Supplementary Material for “High-quality II-VI Films Grown on Amorphous
Substrates using Tunable Tetradymite Templates”
V.A. Stoica,1,a L. Endicott,1 H.H. Shen,2,b W. Liu,1 K. Sun,2 C. Uher,1 and R. Clarke1
1
Department of Physics, University of Michigan, Ann Arbor, Michigan 48109, USA
2
Department of Materials Science and Engineering, University of Michigan, Ann Arbor, Michigan 48109, USA
_____________________________
a)
Author to whom correspondence should be addressed. Electronic mail: vstoica@umich.edu
b)
Permanent address: School of Physical Electronics, University of Electronic Science and Technology of China, Chengdu 610054, China.
A. Substrate Cleaning
Before insertion inside the MBE chamber, the fused silica substrates were initially sonicated in an
ambient temperature acetone bath for 1 h, followed by H2SO4 (80%)/H2O2(20%) acid solution cleaning.
This solution was prepared by adding H2O2 to H2SO4 while boiling at 80°C. The fused silica substrate is
left in the boiling acid solution for 30 minutes, followed by rinsing in boiling de-ionized water and
drying with pressurized air.
The polyimide Kapton HNTM (DuPont) substrates were first cut to the desired size, followed by
immersion in boiling methanol several times, drying with pressurized air and immediate loading into the
MBE chamber.
The single crystalline substrates from MTI Corporation were cleaned as follows: Al2O3 (0001) and
MgAl2O4 (111) substrate cleaning followed by the RCA cleaning procedure,1 modified in this case by
omitting the HF-solution etching. GaP (111) and Ge (111) were cleaned simply by immersion in boiling
de-ionized water and then dried with pressurized air. The water soluble MgO (111) substrate was
cleaned in boiling de-ionized methanol and then dried with pressurized air.
B. MBE Deposition Procedures and in situ RHEED Studies
Before the deposition runs, the sample holder thermocouple was calibrated using another
thermocouple placed on the substrate surface, inside the molecular beam epitaxy (MBE) chamber. After
loading the substrates inside the MBE chamber, the fused silica and Kapton substrates were outgassed at
320°C for ~1 h, before lowering the temperature for deposition. The use of higher outgassing
temperatures for fused silica did not visibly influence the growth results. On the other hand, the single
crystal substrates (Al2O3, MgAl2O4, GaP, Ge, and MgO) were outgassed at higher temperatures of 6001
650°C before deposition, which was important for obtaining clean and atomically abrupt surfaces that
optimize epitaxial growth of II-VI/V-VI heterostructures. Growth rates are calibrated using both a quartz
crystal microbalance and RHEED intensity oscillations recorded during the real-time epitaxial growth of
Bi2Te3 and SbxBi2-xTe3 on Al2O3 (0001) substrates. The resulting film thicknesses were verified with
atomic force microscopy (AFM) at an abrupt film edge step and lastly with acoustic pulse measurements
based on ASOPS measurements.2 The individual growth rates of the Bi2Te3 and Sb2Te3 components
during the growth of ternary SbxBi2-xTe3 solid solutions are used to achieve the desired compositions for
lattice matching.
For the deposition of ternary SbxBi2-xTe3 buffer layers, a substrate temperature was selected in the
range from room temperature (RT) to 320°C, depending on the particular substrate being used. For the
case of fused silica and Kapton substrates, we observed that growing SbxBi2-xTe3 layers at > 200°C
results in film discontinuities, while films become continuous with lower growth temperatures. Nearly
optimal growth conditions that assure good substrate coverage, film adherence and crystallization are
realized when growing first a polycrystalline Bi2Te3 nucleation layer at ~180°C. In the next step, a
highly-oriented, c-plane textured film is obtained when the Bi2Te3 nucleation layer is annealed at 310 ±
10°C. This post-deposition annealing (5-90 minutes) promotes film crystallization and surface flatness
and is performed by exposing the surface to Te flux, which avoids surface stoichiometry changes due to
onset of buffer layer sublimation at > 320°C. The c-plane film texture is maintained through the
overgrowth of additional SbxBi2-xTe3 layers at ~ 250°C. In the final step, the deposition of CdSe and
ZnTe is performed at 300°C.
To characterize the quality of the crystalline surface of the buffer layer at each stage of the
fabrication, we recorded the corresponding RHEED patterns using a KSA -400 data acquisition and
analysis system. For example, in Fig. S1, we show a series of RHEED patterns monitoring the evolution
of the film crystallinity throughout the heterostructure fabrication on fused silica substrate. The RHEED
image from Fig. S1 (a) shows the characteristic diffuse scattering pattern for an amorphous substrate
surface; no diffraction features are observed, as expected. Next, Fig. S1 (b) shows the typical diffraction
pattern for a 20 nm Bi2Te3 buffer layer after deposition at 180°C. A composite diffraction pattern is
observed, where the vertical streaks are due to c-plane oriented grains and the additional diffraction
rings are due to randomly oriented grains. The dominance of diffraction streaks indicates that most of
the grains are c-plane oriented at this Bi2Te3 layer thickness. Further annealing of the buffer layer at
310°C produces entirely c-plane oriented grains and leads to the complete suppression of the randomly
oriented grain contributions. Therefore, the post-deposition annealing is very important to obtain the
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formation of highly-oriented, c-plane texture for the tetradymite buffer layer. After overgrowth of
additional SbxBi2-xTe3 overlayers at 250°C, the c-plane film texture is maintained throughout the
deposition. The corresponding RHEED pattern in Fig. S1 (c) shows the desired streaky pattern for a
Sb1.5Bi0.5Te3 buffer layer recorded just prior to the growth of the CdSe film. Furthermore, the
overgrowth of the CdSe film is seen to maintain a streaky RHEED pattern, as shown in Fig. S1 (d),
demonstrating that optimization of the buffer layer fabrication is key for controlling the highly-oriented
growth of CdSe films on amorphous substrates.
Figure S1. RHEED patterns obtained during progression of heterostructure growth: (a) bare fused silica;
(b) as grown Bi2Te3 (20 nm); (c) annealed Sb1.5Bi0.5Te3 (20 nm); (d) CdSe (20 nm).
Additional insight on the deposition of II-VI/V-VI heterostructures is obtained from growth
experiments on single-crystalline substrates, where the heterostructures are epitaxial, i.e. the buffer layer
and CdSe wurtzite film strictly follows the crystallographic orientation of the substrate. Fig. S2 (a) and
(b) shows the RHEED patterns acquired along the characteristic in-plane azimuths for a
CdSe/Sb1.24Te0.76/Al2O3 (0001) epitaxial heterostructure. Due to the layered structure of the tetradymite
buffer layer, one might expect to observe abrupt van der Waals (vdW) growth3 of II-VI overlayers.
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However, this is not the case. In fact, when using tetradymite layers with variable composition, we
observe a gradual relaxation of the in-plane lattice parameter of the CdSe films, with characteristic
thicknesses for relaxation depending on the degree of mismatch (few nm to hundreds of nm), a behavior
observed also on fused silica substrates. The vdW epitaxy does not occur for CdSe grown on tetradymite
buffer layers, probably due to strong interlayer bonding at the CdSe/Bi2Te3 interface. For ternary buffer
layer compositions close to x=1.6-1.8, no in-plane lattice parameter relaxation is observed for CdSe,
indicating that close matching with the buffer layer is achieved in this composition range. The RHEED
images obtained for an epitaxial zincblende ZnTe layer grown on Ge (111) using Bi 2Te3 buffer layer are
shown in Fig. S2 (c-d). These results show that our technique is applicable to a multitude of II-VI
compounds, with both cubic and hexagonal structures.
It is important to note that the growth mode of CdSe on non-crystalline and crystalline substrates
using tetradymite layers does not belong to the vdW epitaxy class, supporting our observations that
CdSe film improvement results from the compositional tuning of interfacial lattice matching.


and [1 1 00] azimuthal
Figure S2. (a) and (b) CdSe/Sb1.24Bi0.76Te3/Al2O3 (0001) along [11 2 0]


directions, respectively; (c) and (d) ZnTe/Bi2Te3/Ge (111) along [0 1 1] and [1 2 1] azimuthal directions,
respectively.
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C. AFM Studies of Surface Morphology
AFM was used to study the morphology of CdSe films grown on fused silica using Sb xBi2-xTe3
tetradymite buffer layers with varying composition. In Fig. S3 (a-b) we show the surface CdSe
morphologies resulting from growth on lattice-mismatched Bi2Te3 and nearly lattice matched Sb2Te3
buffer layers, respectively. A large difference in grain size, surface roughness and pit density is observed
between these two cases. Fig. S3 (c) shows the scaling between composition and surface roughness with
different pit densities. We observe that interfacial lattice-matching at x=1.6-1.8 favors much better
surface morphology, supporting our approach for the II-VI/V-VI heterostructure design.
Figure S3. AFM of (a) CdSe/Bi2Te3/fused silica and (b) CdSe/Sb2Te3/Bi2Te3/fused silica. (c) RMS
roughness dependence on buffer layer composition, where the uncertainty is the symbol size. The
growth on Sb2Te3 compared to Bi2Te3 reduces the pit density by an order of magnitude.
D. Estimation of Carrier Diffusivity Lengths
To estimate carrier diffusion lengths in our films, we have used an optical non-contact technique
based on femtosecond fiber lasers and asynchronous optical sampling (ASOPS),12 which is adapted here
to transmission geometry. A femtosecond (fs) pump beam (~150 fs duration) at 519 nm and another fs
probe beam at 1550 nm are both focused on the film surface at normal incidence. The probe
transmission through the sample is monitored as a function of time delay relative to the pump excitation.
The resulting time resolved transmission (TRT) of the probe beam is recorded as a function of pump
intensity strength, while the corresponding transient carrier densities are estimated from the absorption
and beam spot size on the sample (~4 μm) of the pump laser.
Fig. S4 (a) shows the carrier density dependence of TRT for a CdSe/Sb1.68Bi0.32Te3/Al2O3 (0001)
heterostructure. It can be seen that the TRT relaxation time is faster at larger carrier densities, which is
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characteristic behavior for all CdSe samples grown on glass and for bulk CdSe single crystals. The TRT
relaxation time dependence on carrier density indicates that we are probing charge carrier dynamics
instead of thermal effects. This is because the thermal relaxation in both CdSe and the buffer layer is
expected not to depend on carrier densities, while the carrier dynamics often shows carrier density
dependence through interplay between inter-carrier scattering, trapping, and impurity or phonon
scattering. The TRT sensitivity to carrier dynamics in CdSe samples is also confirmed by measuring,
separately, the TRT response of buffer layers with variable thickness in the same experimental
geometry. We observe that for time delays > 3 ps, the TRT response is dominated by the carrier
dynamics in the CdSe layer, while the buffer layer contribution to TRT is very weak on this timescale.
This is because the carrier relaxation time in the buffer layer is short (~ 1ps), while the TRT sensitivity
to thermal excitations and their relaxation is negligible at our probe laser wavelength of 1550 nm.
Figure S4 (a) Pump power dependency of carrier relaxation probed by TRT in CdSe on buffer layer
deposited on fused silica; (b) the ratio between TRT intensities acquired with displaced and centered
pump and probe spots for the CdSe single crystal (see text).
TRT traces are also acquired as function of spatial displacement relative to the pump beam position
along the sample surface. This geometry can be used to monitor the spatial distribution of carriers as a
function of time after pump excitation, particularly the carrier diffusion dynamics. Because the pump
spot intensity profile is Gaussian, the generated carrier density along the sample surface has the same
spatial profile, driving the carrier diffusion. Moreover, immediately after the laser pulse but before
recombination has occurred, the generated carriers can diffuse away from the pump beam center and
along the sample surface. The diffusion of carriers is driven by the local carrier density gradient along
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the sample surface, which follows the Gaussian excitation profile. A probe displacement with respect to
the pump spot position can in turn be used to monitor carrier diffusion.4 In this approach, it is assumed
that the relaxation time of carriers is the same at all sample locations within the spatial profile of the
pump beam. We observe that this assumption is not valid in our samples, since the relaxation time is
dependent on the average pump power as we have shown in Fig. S4 (a). Consequently, a different
technique is needed to account for the spatial dependence of carrier relaxation time. To address this, we
perform first a TRT measurement with probe spot displaced from the pump spot center. A smaller initial
carrier density is probed when using probe displacement from the center of the pump spot. Nevertheless,
the initial amplitude of the TRT signal correlates with the local carrier density. Next, we perform a
complimentary measurement with centered pump and probe spots while adjusting the pump beam
intensity to match the TRT amplitude for the displaced spot case. This procedure assures that the initial
average carrier density within the probe spatial spot is the same for the two complimentary TRT
measurements, which significantly suppresses the TRT relaxation time variation with the carrier density.
We then divide the two TRT complimentary measurements to monitor the diffusion component of
carrier dynamics. In Fig. S4 (b) we show the resulting ratio of TRT signals for the CdSe single crystal
and the CdSe film grown on nearly-matched Sb1.68Bi0.32Te3 buffer layer grown on fused silica substrate.
The ratio of TRT signals grows quickly within the first 100 ps, due to carrier diffusion, followed by a
slower growth on a longer time scale. For the case of the CdSe film sample the TRT signal is modulated
by acoustic dynamics inside the buffer layer (not shown), but the rise of slowly-varying TRT
background is still similar to the CdSe bulk sample. This implies that we probe the carrier diffusivity
dynamics at the early time delays in both samples, while at longer time delays the diffusion effects are
negligible. We assign this behavior to free electron diffusion at early time delays, followed by formation
of excitons at longer time delays, which are characterized by much smaller diffusivities. The temporal
dependency of the pump spot Gaussian width, w, due to carrier diffusion can be obtained from:

I (t )
e
I0
d 2  1
1

2  w 2 w0 2




, where I(t) is the TRT corresponding to a displacement, d, between the pump and
probe spots on the sample, I0 is the TRT measured at the center of the pump spot, and w0 is the Gaussian
width before the pump excitation. We can then use the w2 = 11.08 Dt + w02 relation for estimating the
carrier diffusion constant.4
This approach is applied within the first 100 ps corresponding to a time interval preceding the carrier
relaxation and exciton formation. Accordingly, we estimate a carrier diffusivity of ~5 cm 2/s for the CdSe
film and ~20 cm2/s for the CdSe single crystal. Using the TRT relaxation times and these diffusivities,
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we then estimated the carrier diffusion lengths of 500 nm for the CdSe film and 2.5 μm for the CdSe
single crystal.
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V. A. Stoica, Y. M. Sheu, D. A. Reis, and R. Clarke, Opt. Express 16, 2322 (2008).
3
A. Koma, Thin Solid Films 216, 72 (1992).
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L. M. Smith, D. R. Wake, J. P. Wolfe, D. Levi, M. V. Klein, J. Klem, T. Henderson, and H. Morkoc,
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8