swings.

Before these properties can be exploited, however,

several challenges remain to be solved in the area of

materials.

• Silicon carbide (SiC), although it can be grown

in bulk crystals at atmospheric pressure, requires

enormous temperatures, which in turn limits the

size of crystals and leads to many large defects,

known as micropipes.

• GaN bulk growth is not yet a reality, owing to

the enormous rate of decomposition above

1,000 °C. Other potential solutions, such as

compliant substrates, lateral growth, and pendeoepitaxial overgrowth (that is, epitaxy sideways

over a substrate, which reduces defect density

by five orders of magnitude) are under ONRfunded investigation.

• Acceptor dopants are “deep” in SiC, deeper still

in GaN, and too deep in AlN. Thus, such

components as lasers and bipolar transistors,

which depend on low resistivity p-type films, are

most challenging to produce. Solutions to this

particular challenge may come from recent ONR

initiatives in spontaneous and piezo-polarizationinduced hole conduction.

World experts address these and other related wide

bandgap materials and device challenges in this

edition of Naval Research Reviews.

If only life were that simple for Department of

Defense research and development. Although we

prefer “more affordable” to “cheaper,” comparatives

such as “lighter,” “more powerful,” “more efficient,”

“broader band,” “higher frequency,” “lower noise,”

“stealthier,” and “higher resolution” abound in the

modern Navy’s research, development, and

engineering (RD&E) lexicon. Defense electronics

RD&E does not pass muster unless it fits these, and

many more, criteria.

In truth, we in the DoD should use mostly

superlatives, e.g., most efficient, for that is what we

mean: only the best is good enough for the U.S.

military. ONR’s function is to ensure that the best is

what they get.

More than 80% of command, control,

communications, computers, and intelligence (C4I)

system electronics performance is achieved with

silicon (Si; specifically, complementary metal-oxide

on silicon (CMOS)). Similarly, the highest

performance microwave, millimeter wave, and optoelectronic functions depend upon cubic III-V

semiconductors: GaAs, InP, and their alloys with

AlAs, etc. Development, perfection, and insertion of

these compounds were largely the results of ONR’s

(especially Max N. Yoder’s) forward thinking and

oversight, and of sustained investment by ONR and

DARPA over the last 30 years.

For more than 20 years, Max N. Yoder has been

pushing the wide bandgap boat out into a very

skeptical sea. Although the true potential of wide

bandgap semiconductors such as GaN and SiC was

recognized more than 50 years ago, until recently

the materials problems were thought insurmountable.

Multifunction electromagnetic systems under
Naval Research Reviews, a theme-issue quarterly journal,

publishes articles about Department of the Navy Science

and Technology programs conducted by academia,

government laboratories, for-profit and nonprofit

organizations, and industry.

The journal is available on the World Wide Web at

http://www.onr.navy.mil/onr/pubs.htm

Chief of Naval Research

RADM Paul G. Gaffney, II, USN

Executive Director and Technical Director

Dr. Fred Saalfeld

Vice Chief of Naval Research

BGen Timothy Donovan, USMC

Director, Corporate Communications

Liane Young

Publications Coordinator

Cynthia Nishikawa

Scientific Editor

Colin Wood

Managing Editor*

Kathy Parrish

Technical Writer/Editor*

Saul Oresky

Design and DTP*

Jonna Atkinson

Donna Gloystein

Jan D. Morrow

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