Deep ultraviolet photodetectors grown by gas source molecular beam
epitaxy on sapphire and AlGaN/sapphire substrates
M. Holtz*a, V. Kuryatkova, D. Y. Songa, B. Borisova, S. Nikishina, A. Usikovb, V. Dmitrievb,
Yu. Kudryavtsevc, and R. Asomozac
a
Nano Tech Center, Texas Tech University, Lubbock, Texas, 79409, USA
b
TDI, Inc., 12214 Plum Orchard Dr., Silver Spring, MD 20904 USA
c
SIMS Laboratory of SEES, CINVESTAV, Mexico D.F. 07300, Mexico
ABSTRACT
Optically-based chemical and biological sensors require optoelectronic devices with specific emission and detection
wavelength ranges. Semiconductor optoelectronic devices applicable to this sensing are of particular interest due to their
low power consumption, compact size, long lifetime, and low cost. We report the electrical and optical properties of
deep UV p-i-n photodiodes (PDs) based on short period superlattices (SPSLs) of AlN/AlGaN. All device and test
structures are grown by gas source molecular beam epitaxy with ammonia on sapphire and AlGaN/sapphire substrates.
AlGaN/sapphire substrates were grown by stress controlled hydride vapor phase epitaxy (HVPE). The cutoff
wavelength of PDs based on these SPSLs can be varied from 250 to 280 nm by changing the SPSL barrier/well
thickness ratio. For mesa diodes with 150 µm diameter we obtain extremely low dark leakage current of ~ 3 pA/cm2,
and high zero-bias resistance of ~ 6 × 1014 Ω. A cutoff wavelength of 247 nm is obtained for these devices with four
orders of magnitude rejection by 315 nm. We obtain a maximum responsivity of 60 mA/W.
Keywords: photodetector, solar blind, AlGaN, superlattice
1. INTRODUCTION
A solar blind photodetector (SBPD) is sensitive to light in the 200 to 280 nm ultraviolet (UV) range and insensitive to
visible light, particularly the solar blackbody spectrum [1]. Such photodetectors have broad applications including
space-based observational astronomy, sensors for defense, and environmental monitoring. In each case, the presence of
intense visible light is a limitation, particularly when attempting to detect UV radiation at low light levels for obtaining
good limits of detection.
One approach to SBPDs is to utilize a short-pass filter, with cutoff wavelength at ~ 280 nm, in conjunction with a
photomultiplier tube (PMT) as shown in Fig. 1 (upper). A hypothetical spectrum, composed of visible (e.g., blackbody)
radiation and the desired UV light, passes through the filter to reach the PMT. Disadvantages to this approach include
size and weight, the need for substantial additional electronics, and the low shock resistance of the PMT. An alternative
scheme utilizes solid-state photodetectors based on semiconductor materials which have energy gap at the desired
corresponding wavelength. This is depicted in Fig. 1 (lower), where the same general spectrum illuminates the detector.
Photons with energy lower than the semiconductor band gap (longer wavelengths) pass undetected through the sensor.
This approach has the advantages of requiring few components, is lightweight and compact, and is shock resistant.
Furthermore, it is possible to fabricate these in arrays for imaging – a configuration not realized using the PMT
approach.
In the UV, the premier semiconductor materials are based on AlGaN alloys or superlattices. By controlling composition,
the energy gap may be continuously varied from 3.4 (GaN) to 6.1 eV (AlN). One critical issue in growing high-quality
AlGaN materials remains the paucity of bulk substrates for any of the nitride semiconductors. Instead, sapphire, silicon
carbide, and silicon “foreign” substrates have been used with good success, although the dislocation density of the
*
Mark.Holtz@ttu.edu
Optically Based Biological and Chemical Detection for Defence III, edited by
John C. Carrano, Arturas Zukauskas, Proc. of SPIE Vol. 6398, 63980Z, (2006)
0277-786X/06/$15 · doi: 10.1117/12.689938
Proc. of SPIE Vol. 6398 63980Z-1
Fig. 1. Schematic illustration of two approaches for obtaining solar blind photodetection.
epitaxial materials grown on these substrates is still very high in comparison with the cubic column IV and III-V
compound semiconductors. AlGaN “template” substrates have recently been developed to replace these foreign
substrates.
Much progress has been made in developing AlGaN materials for optoelectronic devices. For example, when
considering AlGaN and sibling semiconductor InGaN, LEDs have been fabricated operating from the green to 210 nm
[2]. SBPDs require AlGaN with Al content ≥ 0.5. Such devices have been reported [1,3,4]. SBPDs based on AlGaInN
alloys have been grown primarily by metalorganic chemical vapor deposition (MOCVD) [5,6]. There are also several
reports of SBPDs grown using molecular beam epitaxy (MBE) [7-9]. Common problems to basing SBPDs on these
materials remain the ability to grow n, p, and i type materials, each at the needed composition, the high dislocation
density, and control over the contact formation. In this paper, we report our use of gas source MBE (GSMBE) for
growing short-period superlattices (SPSLs) with high average Al content, and progress toward a SBPD based on this
approach.
2. GROWTH OF AlGaN/AlN SUPERLATTICES
All samples discussed here were grown by GSMBE with ammonia on c-plane sapphire using thin AlN buffer layers or ~
200 nm thick AlxGa1-xN/sapphire templates (x ranging from 0.7 to 0.9) grown by hydride vapor phase epitaxy (HVPE).
The SPSLs examined have periods ranging from 1.25 to 2.75 nm. The barrier material is pure AlN and the well material
is AlxGa1-xN with x ~ 0.08. Growth details have been previously published [10-15].
Proc. of SPIE Vol. 6398 63980Z-2
- 01
00
01
(a)
(b)
Fig. 2. Evolution of RHEED patterns for different stages of GSMBE: (a) formation of 3D AlN islands, (b) 2D
growth of AlN buffer with the corresponding streaky pattern.
Reflection high energy electron diffraction (RHEED) was used for in situ monitoring of the growth mode. A thin AlN
buffer layer was grown on the sapphire substrate for the SPSLs described here. The process was started by exposing the
surface of sapphire or AlN/sapphire templates to ammonia for 15 minutes. Note that incomplete nitridation of sapphire
and low lateral growth rate of Al- and N-face domains results in formation and vertical propagation of inversion
domains through the entire structure. SPSLs with high density of these domains exhibit inferior electrical properties and
cannot be used in the preparation of devices. Note that exposure of the HVPE AlGaN/sapphire templates is also needed
to remove the residual surface oxide. Incomplete nitridation in this case results in elevated dislocation density in the
SPSLs similar to what we obtained for the growth on bulk AlN substrates [15]. We found that nitridation of AlGaN
template substrates is not similar to nitridation of sapphire and bulk AlN substrates. Although further work is required
to lower dislocation densities, we did find that template substrates permit growth of SPSLs directly following
nitridation.
Focusing on the growth of SPSLs on sapphire, the nitridation was immediately followed by exposing the surface to Al
flux for 3 s, followed by 3 s long exposure to ammonia (Al flux off). This alternating exposure process, repeated 5 to 10
times, facilitates formation of oriented AlN islands on sapphire and helps to reach 2D growth conditions on the template
surface. The growth of AlN was then started by turning on ammonia and Al together. Spots observed at this point in the
RHEED pattern, Fig. 2a, indicated formation of oriented three-dimensional (3D) islands of AlN. The formation of the (1
× 1) surface structure is seen by RHEED, Fig. 2b, after depositing 5 to 7 nm of AlN. A completed 2D growth mode is
Fig. 3. TEM cross-section of buffer layer and superlattice at the sapphire substrate.
Proc. of SPIE Vol. 6398 63980Z-3
seen by RHEED after about 40 nm growth. The SPSL growth was started after the AlN buffer layer completed the
transition to the 2D growth mode. Formation of a (2 × 2) surface structure could be seen after deposition of 20 to 50 nm
of SPSL. The (1 × 1) → (2 × 2) transition corresponds to the formation of a single domain SPSL with group-III face
polarity. When we grew on HVPE AlGaN/sapphire template substrates, the RHEED sequence was similar but a
presence of 3D growth was attributed to a residual native oxide. The rapid transition to 2D growth is of interest for two
reasons. First, it results in a decrease in the defect density, as shown in the transmission electron microscope (TEM)
cross section of Fig. 3. Second, the rapid transition is needed to assure crack-free samples. We used these SPSLs grown
on AlN buffer for studying the properties of the materials and for device fabrication.
3. X-RAY PROPERTIES OF AlGaN/AlN SUPERLATTICES
High-resolution X-ray diffraction (HRXRD) studies were performed using double and triple-axis alignment. A fourbounce Bartels Ge-220 monochromator was used for selecting Cu-Kα1 radiation. In order to observe the higher order
superlattice peaks and the (1124) asymmetric diffraction, the monochromator was replaced with a hybrid mirror. A
three-bounce Ge-220 analyzer crystal was used for high-resolution measurements. The 2θ-ω scans of the symmetric
(0002) reflections were employed to determine the SPSL period and the average AlN content, along with the respective
thicknesses of the well and barrier layers. The 2θ scans of the symmetric (0002) and asymmetric (1124) reflections
were used to accurately determine the strained lattice parameters, a and c, following the method described by Fewster et
al. [16]. Precise knowledge of both SPSL lattice parameters is important for accurately determining the stress, degree of
relaxation, and the average composition.
Close examination of the a and c lattice constants, and analysis using Vegard’s Law, shows the SPSLs to be primarily
strain relaxed. Layers of “bulk” AlGaN, with varying composition in this range and grown with comparably ~ 1 µm
total thickness, are likewise relaxed. Figure 4 shows 2θ-ω rocking curve for a SPSL with nominal period of 9
monolayers (ML) comprised of 6 ML AlN barriers and 3 ML Al0.08Ga0.92N wells. Simulation of the 0 and ±1 2θ-ω
diffraction using this ideal structure results in a poor match of the peak positions and line shapes. In order to adequately
model the observed HRXRD we find it necessary, in general, to include ML-level interface roughness and the presence
of an interfacial transition region between the barrier and well layers which has intermediate composition ~ 0.5 [17]. It
is possible that the barrier → well growth transition produces a different interface roughness and interfacial composition
than the well → barrier growth transition. This is because the Al0.08Ga0.92N is observed to result in extremely smooth
surfaces, based on RHEED, even after growing a few ML, while the AlN is difficult to coax into the 2D mode. The
resulting simulation is shown with the data in Fig. 4. In all SPSL cases studied it was necessary to include the interfacial
composition factor, while the presence of ML interface roughness varied from sample to sample. Ref. [17] contains a
detailed description.
Fig. 4. Long scan X-ray diffraction from a test superlattice grown with uniform properties. Solid: data. Dashed:
simulation. The narrow peaks are artifacts of the simulation setup, which is composed of block elements
having the desired period.
Proc. of SPIE Vol. 6398 63980Z-4
Fig. 5. Directly measured reflectance and transmittance spectra of a representative uniform superlattice. The
absorption coefficient is obtained from these two measurements.
4. OPTICAL PROPERTIES
In order to obtain the optical energy gaps of the SPSLs, we measured both reflectance and transmittance spectra at room
temperature. Figure 5 shows reflectance (R) at 5º external angle of incidence and transmittance (T) measurements at
normal incidence. The transmission is high below photon energy ~ 4.8 eV, above which it decreases as the SPSL
becomes opaque. The reflectance spectrum exhibits Fabry-Perot fringes in the transparent range. The energy at which
these fringes are quenched by the onset of high absorption allows estimation of the optical energy gap [18]. The fringes
also provide a rapid estimate of the thicknesses which are consistent with the scanning electron microscopy (SEM)
cross-sections. Thus, these two measurements allow us to obtain the optical gap. Using the method presented by Poruba,
et al. [19] we use reflectance and transmittance spectra are used to obtain the absorption coefficient (α) and index of
refraction. The approach takes into account multiple internal reflections and uses R and T values at each photon energy
(ħω) and the layer thickness from our SEM cross-sections. Self-consistent expressions are numerically solved to obtain
α(ħω).
For transitions across direct bandgap semiconductors, the absorption depends on photon energy according to
α ( hω ) ∝ hω − Eg
2
(1)
above the optical gap Eg. In Fig. 4 we graph α versus photon energy (lower panel), which is expected to rise linearly
above Eg. Extrapolating the linear fit of α2 vs. photon energy data to α2 = 0 provides an estimate for Eg within ± 0.05 eV.
Below the energy gap the index of refraction n ~ 2.3. A weak tail in α(ħω) just below Eg suggests an Urbach region
Proc. of SPIE Vol. 6398 63980Z-5
Fig. 6. Correlation between average composition of superlattices obtained from optical measurements and X-ray
diffraction. Period varies from ~ 1.25 to 2.25 nm. Dashed curve is 1:1 correspondence.
extending ~ 100 meV below the bandgap. Studies of the temperature dependence are needed to fully qualify this
proposed origin of the observed weak tail. We also obtain information on the value of α just above Eg and the steepness
of the absorption edge. We estimate an absorption region approximately 0.2 µm thick will satisfy the 1/e2 requirement.
We note also that a solar blind detector with nominal cutoff wavelength of 260 nm (Eg = 4.77 eV), for example, will
require an absorption region with average AlN mole fraction of ~ 0.60. In order for the surrounding layers to transmit
wavelengths above 240 nm, so that the SBPD is sensitive to a ~ 20 nm range (240 to 260 nm), the layers sandwiching
the absorption region must have a higher AlN mole fraction by ~ 0.15.
Energy gaps thus obtained are found to correlate well with the average SPSL mole fraction <x> obtained by HRXRD.
This agreement is illustrated in Fig. 5, where we graph <x> from the energy gap studies versus <x> from x-ray. For the
former, we use measured Eg and obtain x from
(2)
E g = xE AlN + (1 − x ) EGaN − x (1 − x )b
with bowing coefficient b = 1.0 eV [20]. The correlation between the values of <x> for these two measurements
suggests that the energy gap of our SPSLs is primarily determined by the average composition with minor effect of
particle confinement. For reference, the right-hand axis of Fig. 5 is the energy gap using Eq. (2) and <x> from x-ray
diffraction.
5. p-i-n SBPD DESIGN
The conclusion from our optical studies is that SPSLs with extremely short period have absorption properties similar to
AlGaN random alloys. This allows us to apply similar design concepts used for alloy-based UV photodetectors. The
double hetero structure (DHS) design, in principle, should demonstrate relatively higher external quantum efficiencies,
uniform spectral response, and lower diffusion current compare with homojunction based structures [21]. A schematic
cross-section of our DHS SBPD with cutoff near 250 nm is illustrated in Fig. 7. The structure consists of two main
parts. The first contains a thin, ~ 40 nm, AlN buffer grown on sappire or AlGaN/sapphire template substrates following
by undoped ~ 300 nm thick SPSL with an average xAlN content of ~ 0.7 and transparent to 240 nm. These AlN and
SPSL are required to mitigate dislocation density in the subsequent active layer of the device [15,22,23]. The second
part is composed of n-, i-, and p-type layers, each of AlN/Al0.08Ga0.92N SPSLs having different optical gaps and
thicknesses, comprising the photodetector part of the device. Using the dependences of Eg on SPSL period with 2 ML
and 3 ML thick Al0.08Ga0.92N wells, reported in [15,24], one could estimate the required period and well/barrier
thicknesses of these SPSLs, namely {~ 9 ML (~ 2.25 nm) and ~ 3 ML (~0.75 nm)/~ 6 ML (~1.5 nm)} and {~ 8 ML (~ 2
nm) and ~ 2 ML (~ 0.50 nm)/ ~ 6 ML (~ 1.5 nm)} for i- and p-(n-) regions, respectively. These layers were grown in the
2D growth mode as confirmed by RHEED. The optical bandgaps were controlled by SPSL design: ~ 4.96 eV (~250 nm)
in i-layer to ~ 5.17 eV (~240 nm) in n- and p-type layers. The SPSLs were either not intentionally doped (i-type) or
doped with Si derived from silane (n-type), or Mg derived from a conventional effusion cell (p-type). The final Mg-
Proc. of SPIE Vol. 6398 63980Z-6
Fig. 7. Cross-section of SBPD design. 1 – sapphire or AlGaN/sapphire substrate, 2 – AlN buffer, ~ 40 nm, 3 –
undoped SPSL, ~ 300 nm, 4 – n-type SPSL, Eg ~ 5.2 eV, ~ 300 nm, 5 – i-type, Eg ~ 5.0 eV, ~ 100 nm, 6 –
p-type SPSL, Eg = 5.2 eV, ~ 200 nm. The metal contacts are shown in gray.
doped well layer served as the contact material. This approach permits design of PDs which may be simultaneously
illuminated from front and back. The average concentrations of Si and Mg measured by secondary ion mass
spectrometry (SIMS) were ~ 7 × 1019 cm-3 and ~ 1 × 1020 cm-3, respectively. In parallel experiments we studied the
electrical properties of Si-doped (ND ~ 7 × 1019 cm-3) and Mg-doped (NA ~ 1 × 1020 cm-3) SPSLs with optical bandgap of
~ 5.2 eV. Room temperature Hall measurements yielded an electron concentration n ~ (3-5) × 1019 cm-3 and mobility of
~ 30 cm2/Vs for n-type layers and a hole concentration p ~ (6-8) × 1017 cm-3 and mobilities of 3-6 cm2/Vs for p-type
material. Hall measurements of unintentionally doped SPSLs, grown in separate experiments, exhibit electron
concentrations of (1-3) × 1016 cm-3. The optimal thickness of i-region is determined to be ~ 100 nm [9]. Capacitancevoltage (CV) measurements of the structure shown in Fig. xxx reveal the width of the depletion region to be ~ 150 nm
at zero bias. The SIMS data, not shown here, illustrates substantial diffusion of the Si through the i region extending
into the p-type layer. Although the impurity concentrations of Mg and Si in the p-type region differ by more than two
orders of magnitude, activation of Si donors is more efficient than Mg acceptors. This could be one of the reasons for
the high series resistance observed under forward bias for our p-i-n structure. The resulting non-uniform electric field
across the i region could result in a diminished internal quantum efficiency of the PD. We are currently experimenting
with growth procedures aimed at reducing dopant penetration into the i region.
Mesa photodiodes were prepared using standard processing. The 150 µm diameter mesas were plasma etched with Cl2
using 100-nm thick Ni hard masks. The mesa walls were then passivated in oxygen plasma. Details of the plasma
process were described elsewhere [25,26]. P- and n-type contacts were processed separately after removal of the Ni
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Fig. 8. Current-voltage characteristic of p-i-n photodiode structure in dark (solid) and under 247 nm illumination.
Inset shows dV/dI under dark conditions.
Proc. of SPIE Vol. 6398 63980Z-7
mask. N-type contacts were made using Ti/Al/Ti/Au with the total thickness of ~ 200 nm. Contacts to p-type SPSL were
made using 50 nm of Ni followed by 70 nm of Au. For n- and p-type contacts the lowest specific contact resistance of ~
6 x 10-5 Ω·cm2 and ~ 2 x 10-2 Ω·cm2 were obtained after rapid annealing in N2 ambient at 700 oC and 550 oC,
respectively.
6. CHARACTERISTICS OF THE 247 nm SBPD
A representative current-voltage (I-V) characteristic of our PD structure is plotted in Fig. 8. Low dark leakage current of
~ 0.5 fA is measured at near zero bias. This corresponds to leakage current density of ~ 3 pA/cm2 for the 150 µm
diameter mesa device. The reverse dark current remains below 0.5 µA/cm2 up to a bias of – 10 V and scales with the
device area demonstrating the absence of significant contributions to leakage current from the device mesa sidewalls.
These extremely low values of dark current are comparable to results demonstrated for AlGaN solar blind PDs grown
by MOCVD (e.g., [6]). The dark current at lower bias, up to – 1.5 V, does not significantly depend on applied voltage
and exhibits a slope of ~ 1 to 1.3 in the log-log scale. This low field leakage is due to tunneling of electrons (holes) from
the valence (conduction) band to the conduction (valence) band. The tunneling current is expected to depend on internal
electric fields.
At reverse bias beyond – 3 V the dark current increases with exponent ranging from 3 to 6 in the log-log scale in
different devices and wafers. Field-assisted thermal ionization of carriers from traps in the depletion region or
combination of hopping conductivity and Poole-Frenkel effect are likely mechanisms for this conductivity [27]. The
range of slopes is typical of AlGaN PDs at high reverse bias [23], and the reverse current correlates with high defect
density in structures grown on sapphire. The average density of threading dislocations in our samples, estimated from
X-ray measurements, is ~ 109 cm-2. This is comparable to that reported for AlGaN grown by MOCVD over thin AlN
buffer on sapphire [1,3,4].
From dV/dI (inset to Fig. 8) we calculate zero-bias PD resistance of ~ 6 × 1014 Ω, resulting in the area-resistance
product RoA ~ 1.2 x 1011 Ω cm2. The PD exhibits excellent monochromatic photoresponse at 247 nm, with I247nm/Idark ~
3 x 106 at zero bias.
Responsivity (mA / W)
Spectral responsivity under reverse bias is shown in Fig. 9. The peak response was obtained at a wavelength of ~ 247
nm, in excellent agreement with the expected optical bandgap of ~ 5 eV for the i-region of the PD. The device shows
excellent rejection of visible radiation. Cutoff at 247 nm is obtained for these devices, with three orders of magnitude
decrease in responsivity from 247 to 265 nm and four orders of magnitude rejection by 315 nm. This compares well
with the rejection ratios in state-of-the-art p-i-n PDs with cut off at longer wavelength [1,3-6]. The sub-structure
observed from 265 to 315 nm in the photoresponse is not understood at this time. It may be related to broadening of the
absorption edge due to partial monolayer thickness fluctuations in the well and barrier thickness and formation of the
Responsivity (A / W)
-2
10
-4
10
10
60
40
20
0
5
10
Reverse bias (V)
-6
240
260
280
300
320
Wavelength (nm)
Fig. 9. Spectral Responsivity showing cutoff at 247 nm and decreased sensitivity by three orders of magnitude
above 280 nm. Inset shows device linearity with 247 nm illumination.
Proc. of SPIE Vol. 6398 63980Z-8
well/barrier interfacial layers with intermediate AlN content [17]. Furthermore, there is little information regarding the
effect of intentional doping on the optical absorption properties of AlGaN, which may result in an Urbach tail, and
affect the device efficiency and spectral response. A study of the temperature dependence of the PD spectral response is
needed to examine this hypothesis.
At zero bias the responsivity of 16 mA/W was measured with illumination at 247 nm. This corresponds to an external
quantum efficiency of ~ 12.5%. The responsivity linearly increases with reverse bias, as shown in inset of Fig. 9, and
reaches its maximum value of 62 mA/W at – 10 V. This corresponds to external quantum efficiency of 30%. This is less
than the highest responsivity reported in solar-blind p-i-n PD of ~ 72% at 282 nm [23] grown by MOCVD. Optimizing
the dopant distribution throughout the device, but particularly in the i-region, should improve the external quantum
efficiency of our PDs. Nevertheless, using the RoA determined above we calculate the thermal noise limited specific
detectivity, at zero-bias, of D*=4.5x1013 cmHz1/2W-1.
7. SUMMARY
In summary, we have successfully grown SPSLs of AlGaN/AlN on select substrates. The materials show excellent
morphological properties and optical properties consistent with alloys of similar average composition. We have
fabricated p-i-n SBPDs using mesa etching. The cutoff wavelength of these devices varies with the SPSL characteristics
(energy gap), and we have obtained devices with cutoff at 247 nm and over three orders of magnitude rejection at and
above 280 nm. Electrical measurements show excellent dark current and good linearity.
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