- Email: [email protected]
Step graded buffer for (110) InSb quantum wells grown by molecular beam epitaxy
Author's Accepted Manuscript
Step graded buffer for (110) InSb quantum Wells grown by molecular beam epitaxy Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno, Brian R. Bennett
www.elsevier.com/locate/jcrysgro
PII: DOI: Reference:
S0022-0248(14)00501-6 http://dx.doi.org/10.1016/j.jcrysgro.2014.07.014 CRYS22337
To appear in:
Journal of Crystal Growth
Received date: 8 April 2014 Revised date: 1 July 2014 Accepted date: 6 July 2014 Cite this article as: Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno, Brian R. Bennett, Step graded buffer for (110) InSb quantum Wells grown by molecular beam epitaxy, Journal of Crystal Growth, http://dx.doi.org/ 10.1016/j.jcrysgro.2014.07.014 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Step Graded Buffer for (110) InSb Quantum Wells grown by Molecular Beam Epitaxy Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno and Brian R. Bennett Electronic Science and Technology Division, Naval Research Laboratory, Washington, DC Abstract We report on a two step buffer layer preparation for the growth of InSb quantum wells on a (110) GaAs surface. At each buffer layer step, layer conditions were optimized to produce smooth surfaces compatible with InSb quantum wells. Through varying growth rate, group V/III flux ratio, substrate temperature, and the addition of in situ annealing, we are able to grow In0.85Al0.15Sb on a GaAs substrate with an RMS surface roughness of approximately 2 nm. Surface morphology and cross-sectional transmission electron microscopy (TEM) were analyzed to understand the formation of threading dislocations, inclusions and dislocation filtering. This work presents an initial study for the growth of large lattice mismatched III-V materials on the (110) surface. Keywords: A3. molecular beam epitaxy, B2. Semiconducting III-V materials, B1. Antimonides, A1. Surfaces, A1. Crystal morphology 1. Introduction Understanding the details needed to grow semiconductors on a range of different substrate orientations will enable the fabrication of increasingly complex III-V heterostructures for advanced technologies. III-V growth by molecular beam epitaxy (MBE) has traditionally focused on (100) substrates in part due to the wide range of growth conditions which result in excellent epitaxial quality. Non (100) surface growth is also needed to take advantage of the anisotropic properties of the III-V semiconductors. Despite being the natural cleavage plane in 1
the zinc-blende structure, growth on the (110) surface has often been overlooked. Growth on the (110) surface could allow the development of new devices based on strong spin orbital coupling with applications to Majorana Fermions and spin-based devices due to the electron spin relaxation times of several nanoseconds at room temperature.1,2 The results of this work will also lead to a better understanding of cleaved edge overgrowth and to the overgrowth on patterned semiconductors containing edges with (110) surfaces that are useful for non-planar structures and devices.3 Further developments in the fabrication of device structures on the (110) surface requires a greater understanding of the growth processes and more specifically, the relationship between the growth conditions and the surface morphology of the films. Though the benefit of the (110) surface has been known for many years, little work has been done beyond the growth of InAs and GaAs/AlGaAs (110) quantum-well structures.4,5,6,7 The nature of the (110) surface, being non-polar as opposed to the standard (100) polar surface, leads to complications arising from the low incorporation rate of group V atoms onto the surface of the film as described by the work of Tok et al.
They have shown that the arsenic
incorporation kinetics vary greatly between the (100) and (110) surfaces for GaAs.8 This arises from the (110) surface having a single dangling adsorption site as opposed to the two dangling bonds found on the (100) surface. Therefore, the shorter residence time on the (110) surface for the group V means that its availability for incorporation is reduced. An increase in growth temperature will have a strong influence on the arsenic desorption rate and therefore the temperature of growth is of utmost importance. As a result, the general rule is that the growth on (110) requires higher group V fluxes and lower temperatures to allow for group V incorporation.9 Due to the orientation of the growth direction, two strain relief mechanisms are observed in the two orthogonal directions [1 1 0] and [001] which lead to strain relief. 60o 2
dislocations on the (110) mismatched growth have also been responsible for observable tilting of the surface which further complicates the growth as opposed to growths on the traditional (100) surface.10 A wide variety of growth regimes exist on the (110) surface particularly due to the difficulty of group V incorporation.
A systematic study on the effects of the deposition
conditions on the surface morphology is required to define the conditions under which highquality films can be produced. Antimonide based semiconductors are candidates for high-speed, low-power electronic devices,11 and have the highest electron g-factor amongst all III-V semiconductors.12 However, we are not aware of any systematic study on the growth of buffer layers for (110) antimonides on semi-insulating wafers, namely (110) GaAs. In this work, we investigated the growth conditions for reducing surface roughness using a two step buffer layer process to accommodate the 13.8% lattice mismatch between the GaAs substrate and an In0.85Al0.15Sb layer. We varied the flux rate and growth temperature at each buffer layer and determined surface roughness and morphology with atomic force microscopy (AFM). Each buffer layer step was investigated until the lowest roughness was achieved within our test parameters (as measured by AFM) before using it as a buffer for subsequent layer growths. The smoothest films were grown at low substrate temperatures, with the use of multiple in situ annealing steps and the careful control of group V flux rates. We fabricated and characterized an active QW structure after the optimized buffer layer was determined. 2. Experimental Procedures Samples were grown in a Riber 32P MBE system using cracked As2 and Sb2 sources on semi-insulating GaAs (110) on-axis substrates. Prior to growth, the native oxide layer on the
3
surface was removed by heating the substrate to 620oC for 10 minutes under an As flux. Temperature was measured with a transmission spectroscopy system to measure the band edge of GaAs, using the substrate heater as a light source. After removing the native oxide layer, the (110) surface displayed a (1 x 1) streaked reflection high-energy electron diffraction (RHEED) pattern.
Growth rates were obtained by RHEED oscillations on (100) GaAs wafers prior to
deposition. The two step graded buffer layer structure is shown by the schematic diagram in Fig.1. The structure of this paper is the optimization of the heteroepitaxial Al0.91Ga0.09Sb on GaAs (110) followed by the optimization of the In0.85Al0.15Sb (110) layer on top of the optimized Al0.91Ga0.09Sb layer. The effect of in-situ annealing on the surface morphology of the buffer layers is also investigated. We varied substrate temperatures (320oC to 440oC) and fluxes in order to optimize growth conditions. In situ annealing was performed under group V flux. The surface was characterized by atomic force microscopy (AFM) to determine the surface morphology as well as high resolution X-ray diffraction using a (022) Ge monochromater Cu K beam to determine the lattice constant and composition. After the optimized buffer layer structure was completed, a quantum well layer was deposited as a proof of concept and the electronic states were probed using photoluminescence. 3. Results and Discussion Previous to the metamorphic growth on (110) GaAs, a homoepitaxial smoothing layer was deposited at a thickness of 30nm. Standard MBE growth was used in the deposition of the smoothening layer with an optimized surface roughness of approximately 0.3nm being obtained. Homoepitaxial growth of GaAs has been studied in depth elsewhere. 8,13,14,15 Our optimized parameters consisted of a As/Ga beam-equivalent pressure (BEP) 138 (as determined by flux gauge measurements) and a deposition rate of 0.15ML/s at a temperature of 440oC. It should be 4
noted that there are other methods to obtain atomic-terraced homoepitaxial (110) GaAs such as migration enhanced epitaxy (MEE). In the MEE method, Group III atoms are supplied on the surface under Group-V-free atmosphere or a very low Group V pressure allowing the Group III atoms to migrate for long distances.16,17,18 A.
Optimization of AlGaSb (110) buffer layer growth
In order to accommodate the large lattice mismatch between the GaAs substrate and the InSb Quantum Well (~14.7%), an intermediate layer was used on top of the prepared GaAs surface discussed in the previous section. Rather than going directly to an In0.85Al0.15Sb buffer layer which would have a lattice mismatch of approximately 13.8%, we instead grew an Al0.91Ga0.09Sb layer to help accommodate the lattice mismatch (~8.6% between GaAs and Al0.91Ga0.09Sb). A similar layer structure was used by Mishima et al.19 who showed a reduction in densities of both dislocations and microtwins with the incorporation of a 1 m thick AlSb layer which lead to better InSb QW properties, albeit in the (100) direction.
Using the
intermediate lattice constant of Al0.91Ga0.09Sb (6.132 Å) allows for a wide growth temperature window and the intermediate lattice constant leads to a lower density of defects due to reduced lattice mismatch.20 The ability to grow smooth high-quality Al0.91Ga0.09Sb barrier layers, would also provide an ideal barrier for the growth of InAs QWs. The addition of the interface between Al0.91Ga0.09Sb and In0.85Al0.15Sb will also facilitate the annihilation of threading dislocations, as the interface will act as a sink for the termination of dislocations (as will be shown subsequently). A composition of Al0.91Ga0.09Sb was used in order to maintain a large bandgap and to minimize oxidation of pure AlSb in post-growth processing.21 Growth of the Al0.91Ga0.09Sb layer was optimized with the results presented in Fig.2. A growth rate of 0.50 ML/s (corresponding to a Sb/(Al+Ga) BEP ratio of 43) was used, and the 5
temperature was varied between 440oC and 360oC to grow a 500 nm thick layer. The surface roughness with these conditions has a minimum at 400oC, and was lowered by reducing the growth rate to 0.25 ML/s. As observed in Fig.2, the growth done at 360oC is observed to have clusters most likely formed from the unincorporated adatoms on the surface in addition to dislocations. The films grown at 440oC and 400oC have a similar morphology which arises from the large lattice mismatch between the homoepitaxial layer and the Al0.91Ga0.09Sb layer. This leads to an “orange peel” morphology due to the large density of threading dislocations. This morphology arises from dislocation loop nucleation at the interface forming 60o dislocations. These threading dislocations relieve strain in the [001] direction and are analogous to those observed on InAs/GaAs (110) quantum well heterostructures as described by Jaszek et al.22 This may account for the roughening observed at the surface. Through cross-sectional TEM (XTEM) analysis, we will show that the Al0.91Ga0.09Sb layer is important for removing a large density of threading dislocations. For subsequent growths, the layer thickness was maintained at 250 nm which yielded a roughness of 0.63nm B. Optimization InAlSb (110) buffer layer growth The final layer in the step graded structure is an In0.85Al0.15Sb alloy grown to a thickness of 500 nm. In0.85Al0.15Sb was chosen since it has a lattice constant near the In0.8Al0.2Sb barrier and InSb quantum well.
The In0.8Al0.2Sb/InSb heterojunction has a conduction band offset
(~0.2eV) that is large enough to confine electrons in the InSb, but with a lattice mismatch that is small enough (~1.1%) for coherent growth of the InSb QW. The growth temperature was varied between 440ºC and 320ºC with the growth rate kept at 0.25 ML/s. The Sb/(Al+In) BEP ratio was 17.5. This layer has a narrower bandgap than the
6
layers already deposited so it absorbs more infrared light from the heater causing the substrate temperature to increase as the layer gets thicker. To maintain a constant growth temperature it is necessary to continually reduce the power supplied to the heater. We maintained constant temperature by continually monitoring the GaAs band edge spectra to adjust the heater power. The surface roughness as a function of growth temperature is shown in Fig.3.
Larger
micrographs of the surface morphology and surface profiles are given in Fig.4. At temperatures of 440ºC (Fig.4A), very distinct morphology was observed consisting of large randomly oriented mesas on the surface. The peak to valley height is approximately 130 nm and the distance between mesas is approximately 2 m laterally. Again, this can be attributed to the high temperature growth and the non-polar nature of the surface which allows for mass transport at the surface of the growth. The atoms are not held down by the polar surface therefore the added mobility leads them to more preferred orientations, of which the (001) is the most preferred growth plane. The instability at the surface is also influenced by the dislocations which preferentially grow along the [1 1 0] direction as will be discussed in more detail to follow. Upon lowering the growth temperature, the growth surface (as measured by AFM) flattened significantly until a minimum was reached at approximately 330ºC. The morphology of ridges along the [001] direction is reduced. As observed in the layers grown at 410oC (Fig.4B) and 360oC (Fig.4C), the angled steps are still prevalent. However, the steps become smaller as we go to lower temperatures. At temperatures below 340oC, substrate morphology is dominated by Sb segregation/clusters which account for the sharp noise associated with the linear height scan. At a substrate growth temperature of 320ºC, the surface (not shown) began to significantly
7
roughen. The RHEED images indicated that the growth proceeded in a 3D mode from the onset; it had a surface roughness of 105 nm. The wafer was too cold leading to stochastic roughening thereby putting a lower limit on the growth temperature.23 In all other previous growths, the growth started in a 3D RHEED pattern and within a few monolayers began to elongate into a 2D type growth. During growth the Sb/(Al+In) flux is kept at a BEP of 17.5 for the same reasons as the high As flux is needed to grow high quality arsenides. However, since the temperature is much lower than that at which the other layers are grown (~335ºC vs. 440ºC and 400ºC.) the sticking of the Sb to the surface increases and the BEP pressure of 17.5 is high enough to support the growth on the surface. Even at a low V/III flux of 17.5, too much Sb is being deposited onto the surface which leads to clusters of Sb. The Sb clusters have a distinct pyramidal morphology and a height of approximately 15 nm. A key concern for the growth of the In0.85Al0.15Sb arises from the large pits observed on the order of 50-100 nm deep. These are observed in Fig.4D,E,F The pitting was limited to samples grown at 360ºC and below (Figs.4D,E,F) whereas the samples grown above that temperature did not have pitting but rather a distinct morphology of hills and valleys.A distinct crystallographic growth direction is observed for the pitting with angled walls and arising from a central point. Due to their prevalence and large size, an understanding of the cause and how to eliminate the pitting is of utmost concern. In order to determine the cause of the pitting as well as to observe the morphology and dislocations of the structure, a cross sectional TEM image was taken of the sample grown at 335oC.
8
C.
TEM analysis of In0.85Al0.15Sb Buffer Layer
XTEM analysis was used to determine the cause of the pitting by examining the sample that was grown at 335ºC as shown in Fig.5. The bright field (BF) and dark field (DF) images are given of two separate areas in Figs.5A,C and Figs.5B,D respectively. The [1 1 0] direction is into the page. From Fig.5A (BF) and Fig.5B (DF), it is observed that the Al0.91Ga0.09Sb region has the bulk of the defects present that terminate at the interface with the In0.85Al0.15Sb layer. These are therefore looking down the core of the 60o dislocations which lay along the [1 1 0] directions. The threading dislocation density of the Al0.91Ga0.09Sb layer is approximately 3x1010/cm2. Many of the dislocations and defects are filtered at the Al0.91Ga0.09Sb/In0.85Al0.15Sb interface, with the result that the dislocation density is reduced to 3x109/cm2. In the In0.85Al0.15Sb layer, dislocations are seen to thread along the [110] direction as reported for other heteroepitaxial (110) growths.6,24 It is observed that the 60o dislocations thread from the AlGaSb/InAlSb interface to the film surface. The 60o type dislocations move along the (111) and (11 1 ) planes until they penetrate the surface of the film.24 From observing the bright field and dark field images in Figs.5C,D it is apparent that a pyramidal precipitate occurring at the Al0.91Ga0.09Sb/In0.85Al0.15Sb interface results in the pitting of the surface discussed previously. Transmission electron diffraction indicates that this is an ordered precipitate of the CuPt structure, which has been observed in MBE-grown InAs125,26,27
xSbx.
The rest of the film showed a dense defect structure with most of the dislocations
terminating at the Al0.91Ga0.09Sb /In0.85Al0.15Sb interface. pyramidal precipitate to excess Sb during deposition. 9
We attribute the formation of the While the deposition of GaAs and
Al0.91Ga0.09Sb required a very high flux of As and Sb, respectively, the low melting temperature of In0.85Al0.15Sb requires a much lower growth temperature (~335ºC). At this temperature, excess buildup of Sb on the surface (observed as clusters) leaves Sb crystals that nucleate on the surface, such as those observed in Fig.4D,E. These Sb crystals may assemble into larger crystals. This formation of Sb crystals at the surface during growth or even Sb buildup on the wafer holder ring causing pieces to flake off onto the surface of the wafer, may lead to the formation of pyramidal precipitates, as observed by XTEM in Figs.5D,C.
By observing the surface in the
XTEM image, it is apparent that the pyramidal precipitates lead to the pitting of the film surface. These types of pyramidal precipitates have been observed elsewhere in III-V growths. However, the cause of their growth in those systems does not correspond to the inclusions we observe.28 In a work by Xiong et al., (studying InSb growth on (100) GaAs) it was observed that hillock structures would form at the surface with excess Sb incorporation as well as be a driving cause for ripples at the surface (which may be the cause for the ripples observed in Fig.4).29 Therefore, excess Sb must be carefully controlled. Unlike in the (100) oriented system, where the polar surface easily holds the Sb forming full coverage, the (110) surface leads to difficulty in controlling full site incorporation. Though contrary to what was described as hillocks by Xiong et al., we believe that at such low growth temperatures, ~335oC, excess Sb buildup will occur. The excess buildup leads to segregation which forms into pyramidal shapes as the material is deposited. Once the material is used up, the film will continue to grow and slowly cover the inclusion. As can be seen from Fig.5D, the pyramid is slowly covered. It is understood that if we grew to a thicker film, the pyramid would be fully covered.
10
D. Effect of in situ annealing on surface characteristics Molecular beam epitaxy growth is a non-equilibrium process with thermodynamics providing the driving force for the morphological evolution of the surface but growth kinetics determining the surface morphology evolution.
The surface structures that are created due to
kinetic instabilities are metastable and can be annealed away after the conclusion of growth, provided the temperature is high enough for mass transport across the surface.
Though the
growth process by itself of the two layers yielded a surface roughness of ~5nm (Fig.4E), in situ annealing of each subsequent layer should bring the roughness down considerably. The result of annealing the Al0.91Ga0.09Sb layer (annealed at 530oC, not shown) was counter to that which was expected. Rather than smoothing the surface, the film roughness was essentially unchanged (0.90 nm to 0.91 nm RMS). This step, does, however, result in a way to mitigate dislocations at the next interface. Due to the large lattice mismatch between the GaAs and Al0.91Ga0.09Sb layers, a large density of dislocations are present. The two-layer structure was grown adding an in situ anneal after the first layer (Al0.91Ga0.09Sb) as well as the second layer in order to observe the effect of the anneal on the surface as shown in Fig.6A. Annealing after the Al0.91Ga0.09Sb layer only reduced film roughness (by giving a smoother growth layer); however, the Sb segregation still occurred at the surface of the In0.85Al0.15Sb layer. The large density of pitting did decrease due to removal of the Sb buildup from the wafer holder which was theorized to flake onto the surface of the wafer. The addition of the annealing step after the In0.85Al0.15Sb growth reduced the Sb surface segregation. The annealing took place at 425oC due to the lower melting temperature. The large degree of pitting may be due to the flaking of Sb from the
11
mounting ring during the growth process. However, this pitting was reduced by the annealing after the Al0.91Ga0.09Sb growth. The anneal after the growth of the Al0.91Ga0.09Sb layer was conducted at a sufficiently high temperature to evaporate the excess Sb from the wafer surface and therefore retard the formation of inclusions at the start of the In0.85Al0.15Sb deposition. However, this high temperature annealing treatment did not completely prevent the inclusion formation during the growth of In0.85Al0.15Sb from Sb (i.e. large pitting) that was still prevalent in the annealed films. A symmetric x-ray scan of the sample prepared with in situ annealed Al0.91Ga0.09Sb and in situ annealed In0.85Al0.15Sb is shown in Fig.6B. The sample shows only (220) diffraction peaks of the layered structure. Composition of all layers was verified by lattice parameter calculations. The rocking curve of the In0.85Al0.15Sb layer had a full width half maximum of 980 arc-seconds which most likely arises from the large pitting observed in the samples as well as the large degree of relaxation. The full width half maximum of the AlGaSb layer was slightly better with a value of 880 arc-seconds. In order to remove the large pitting caused by the Sb inclusions, a lower flux of Sb was used during the growth of the In0.85Al0.15Sb layer in order to prevent the buildup of antimony on the mounting ring. The lower melting temperature of In0.85Al0.15Sb, combined with the much higher incorporation coefficients observed at lower temperatures for (110) surfaces, allows a lower Sb flux rate to lead to smooth layered growth.8 The Sb valve position was reduced from 300 (fully open) to 100, which corresponds to a 5x reduction in the Sb BEP. Multiple annealing steps were also used during this growth of each layer. The surface roughness was observed to be 1.49 nm over an area of 5x5m as well as displaying the removal of all large pits due to the lower flux. The films appeared to be dense with no pinhole inclusions. It should be noted from 12
the XTEM images in Fig.5 that the [110] threading dislocations penetrate the surface with a density of approximately 1.5/m2 and an average length of approximately 500nm. These 60o threading dislocations terminate approximately 2nm above the surface of the film, which contributes to the roughness of the film. The final growth conditions incorporating all optimized parameters is given in Figure 7A. E. Characterization of (110) InSb Quantum Well InSb quantum wells were grown on an optimized buffer layer (each layer being in situ annealed). The heterostructure consisted of 60 nm In0.8Al0.2Sb on either side of a 13 nm or 30 nm InSb quantum well.
Si delta doping of 3x1012/cm2 are located symmetrically 20 nm from the
quantum well. The top of the active layer was capped with 20nm In0.8Al0.2Sb which is delta doped 20nm below the surface with Si (3x1012/cm2) in order to compensate the charge on the surface. The X-ray diffraction pattern for the sample and the AFM image of the 13 nm QW surface is given in Figs.7B,C respectively. The diffraction pattern for the 13nm InSb QW structure is given showing the well-defined peaks for the In0.8Al0.2Sb and InSb which arise from the active layer. The mobility was measured in a Van der Pauw geometry at 300K and 77K. Mobility for the 13 nm QW was 3,320 cm2/Vs at 300K (1.7x1012/cm2 electrons) and 3,520 at 77K (1.2x1012/cm2 electrons). For the 30nm QW, mobility was measured to be 3,290 cm2/Vs at room temperature (3.2x1012/cm2 electrons) and 3,620 at 77K (2.2x1012/cm2 electrons). All samples were measured with the lights on. Alhough these values are far lower than those reported for InSb/InAlSb quantum wells on (001) GaAs (electron mobility of 38,000 and 121,000 cm2/Vs at 300K and 77K, respectively with an electron density 4.0x1011/cm2)30, as an initial value for a new quantum well structure, there is much promise and room for improvement through the further reduction in roughness, elimination of inclusions and higher quality films. 13
In order to better quantify the confinement in the quantum well structure, photoluminescence (PL) spectra was taken. The samples were cooled down to 4.2K in a He cryostat with ZnSe windows. The samples were optically excited with 170mW over a 4mm2 area by an 808nm diode laser chopped at 18 kHz. The signal was detected with an InSb detector mounted on a Fourier Transform Infrared (FTIR) spectrometer. The 18 kHz modulated signal from the InSb detector was sent to a lock-in amplifier and the lock-in output was sent to the FTIR to transform the signal and obtain the spectral dependence of the PL. Figure 8A shows photoluminescence spectra for five quantum well samples, three grown on the (100) surface and two grown on the (110) surface. Quantum wells observed are of varying thickness; however, samples grown on (100) follow as consistent trend with the literature as a function of thickness for (100) InSb QWs.31 All samples had the same quantum well structure even in the case of different substrate surface orientation. (100) samples were grown using the same MBE, however with different optimized conditions due to the different surface orientation (i.e. lower group V overpressure and higher temperatures). The energy peaks vary from ~250 meV to ~500 meV. Two energy peaks are observed, one at ~450 meV that we identified as band to band recombination at the In0.85Al0.15Sb layers, and another one that varies from ~250 meV to ~370 meV as the quantum well width is decreased. The latter transition is identified as the ground state energy within the InSb quantum well, since the energy positions for all three samples grown on the (100) are consistent with previous measurements on similar samples as shown in Figure 8B. 31,32 The structure observed for the 8 nm quantum well at 350 meV is just an artifact as a result of an absorption line of the H-O stretch modes of water in the atmosphere.33 The line widths of all samples is ~50 meV, which is broader than studies on multiple quantum wells by Zhang et al.31 but are consistent with line widths reported in studies by Smith et al.34 14
and by Tenev et al.35 on (100) InSb quantum wells. Although the reason for these broad line widths merits further study, one possible explanation is that we are observing luminescence from several electron quantum well confined bands due to a large doping (of the order of 1012/cm2 electrons).35 The observation of luminescence from single InSb/InAlSb quantum wells grown on GaAs substrates confirms the relatively high quality of the materials grown. 4. Conclusion Due to the non-polar nature of the (110) surface, a wide variety of growth modes exist which make growing smooth buffer layers difficult. We have successfully grown a smooth buffer layer structure for (110) InSb growth and fabricated and characterized the active quantum well structures. We have significantly reduced the film roughness by varying the growth temperature, flux, and annealing of the films, leading to suitable conditions for quantum well structures. In general, a high group V to group III flux ratio is needed due to the difficulty in covering the surface with group V material. A lower deposition temperature is needed in order to allow for greater sticking of the group V materials to the surface.
These techniques should result in
superior InSb quantum wells for new and interesting studies in spin-based devices.
Acknowledgements: This work was funded by the Office of Naval Research. The authors thank Dr. Shawn Mack for comments and discussions.
15
Figure 1: Cross sectional schematic of the buffer layer structure. Figure 2: Surface roughness as a function of substrate deposition temperature for heteroepitaxial Al0.91Ga0.09Sb (110). All samples were 500nm except where indicated. Insets are AFM height sensor micrographs of 1 x 1 m scan areas. Figure 3: Surface roughness as a function of substrate deposition temperature for 500nm thick heteroepitaxial In0.85Al0.15Sb (110). Insets are AFM height sensor micrographs of 10x10m scan areas. Figure 4: AFM micrograph of the In0.85Al0.15Sb surface as a function of substrate deposition temperature from 440oC to 330oC (A-F). Figure 5: Two separate regions, Area #1 (A,B) and Area #2 (C,D), as observed by XTEM. (A,C) are bright field images recorded using the g=00 2 reflection, while (B,D) are dark field images recording using the g=002 reflection. Figure 6:
(A) Surface roughness comparison of In0.85Al0.15Sb buffer layer grown without
annealing, with annealing of only the Al0.91Ga0.09Sb layer and annealing of both the Al0.91Ga0.09Sb and In0.85Al0.15Sb layers. Note the removal of Sb clusters after the In0.85Al0.15Sb anneal. Insets are AFM height sensor micrographs of 5 x 5 m scan areas. (B) XRD of the sample with in situ annealed In0.85Al0.15Sb and Al0.91Ga0.09Sb layers. Only peaks from the (220) reflection are observed. Figure 7:
(A) Optimized growth parameters for each layer which were obtained from
experimentation. (B) X-ray diffraction pattern of the 13nm InSb QW. (C) Surface of the 13nm InSb QW. 16
Figure 8: (A) Photoluminescence spectra for (110) and (100) oriented quantum well at 4.2K. (B) Photoluminescence energy as a function of quantum well for (110) and (100) oriented quantum well at 4.2K as compared with results from reference [31]. The dashed green line is just a guide for the eye.
HighlightsforReview x x x x
HeteroepitaxialgrowthofhighqualityInAlSbbufferlayeron(110)GaAs Developmentoftechniquetodepositsmoothbufferlayerfor(110)InSbandInAs. Characterizationof(110)orientedInSbquantumwells Understandingofdislocationsandmorphologyfor(110)orientedfilmgrowth.
1
Y. Ohno, R. Terauchi, T. Adachi, F. Matsukura, H. Ohno, Phys. Rev. Letters 83 (1999) 4196. M. Franz Physics 3 (2010) 24. 3 L. Pfeiffer, K. W. West, H. L. Stormer, J. P. Eisenstein, K. W. Baldwin, D. Gershoni, J. Spector, Appl. Phys. Letters 56 (1990) 1697. 4 S. Koh, K. Ikeda, H. Kawaguchi, J. Appl. Phys. 110 (2011) 043516. 5 J. G. Belk, D.W. Pashley, C.F. McConville, B. A. Joyce, T. S. Jones, Surface Science 410 (1998) 82. 6 R. Jaszek, J. J. Harris, C. Roberts, J. Appl. Phys. 85 (1999) 959. 7 S. Iba, H. Fujino, T. Fujimoto, S. Koh, H. Kawaguchi, Physica E 41 (2009) 870. 8 E. S. Tok, T. S. Jones, J. H. Neave, J. Zhang, B. A. Joyce, Appl. Phys. Letters 71 (1997) 3278. 9 D. M. Holmes, E. S. Tok, J. L. Sudijono, T. S. Jones, B. A Joyce, J. Crystal Growth 192 (1998) 33. 10 X. Zhang, D. W. Pashley, L. Hart, J. H. Neave, P. N. Fawcett, B. A. Joyce, J. Crystal Growth 131 (1993) 300. 11 B. R. Bennett, R. Magno, J. B. Boos, W. Kruppa, M. G. Acona, Solid-State Electronics 49 (2005) 1875. 12 H. Kosaka, A. A. Kiselev, F. A. Baron, K. W. Kim, E. Yablonovitch, Electronic Letters 37 (2001) 464. 13 M. L. Crespillo, J. L. Sacedon, B. A. Joyce, P. Tejedor, Microelectronic Journal 36 (2005) 581. 14 P. Tejedor, M. L. Crespillo, B. A. Joyce, Mat. Sci. and Eng. C 26 (2006) 852. 15 L. S. Liu, W. X. Wang, Z. H. Li, B. L. Liu, H. M. Zhao, J. Wang, H. C. Gao, Z. W. Jiang, S. Liu, H. Chen, J. M. Zhou, J. Crystal Growth 301 (2007) 93. 2
17
R. Hey, A. Trampert, U. Jahn, O. D. D. Couto Jr., P. Santos J. Crystal Growth 301-302 158 (2007). 17 J.-H. Noh, S. Hasegawa, T. Suzuki, T. Arakawa, K. Tada and H. Asahi Physica E 23 (2004) 482. 18 Y. Horikoshi, M. Kawashima, H. Yamaguchi, Jpn. J. Appl. Phys. 25 (1988) L868. 19 T.D. Mishima, M. B. Santos, J. Vac. Sci. Technol. B 22 (2004) 1472. 20 L. K. Li, Y. Hsu, W. I. Wang, J Vac. Sci. Tech. B 11 (1993) 872. 21 B.R. Bennett, S.A. Khan, J.B. Boos, N.A. Papanicolaou, V.V. Kuznetsov, J. Electronic Materials 39 (2010) 2196. 22 R. Jaszek, J. J. Harris, C. Roberts, J. Appl. Phys. 85 (1999) 959. 23 W. M. Tong, R. S. Williams, Rev. Phys. Chem. 45 (1994) 401. 24 X. Zhang, D. W. Pashley, L. Hart, J. H. Neave, P. N. Fawcett, B. A. Joyce, J. Crystal Growth 131 (1993) 300. 25 A. G. Norman, T.-Y. Seong, I. T. Ferguson, G. R. Booker, B. A. Joyce, Semicond. Sci. Technol. 8 (1993) S9. 26 T.-Y. Seong, G. R. Booker, A. G. Norman, I. T. Ferguson, Appl. Phys. Letters 64 (1994) 3593. 27 D. M. Follstaedt R. M. Biefeld, S. R. Kurtz, K. C. Baucom, J. Electron. Mater. 24 (1995) 819. 28 D. M. Tricker, P. D. Brown, T. S. Cheng, C.T. Foxon, C. J. Humphrys, Appl. Surf. Sci. 123/124 (1998) 22. 29 M. Xiong, M. Li, L. Zhao, Phys. Stat. Sol. (RRL) 1 (2007) 80. 30 M. Edirisooriya, T. D. Mishima, C. K. Gaspe, K. Bottoms, R. J. Hauenstein, M.B. Santos, J. Crystal Growth 311 (2009) 1972. 31 X. H. Zhang, R. E. Doezema, N. Goel, S. J. Chung, M. B. Santos, N. Dai, F. H. Zhao, Z. S. Shi, Appl. Phys. Letters 89 (2006) 021907. 32 N. Dai, F. Brown, P. Barsic, G. A. Khodaparast, R. E. Doezema, M. B. Johnson, S. J. Chung, K. J. Goldammer, M. B. Santos, Appl. Phys. Letters 73 (1998) 1101. 33 J. E. Boercker, E. M. Clifton, J. G. Tischler, E. E. Foos. T. J. Zega, M. E. Twigg, R. M. Stroud, J. Phys. Chem. Letters 2 (2011) 527. 34 S. J. Smith, G. R. Nash, M. Fearn, L. Buckle, M. T. Emeny, T. Ashley, Appl. Phys.Letters 88 (2006) 051107. 35 T. G. Tenev, A. Palyi, B. I. Mirza, G. R. Nash, M. Fearn, S. J. Smith, L. Buckle, M. T. Emeny, T. Ashley J. H. Jefferson, C. J. Lambert, Phys. Rev. B 79 (2009) 085301. 16
18
Figure8
Figure7
Figure6
Figure5
Figure4
Figure3
Figure2
Figure1
Step graded buffer for (110) InSb quantum Wells grown by molecular beam epitaxy Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno, Brian R. Bennett
www.elsevier.com/locate/jcrysgro
PII: DOI: Reference:
S0022-0248(14)00501-6 http://dx.doi.org/10.1016/j.jcrysgro.2014.07.014 CRYS22337
To appear in:
Journal of Crystal Growth
Received date: 8 April 2014 Revised date: 1 July 2014 Accepted date: 6 July 2014 Cite this article as: Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno, Brian R. Bennett, Step graded buffer for (110) InSb quantum Wells grown by molecular beam epitaxy, Journal of Crystal Growth, http://dx.doi.org/ 10.1016/j.jcrysgro.2014.07.014 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Step Graded Buffer for (110) InSb Quantum Wells grown by Molecular Beam Epitaxy Adrian Podpirka, Mark E. Twigg, Joseph G. Tischler, Richard Magno and Brian R. Bennett Electronic Science and Technology Division, Naval Research Laboratory, Washington, DC Abstract We report on a two step buffer layer preparation for the growth of InSb quantum wells on a (110) GaAs surface. At each buffer layer step, layer conditions were optimized to produce smooth surfaces compatible with InSb quantum wells. Through varying growth rate, group V/III flux ratio, substrate temperature, and the addition of in situ annealing, we are able to grow In0.85Al0.15Sb on a GaAs substrate with an RMS surface roughness of approximately 2 nm. Surface morphology and cross-sectional transmission electron microscopy (TEM) were analyzed to understand the formation of threading dislocations, inclusions and dislocation filtering. This work presents an initial study for the growth of large lattice mismatched III-V materials on the (110) surface. Keywords: A3. molecular beam epitaxy, B2. Semiconducting III-V materials, B1. Antimonides, A1. Surfaces, A1. Crystal morphology 1. Introduction Understanding the details needed to grow semiconductors on a range of different substrate orientations will enable the fabrication of increasingly complex III-V heterostructures for advanced technologies. III-V growth by molecular beam epitaxy (MBE) has traditionally focused on (100) substrates in part due to the wide range of growth conditions which result in excellent epitaxial quality. Non (100) surface growth is also needed to take advantage of the anisotropic properties of the III-V semiconductors. Despite being the natural cleavage plane in 1
the zinc-blende structure, growth on the (110) surface has often been overlooked. Growth on the (110) surface could allow the development of new devices based on strong spin orbital coupling with applications to Majorana Fermions and spin-based devices due to the electron spin relaxation times of several nanoseconds at room temperature.1,2 The results of this work will also lead to a better understanding of cleaved edge overgrowth and to the overgrowth on patterned semiconductors containing edges with (110) surfaces that are useful for non-planar structures and devices.3 Further developments in the fabrication of device structures on the (110) surface requires a greater understanding of the growth processes and more specifically, the relationship between the growth conditions and the surface morphology of the films. Though the benefit of the (110) surface has been known for many years, little work has been done beyond the growth of InAs and GaAs/AlGaAs (110) quantum-well structures.4,5,6,7 The nature of the (110) surface, being non-polar as opposed to the standard (100) polar surface, leads to complications arising from the low incorporation rate of group V atoms onto the surface of the film as described by the work of Tok et al.
They have shown that the arsenic
incorporation kinetics vary greatly between the (100) and (110) surfaces for GaAs.8 This arises from the (110) surface having a single dangling adsorption site as opposed to the two dangling bonds found on the (100) surface. Therefore, the shorter residence time on the (110) surface for the group V means that its availability for incorporation is reduced. An increase in growth temperature will have a strong influence on the arsenic desorption rate and therefore the temperature of growth is of utmost importance. As a result, the general rule is that the growth on (110) requires higher group V fluxes and lower temperatures to allow for group V incorporation.9 Due to the orientation of the growth direction, two strain relief mechanisms are observed in the two orthogonal directions [1 1 0] and [001] which lead to strain relief. 60o 2
dislocations on the (110) mismatched growth have also been responsible for observable tilting of the surface which further complicates the growth as opposed to growths on the traditional (100) surface.10 A wide variety of growth regimes exist on the (110) surface particularly due to the difficulty of group V incorporation.
A systematic study on the effects of the deposition
conditions on the surface morphology is required to define the conditions under which highquality films can be produced. Antimonide based semiconductors are candidates for high-speed, low-power electronic devices,11 and have the highest electron g-factor amongst all III-V semiconductors.12 However, we are not aware of any systematic study on the growth of buffer layers for (110) antimonides on semi-insulating wafers, namely (110) GaAs. In this work, we investigated the growth conditions for reducing surface roughness using a two step buffer layer process to accommodate the 13.8% lattice mismatch between the GaAs substrate and an In0.85Al0.15Sb layer. We varied the flux rate and growth temperature at each buffer layer and determined surface roughness and morphology with atomic force microscopy (AFM). Each buffer layer step was investigated until the lowest roughness was achieved within our test parameters (as measured by AFM) before using it as a buffer for subsequent layer growths. The smoothest films were grown at low substrate temperatures, with the use of multiple in situ annealing steps and the careful control of group V flux rates. We fabricated and characterized an active QW structure after the optimized buffer layer was determined. 2. Experimental Procedures Samples were grown in a Riber 32P MBE system using cracked As2 and Sb2 sources on semi-insulating GaAs (110) on-axis substrates. Prior to growth, the native oxide layer on the
3
surface was removed by heating the substrate to 620oC for 10 minutes under an As flux. Temperature was measured with a transmission spectroscopy system to measure the band edge of GaAs, using the substrate heater as a light source. After removing the native oxide layer, the (110) surface displayed a (1 x 1) streaked reflection high-energy electron diffraction (RHEED) pattern.
Growth rates were obtained by RHEED oscillations on (100) GaAs wafers prior to
deposition. The two step graded buffer layer structure is shown by the schematic diagram in Fig.1. The structure of this paper is the optimization of the heteroepitaxial Al0.91Ga0.09Sb on GaAs (110) followed by the optimization of the In0.85Al0.15Sb (110) layer on top of the optimized Al0.91Ga0.09Sb layer. The effect of in-situ annealing on the surface morphology of the buffer layers is also investigated. We varied substrate temperatures (320oC to 440oC) and fluxes in order to optimize growth conditions. In situ annealing was performed under group V flux. The surface was characterized by atomic force microscopy (AFM) to determine the surface morphology as well as high resolution X-ray diffraction using a (022) Ge monochromater Cu K beam to determine the lattice constant and composition. After the optimized buffer layer structure was completed, a quantum well layer was deposited as a proof of concept and the electronic states were probed using photoluminescence. 3. Results and Discussion Previous to the metamorphic growth on (110) GaAs, a homoepitaxial smoothing layer was deposited at a thickness of 30nm. Standard MBE growth was used in the deposition of the smoothening layer with an optimized surface roughness of approximately 0.3nm being obtained. Homoepitaxial growth of GaAs has been studied in depth elsewhere. 8,13,14,15 Our optimized parameters consisted of a As/Ga beam-equivalent pressure (BEP) 138 (as determined by flux gauge measurements) and a deposition rate of 0.15ML/s at a temperature of 440oC. It should be 4
noted that there are other methods to obtain atomic-terraced homoepitaxial (110) GaAs such as migration enhanced epitaxy (MEE). In the MEE method, Group III atoms are supplied on the surface under Group-V-free atmosphere or a very low Group V pressure allowing the Group III atoms to migrate for long distances.16,17,18 A.
Optimization of AlGaSb (110) buffer layer growth
In order to accommodate the large lattice mismatch between the GaAs substrate and the InSb Quantum Well (~14.7%), an intermediate layer was used on top of the prepared GaAs surface discussed in the previous section. Rather than going directly to an In0.85Al0.15Sb buffer layer which would have a lattice mismatch of approximately 13.8%, we instead grew an Al0.91Ga0.09Sb layer to help accommodate the lattice mismatch (~8.6% between GaAs and Al0.91Ga0.09Sb). A similar layer structure was used by Mishima et al.19 who showed a reduction in densities of both dislocations and microtwins with the incorporation of a 1 m thick AlSb layer which lead to better InSb QW properties, albeit in the (100) direction.
Using the
intermediate lattice constant of Al0.91Ga0.09Sb (6.132 Å) allows for a wide growth temperature window and the intermediate lattice constant leads to a lower density of defects due to reduced lattice mismatch.20 The ability to grow smooth high-quality Al0.91Ga0.09Sb barrier layers, would also provide an ideal barrier for the growth of InAs QWs. The addition of the interface between Al0.91Ga0.09Sb and In0.85Al0.15Sb will also facilitate the annihilation of threading dislocations, as the interface will act as a sink for the termination of dislocations (as will be shown subsequently). A composition of Al0.91Ga0.09Sb was used in order to maintain a large bandgap and to minimize oxidation of pure AlSb in post-growth processing.21 Growth of the Al0.91Ga0.09Sb layer was optimized with the results presented in Fig.2. A growth rate of 0.50 ML/s (corresponding to a Sb/(Al+Ga) BEP ratio of 43) was used, and the 5
temperature was varied between 440oC and 360oC to grow a 500 nm thick layer. The surface roughness with these conditions has a minimum at 400oC, and was lowered by reducing the growth rate to 0.25 ML/s. As observed in Fig.2, the growth done at 360oC is observed to have clusters most likely formed from the unincorporated adatoms on the surface in addition to dislocations. The films grown at 440oC and 400oC have a similar morphology which arises from the large lattice mismatch between the homoepitaxial layer and the Al0.91Ga0.09Sb layer. This leads to an “orange peel” morphology due to the large density of threading dislocations. This morphology arises from dislocation loop nucleation at the interface forming 60o dislocations. These threading dislocations relieve strain in the [001] direction and are analogous to those observed on InAs/GaAs (110) quantum well heterostructures as described by Jaszek et al.22 This may account for the roughening observed at the surface. Through cross-sectional TEM (XTEM) analysis, we will show that the Al0.91Ga0.09Sb layer is important for removing a large density of threading dislocations. For subsequent growths, the layer thickness was maintained at 250 nm which yielded a roughness of 0.63nm B. Optimization InAlSb (110) buffer layer growth The final layer in the step graded structure is an In0.85Al0.15Sb alloy grown to a thickness of 500 nm. In0.85Al0.15Sb was chosen since it has a lattice constant near the In0.8Al0.2Sb barrier and InSb quantum well.
The In0.8Al0.2Sb/InSb heterojunction has a conduction band offset
(~0.2eV) that is large enough to confine electrons in the InSb, but with a lattice mismatch that is small enough (~1.1%) for coherent growth of the InSb QW. The growth temperature was varied between 440ºC and 320ºC with the growth rate kept at 0.25 ML/s. The Sb/(Al+In) BEP ratio was 17.5. This layer has a narrower bandgap than the
6
layers already deposited so it absorbs more infrared light from the heater causing the substrate temperature to increase as the layer gets thicker. To maintain a constant growth temperature it is necessary to continually reduce the power supplied to the heater. We maintained constant temperature by continually monitoring the GaAs band edge spectra to adjust the heater power. The surface roughness as a function of growth temperature is shown in Fig.3.
Larger
micrographs of the surface morphology and surface profiles are given in Fig.4. At temperatures of 440ºC (Fig.4A), very distinct morphology was observed consisting of large randomly oriented mesas on the surface. The peak to valley height is approximately 130 nm and the distance between mesas is approximately 2 m laterally. Again, this can be attributed to the high temperature growth and the non-polar nature of the surface which allows for mass transport at the surface of the growth. The atoms are not held down by the polar surface therefore the added mobility leads them to more preferred orientations, of which the (001) is the most preferred growth plane. The instability at the surface is also influenced by the dislocations which preferentially grow along the [1 1 0] direction as will be discussed in more detail to follow. Upon lowering the growth temperature, the growth surface (as measured by AFM) flattened significantly until a minimum was reached at approximately 330ºC. The morphology of ridges along the [001] direction is reduced. As observed in the layers grown at 410oC (Fig.4B) and 360oC (Fig.4C), the angled steps are still prevalent. However, the steps become smaller as we go to lower temperatures. At temperatures below 340oC, substrate morphology is dominated by Sb segregation/clusters which account for the sharp noise associated with the linear height scan. At a substrate growth temperature of 320ºC, the surface (not shown) began to significantly
7
roughen. The RHEED images indicated that the growth proceeded in a 3D mode from the onset; it had a surface roughness of 105 nm. The wafer was too cold leading to stochastic roughening thereby putting a lower limit on the growth temperature.23 In all other previous growths, the growth started in a 3D RHEED pattern and within a few monolayers began to elongate into a 2D type growth. During growth the Sb/(Al+In) flux is kept at a BEP of 17.5 for the same reasons as the high As flux is needed to grow high quality arsenides. However, since the temperature is much lower than that at which the other layers are grown (~335ºC vs. 440ºC and 400ºC.) the sticking of the Sb to the surface increases and the BEP pressure of 17.5 is high enough to support the growth on the surface. Even at a low V/III flux of 17.5, too much Sb is being deposited onto the surface which leads to clusters of Sb. The Sb clusters have a distinct pyramidal morphology and a height of approximately 15 nm. A key concern for the growth of the In0.85Al0.15Sb arises from the large pits observed on the order of 50-100 nm deep. These are observed in Fig.4D,E,F The pitting was limited to samples grown at 360ºC and below (Figs.4D,E,F) whereas the samples grown above that temperature did not have pitting but rather a distinct morphology of hills and valleys.A distinct crystallographic growth direction is observed for the pitting with angled walls and arising from a central point. Due to their prevalence and large size, an understanding of the cause and how to eliminate the pitting is of utmost concern. In order to determine the cause of the pitting as well as to observe the morphology and dislocations of the structure, a cross sectional TEM image was taken of the sample grown at 335oC.
8
C.
TEM analysis of In0.85Al0.15Sb Buffer Layer
XTEM analysis was used to determine the cause of the pitting by examining the sample that was grown at 335ºC as shown in Fig.5. The bright field (BF) and dark field (DF) images are given of two separate areas in Figs.5A,C and Figs.5B,D respectively. The [1 1 0] direction is into the page. From Fig.5A (BF) and Fig.5B (DF), it is observed that the Al0.91Ga0.09Sb region has the bulk of the defects present that terminate at the interface with the In0.85Al0.15Sb layer. These are therefore looking down the core of the 60o dislocations which lay along the [1 1 0] directions. The threading dislocation density of the Al0.91Ga0.09Sb layer is approximately 3x1010/cm2. Many of the dislocations and defects are filtered at the Al0.91Ga0.09Sb/In0.85Al0.15Sb interface, with the result that the dislocation density is reduced to 3x109/cm2. In the In0.85Al0.15Sb layer, dislocations are seen to thread along the [110] direction as reported for other heteroepitaxial (110) growths.6,24 It is observed that the 60o dislocations thread from the AlGaSb/InAlSb interface to the film surface. The 60o type dislocations move along the (111) and (11 1 ) planes until they penetrate the surface of the film.24 From observing the bright field and dark field images in Figs.5C,D it is apparent that a pyramidal precipitate occurring at the Al0.91Ga0.09Sb/In0.85Al0.15Sb interface results in the pitting of the surface discussed previously. Transmission electron diffraction indicates that this is an ordered precipitate of the CuPt structure, which has been observed in MBE-grown InAs125,26,27
xSbx.
The rest of the film showed a dense defect structure with most of the dislocations
terminating at the Al0.91Ga0.09Sb /In0.85Al0.15Sb interface. pyramidal precipitate to excess Sb during deposition. 9
We attribute the formation of the While the deposition of GaAs and
Al0.91Ga0.09Sb required a very high flux of As and Sb, respectively, the low melting temperature of In0.85Al0.15Sb requires a much lower growth temperature (~335ºC). At this temperature, excess buildup of Sb on the surface (observed as clusters) leaves Sb crystals that nucleate on the surface, such as those observed in Fig.4D,E. These Sb crystals may assemble into larger crystals. This formation of Sb crystals at the surface during growth or even Sb buildup on the wafer holder ring causing pieces to flake off onto the surface of the wafer, may lead to the formation of pyramidal precipitates, as observed by XTEM in Figs.5D,C.
By observing the surface in the
XTEM image, it is apparent that the pyramidal precipitates lead to the pitting of the film surface. These types of pyramidal precipitates have been observed elsewhere in III-V growths. However, the cause of their growth in those systems does not correspond to the inclusions we observe.28 In a work by Xiong et al., (studying InSb growth on (100) GaAs) it was observed that hillock structures would form at the surface with excess Sb incorporation as well as be a driving cause for ripples at the surface (which may be the cause for the ripples observed in Fig.4).29 Therefore, excess Sb must be carefully controlled. Unlike in the (100) oriented system, where the polar surface easily holds the Sb forming full coverage, the (110) surface leads to difficulty in controlling full site incorporation. Though contrary to what was described as hillocks by Xiong et al., we believe that at such low growth temperatures, ~335oC, excess Sb buildup will occur. The excess buildup leads to segregation which forms into pyramidal shapes as the material is deposited. Once the material is used up, the film will continue to grow and slowly cover the inclusion. As can be seen from Fig.5D, the pyramid is slowly covered. It is understood that if we grew to a thicker film, the pyramid would be fully covered.
10
D. Effect of in situ annealing on surface characteristics Molecular beam epitaxy growth is a non-equilibrium process with thermodynamics providing the driving force for the morphological evolution of the surface but growth kinetics determining the surface morphology evolution.
The surface structures that are created due to
kinetic instabilities are metastable and can be annealed away after the conclusion of growth, provided the temperature is high enough for mass transport across the surface.
Though the
growth process by itself of the two layers yielded a surface roughness of ~5nm (Fig.4E), in situ annealing of each subsequent layer should bring the roughness down considerably. The result of annealing the Al0.91Ga0.09Sb layer (annealed at 530oC, not shown) was counter to that which was expected. Rather than smoothing the surface, the film roughness was essentially unchanged (0.90 nm to 0.91 nm RMS). This step, does, however, result in a way to mitigate dislocations at the next interface. Due to the large lattice mismatch between the GaAs and Al0.91Ga0.09Sb layers, a large density of dislocations are present. The two-layer structure was grown adding an in situ anneal after the first layer (Al0.91Ga0.09Sb) as well as the second layer in order to observe the effect of the anneal on the surface as shown in Fig.6A. Annealing after the Al0.91Ga0.09Sb layer only reduced film roughness (by giving a smoother growth layer); however, the Sb segregation still occurred at the surface of the In0.85Al0.15Sb layer. The large density of pitting did decrease due to removal of the Sb buildup from the wafer holder which was theorized to flake onto the surface of the wafer. The addition of the annealing step after the In0.85Al0.15Sb growth reduced the Sb surface segregation. The annealing took place at 425oC due to the lower melting temperature. The large degree of pitting may be due to the flaking of Sb from the
11
mounting ring during the growth process. However, this pitting was reduced by the annealing after the Al0.91Ga0.09Sb growth. The anneal after the growth of the Al0.91Ga0.09Sb layer was conducted at a sufficiently high temperature to evaporate the excess Sb from the wafer surface and therefore retard the formation of inclusions at the start of the In0.85Al0.15Sb deposition. However, this high temperature annealing treatment did not completely prevent the inclusion formation during the growth of In0.85Al0.15Sb from Sb (i.e. large pitting) that was still prevalent in the annealed films. A symmetric x-ray scan of the sample prepared with in situ annealed Al0.91Ga0.09Sb and in situ annealed In0.85Al0.15Sb is shown in Fig.6B. The sample shows only (220) diffraction peaks of the layered structure. Composition of all layers was verified by lattice parameter calculations. The rocking curve of the In0.85Al0.15Sb layer had a full width half maximum of 980 arc-seconds which most likely arises from the large pitting observed in the samples as well as the large degree of relaxation. The full width half maximum of the AlGaSb layer was slightly better with a value of 880 arc-seconds. In order to remove the large pitting caused by the Sb inclusions, a lower flux of Sb was used during the growth of the In0.85Al0.15Sb layer in order to prevent the buildup of antimony on the mounting ring. The lower melting temperature of In0.85Al0.15Sb, combined with the much higher incorporation coefficients observed at lower temperatures for (110) surfaces, allows a lower Sb flux rate to lead to smooth layered growth.8 The Sb valve position was reduced from 300 (fully open) to 100, which corresponds to a 5x reduction in the Sb BEP. Multiple annealing steps were also used during this growth of each layer. The surface roughness was observed to be 1.49 nm over an area of 5x5m as well as displaying the removal of all large pits due to the lower flux. The films appeared to be dense with no pinhole inclusions. It should be noted from 12
the XTEM images in Fig.5 that the [110] threading dislocations penetrate the surface with a density of approximately 1.5/m2 and an average length of approximately 500nm. These 60o threading dislocations terminate approximately 2nm above the surface of the film, which contributes to the roughness of the film. The final growth conditions incorporating all optimized parameters is given in Figure 7A. E. Characterization of (110) InSb Quantum Well InSb quantum wells were grown on an optimized buffer layer (each layer being in situ annealed). The heterostructure consisted of 60 nm In0.8Al0.2Sb on either side of a 13 nm or 30 nm InSb quantum well.
Si delta doping of 3x1012/cm2 are located symmetrically 20 nm from the
quantum well. The top of the active layer was capped with 20nm In0.8Al0.2Sb which is delta doped 20nm below the surface with Si (3x1012/cm2) in order to compensate the charge on the surface. The X-ray diffraction pattern for the sample and the AFM image of the 13 nm QW surface is given in Figs.7B,C respectively. The diffraction pattern for the 13nm InSb QW structure is given showing the well-defined peaks for the In0.8Al0.2Sb and InSb which arise from the active layer. The mobility was measured in a Van der Pauw geometry at 300K and 77K. Mobility for the 13 nm QW was 3,320 cm2/Vs at 300K (1.7x1012/cm2 electrons) and 3,520 at 77K (1.2x1012/cm2 electrons). For the 30nm QW, mobility was measured to be 3,290 cm2/Vs at room temperature (3.2x1012/cm2 electrons) and 3,620 at 77K (2.2x1012/cm2 electrons). All samples were measured with the lights on. Alhough these values are far lower than those reported for InSb/InAlSb quantum wells on (001) GaAs (electron mobility of 38,000 and 121,000 cm2/Vs at 300K and 77K, respectively with an electron density 4.0x1011/cm2)30, as an initial value for a new quantum well structure, there is much promise and room for improvement through the further reduction in roughness, elimination of inclusions and higher quality films. 13
In order to better quantify the confinement in the quantum well structure, photoluminescence (PL) spectra was taken. The samples were cooled down to 4.2K in a He cryostat with ZnSe windows. The samples were optically excited with 170mW over a 4mm2 area by an 808nm diode laser chopped at 18 kHz. The signal was detected with an InSb detector mounted on a Fourier Transform Infrared (FTIR) spectrometer. The 18 kHz modulated signal from the InSb detector was sent to a lock-in amplifier and the lock-in output was sent to the FTIR to transform the signal and obtain the spectral dependence of the PL. Figure 8A shows photoluminescence spectra for five quantum well samples, three grown on the (100) surface and two grown on the (110) surface. Quantum wells observed are of varying thickness; however, samples grown on (100) follow as consistent trend with the literature as a function of thickness for (100) InSb QWs.31 All samples had the same quantum well structure even in the case of different substrate surface orientation. (100) samples were grown using the same MBE, however with different optimized conditions due to the different surface orientation (i.e. lower group V overpressure and higher temperatures). The energy peaks vary from ~250 meV to ~500 meV. Two energy peaks are observed, one at ~450 meV that we identified as band to band recombination at the In0.85Al0.15Sb layers, and another one that varies from ~250 meV to ~370 meV as the quantum well width is decreased. The latter transition is identified as the ground state energy within the InSb quantum well, since the energy positions for all three samples grown on the (100) are consistent with previous measurements on similar samples as shown in Figure 8B. 31,32 The structure observed for the 8 nm quantum well at 350 meV is just an artifact as a result of an absorption line of the H-O stretch modes of water in the atmosphere.33 The line widths of all samples is ~50 meV, which is broader than studies on multiple quantum wells by Zhang et al.31 but are consistent with line widths reported in studies by Smith et al.34 14
and by Tenev et al.35 on (100) InSb quantum wells. Although the reason for these broad line widths merits further study, one possible explanation is that we are observing luminescence from several electron quantum well confined bands due to a large doping (of the order of 1012/cm2 electrons).35 The observation of luminescence from single InSb/InAlSb quantum wells grown on GaAs substrates confirms the relatively high quality of the materials grown. 4. Conclusion Due to the non-polar nature of the (110) surface, a wide variety of growth modes exist which make growing smooth buffer layers difficult. We have successfully grown a smooth buffer layer structure for (110) InSb growth and fabricated and characterized the active quantum well structures. We have significantly reduced the film roughness by varying the growth temperature, flux, and annealing of the films, leading to suitable conditions for quantum well structures. In general, a high group V to group III flux ratio is needed due to the difficulty in covering the surface with group V material. A lower deposition temperature is needed in order to allow for greater sticking of the group V materials to the surface.
These techniques should result in
superior InSb quantum wells for new and interesting studies in spin-based devices.
Acknowledgements: This work was funded by the Office of Naval Research. The authors thank Dr. Shawn Mack for comments and discussions.
15
Figure 1: Cross sectional schematic of the buffer layer structure. Figure 2: Surface roughness as a function of substrate deposition temperature for heteroepitaxial Al0.91Ga0.09Sb (110). All samples were 500nm except where indicated. Insets are AFM height sensor micrographs of 1 x 1 m scan areas. Figure 3: Surface roughness as a function of substrate deposition temperature for 500nm thick heteroepitaxial In0.85Al0.15Sb (110). Insets are AFM height sensor micrographs of 10x10m scan areas. Figure 4: AFM micrograph of the In0.85Al0.15Sb surface as a function of substrate deposition temperature from 440oC to 330oC (A-F). Figure 5: Two separate regions, Area #1 (A,B) and Area #2 (C,D), as observed by XTEM. (A,C) are bright field images recorded using the g=00 2 reflection, while (B,D) are dark field images recording using the g=002 reflection. Figure 6:
(A) Surface roughness comparison of In0.85Al0.15Sb buffer layer grown without
annealing, with annealing of only the Al0.91Ga0.09Sb layer and annealing of both the Al0.91Ga0.09Sb and In0.85Al0.15Sb layers. Note the removal of Sb clusters after the In0.85Al0.15Sb anneal. Insets are AFM height sensor micrographs of 5 x 5 m scan areas. (B) XRD of the sample with in situ annealed In0.85Al0.15Sb and Al0.91Ga0.09Sb layers. Only peaks from the (220) reflection are observed. Figure 7:
(A) Optimized growth parameters for each layer which were obtained from
experimentation. (B) X-ray diffraction pattern of the 13nm InSb QW. (C) Surface of the 13nm InSb QW. 16
Figure 8: (A) Photoluminescence spectra for (110) and (100) oriented quantum well at 4.2K. (B) Photoluminescence energy as a function of quantum well for (110) and (100) oriented quantum well at 4.2K as compared with results from reference [31]. The dashed green line is just a guide for the eye.
HighlightsforReview x x x x
HeteroepitaxialgrowthofhighqualityInAlSbbufferlayeron(110)GaAs Developmentoftechniquetodepositsmoothbufferlayerfor(110)InSbandInAs. Characterizationof(110)orientedInSbquantumwells Understandingofdislocationsandmorphologyfor(110)orientedfilmgrowth.
1
Y. Ohno, R. Terauchi, T. Adachi, F. Matsukura, H. Ohno, Phys. Rev. Letters 83 (1999) 4196. M. Franz Physics 3 (2010) 24. 3 L. Pfeiffer, K. W. West, H. L. Stormer, J. P. Eisenstein, K. W. Baldwin, D. Gershoni, J. Spector, Appl. Phys. Letters 56 (1990) 1697. 4 S. Koh, K. Ikeda, H. Kawaguchi, J. Appl. Phys. 110 (2011) 043516. 5 J. G. Belk, D.W. Pashley, C.F. McConville, B. A. Joyce, T. S. Jones, Surface Science 410 (1998) 82. 6 R. Jaszek, J. J. Harris, C. Roberts, J. Appl. Phys. 85 (1999) 959. 7 S. Iba, H. Fujino, T. Fujimoto, S. Koh, H. Kawaguchi, Physica E 41 (2009) 870. 8 E. S. Tok, T. S. Jones, J. H. Neave, J. Zhang, B. A. Joyce, Appl. Phys. Letters 71 (1997) 3278. 9 D. M. Holmes, E. S. Tok, J. L. Sudijono, T. S. Jones, B. A Joyce, J. Crystal Growth 192 (1998) 33. 10 X. Zhang, D. W. Pashley, L. Hart, J. H. Neave, P. N. Fawcett, B. A. Joyce, J. Crystal Growth 131 (1993) 300. 11 B. R. Bennett, R. Magno, J. B. Boos, W. Kruppa, M. G. Acona, Solid-State Electronics 49 (2005) 1875. 12 H. Kosaka, A. A. Kiselev, F. A. Baron, K. W. Kim, E. Yablonovitch, Electronic Letters 37 (2001) 464. 13 M. L. Crespillo, J. L. Sacedon, B. A. Joyce, P. Tejedor, Microelectronic Journal 36 (2005) 581. 14 P. Tejedor, M. L. Crespillo, B. A. Joyce, Mat. Sci. and Eng. C 26 (2006) 852. 15 L. S. Liu, W. X. Wang, Z. H. Li, B. L. Liu, H. M. Zhao, J. Wang, H. C. Gao, Z. W. Jiang, S. Liu, H. Chen, J. M. Zhou, J. Crystal Growth 301 (2007) 93. 2
17
R. Hey, A. Trampert, U. Jahn, O. D. D. Couto Jr., P. Santos J. Crystal Growth 301-302 158 (2007). 17 J.-H. Noh, S. Hasegawa, T. Suzuki, T. Arakawa, K. Tada and H. Asahi Physica E 23 (2004) 482. 18 Y. Horikoshi, M. Kawashima, H. Yamaguchi, Jpn. J. Appl. Phys. 25 (1988) L868. 19 T.D. Mishima, M. B. Santos, J. Vac. Sci. Technol. B 22 (2004) 1472. 20 L. K. Li, Y. Hsu, W. I. Wang, J Vac. Sci. Tech. B 11 (1993) 872. 21 B.R. Bennett, S.A. Khan, J.B. Boos, N.A. Papanicolaou, V.V. Kuznetsov, J. Electronic Materials 39 (2010) 2196. 22 R. Jaszek, J. J. Harris, C. Roberts, J. Appl. Phys. 85 (1999) 959. 23 W. M. Tong, R. S. Williams, Rev. Phys. Chem. 45 (1994) 401. 24 X. Zhang, D. W. Pashley, L. Hart, J. H. Neave, P. N. Fawcett, B. A. Joyce, J. Crystal Growth 131 (1993) 300. 25 A. G. Norman, T.-Y. Seong, I. T. Ferguson, G. R. Booker, B. A. Joyce, Semicond. Sci. Technol. 8 (1993) S9. 26 T.-Y. Seong, G. R. Booker, A. G. Norman, I. T. Ferguson, Appl. Phys. Letters 64 (1994) 3593. 27 D. M. Follstaedt R. M. Biefeld, S. R. Kurtz, K. C. Baucom, J. Electron. Mater. 24 (1995) 819. 28 D. M. Tricker, P. D. Brown, T. S. Cheng, C.T. Foxon, C. J. Humphrys, Appl. Surf. Sci. 123/124 (1998) 22. 29 M. Xiong, M. Li, L. Zhao, Phys. Stat. Sol. (RRL) 1 (2007) 80. 30 M. Edirisooriya, T. D. Mishima, C. K. Gaspe, K. Bottoms, R. J. Hauenstein, M.B. Santos, J. Crystal Growth 311 (2009) 1972. 31 X. H. Zhang, R. E. Doezema, N. Goel, S. J. Chung, M. B. Santos, N. Dai, F. H. Zhao, Z. S. Shi, Appl. Phys. Letters 89 (2006) 021907. 32 N. Dai, F. Brown, P. Barsic, G. A. Khodaparast, R. E. Doezema, M. B. Johnson, S. J. Chung, K. J. Goldammer, M. B. Santos, Appl. Phys. Letters 73 (1998) 1101. 33 J. E. Boercker, E. M. Clifton, J. G. Tischler, E. E. Foos. T. J. Zega, M. E. Twigg, R. M. Stroud, J. Phys. Chem. Letters 2 (2011) 527. 34 S. J. Smith, G. R. Nash, M. Fearn, L. Buckle, M. T. Emeny, T. Ashley, Appl. Phys.Letters 88 (2006) 051107. 35 T. G. Tenev, A. Palyi, B. I. Mirza, G. R. Nash, M. Fearn, S. J. Smith, L. Buckle, M. T. Emeny, T. Ashley J. H. Jefferson, C. J. Lambert, Phys. Rev. B 79 (2009) 085301. 16
18
Figure8
Figure7
Figure6
Figure5
Figure4
Figure3
Figure2
Figure1