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TABLE OF CONTENTS (TOC)
Nano Research DOI 10.1007/s12274-015-0940-6
www.theNanoResearch.com∣www.Springer.com/journal/12274 | Nano Research
64 Nano Res.
Understanding and optimizing spin injection in
self-assembled InAs/GaAs quantum-dot molecular
structures
Y. Q. Huang, Y. Puttisong, I.A. Buyanova, W.M. Chen*
Department of Physics, Chemistry and Biology,
Linköping University, S-581 83 Linköping, Sweden
Spin injection into self-organized InAs/GaAs quantum-dot molecular
structures (QMSs) is shown to originate from the regions of the barrier
layers immediately surrounding the QMSs with local potentials, which
significantly deviate from the commonly believed 2D or 3D nature and
lead to severe spin loss. QMSs with a high symmetry of lateral
arrangement are identified as the optimal structure, not only for spin
detection, but also for spin injection.
Understanding and optimizing spin injection in self-assembled InAs/GaAs quantum-dot molecular structures
Yuqing Huang1, Yuttapoom Puttisong1, Irina A. Buyanova1, Weimin M. Chen1 (*)
1 Department of Physics, Chemistry and Biology, Linköping University, S-581 83 Linköping, Sweden
Semiconductor quantum-dot (QD) structures are promising for spintronic applications owing to strong quenching of spin relaxation processes promoted by carrier and excitons motions. Unfortunately, spin injection efficiency in such nanostructures remains very low and the exact physical mechanism for the spin loss is still not fully understood. Here, we show that exciton spin injection in self-assembled InAs/GaAs QDs and quantum-dot molecular structures (QMSs) is dominated by localized excitons confined within the QD-like regions of the wetting layer (WL) and GaAs barrier layer immediately surrounding QDs and QMSs that in fact lack the commonly believed 2D and 3D character with an extended wavefunction. We identify the microscopic origin of the observedsevere spin loss during spin injection as being due to a sizable anisotropic exchange interaction (AEI) of the localized excitons in the WL and GaAs barrier layer, which has so far been overlooked. We find that the AEI of the injected excitons and thus the efficiency of the spin injection processes are correlated with the overall geometric symmetry of the QMSs, as the latter largely defines the anisotropy of the confinement potential of the localized excitons in the surrounding WL and GaAs barrier. These results pave the way for a better understanding of spin injection processes and the microscopic origin of spin loss in QD structures, which in turn provides a useful guideline to significantly improve spin injection efficiency by optimizing the lateral arrangement of the QMSs thereby overcoming a major bottleneck in spintronic device applications utilizing semiconductor QDs.
Nano Research DOI (automatically inserted by the publisher) Research Article
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2 Nano Res.
While conventional electronics only relies on manipulation of electronic charge, a promising direction for next-generation electronic and photonic devices is expected to benefit additionally from exploiting the spin degree of freedom [1-7]. Among vital ingredients in achieving spin functionalities are efficient injection of spin polarized carriers or excitons, prolonged spin dephasing/relaxation time as well as our ability to manipulate spin in a desirable way. Semiconductor quantum dots (QDs) have emerged as an excellent materials system in this context, as motion-induced spin relaxation that is dominant in bulk and 2D semiconductor structures is quenched due to the three-dimensional confinement [2,6,8-12]. Indeed, appreciably long spin lifetimes have been reported in semiconductor QDs and a high degree of electron spin polarization has also been demonstrated through optical orientation under resonant excitation of QD excitons [13-15]. However, upon spin injection from adjacent layers such as a spin-filtering layer, a wetting layer (WL) and a barrier layer, that is a necessary step required for operation of most envisaged spintronic devices based on QDs, reported values of electron spin polarization degree of QDs remain very low so far [16-23]. The exact physical mechanism for the observed low spin injection efficiency is still not understood, unfortunately, which can originate from poor spin generation and spin loss along the path of spin injection. The former was generally believed to be associated with accelerated spin relaxation in WL and barrier layers when carrier and exciton motions (i.e. with a non-zero momentum k) promote spin-orbit mediated spin relaxation [21,24,25]. Up to now, WL and barrier layers surrounding QDs are commonly regarded as being of “ideal” 2D and 3D characters, respectively.
In this work we examine the exact path of spin injection from WL and barriers into QDs in self-assembled InAs/GaAs QDs and quantum-dot molecular structures (QMSs), including single quantum dots (SQDs), laterally-aligned double QDs (DQDs), quantum rings (QRs), and quantum clusters (QCs). Surprisingly, we find that the dominant exciton spin injection into QDs is in fact undertaken from localized regions of the WL and barrier layers immediately surrounding the QDs, in sharp contrast to the commonly believed 2D and 3D regions. The strong localization of excitons in such WL and barrier regions facilitates an
anisotropic exchange interaction (AEI), which leads to strong mixing of electron/exciton spin states and thereby reduces spin injection efficiency. We further show that the efficiency of spin injection is correlated with the “overall” geometric symmetry of the QMSs, implying a close relation between spin loss and the potential anisotropy within and around the QMSs. 2 Results and discussion In our studies, spin-polarized carriers and excitons were generated by employing the optical orientation technique [26] with the aid of circularly polarized light from a wavelength-tunable laser [18,27]. By tuning the excitation photon energy to match the bandgap energies of the WL and GaAs barrier, spin generation in these layers can be individually selected and subsequent spin injection into QDs and QMSs can be monitored by the resulting photoluminescence (PL) polarization from the QDs and QMSs. As spin relaxation is known to be much more severe for holes than electrons in the studied semiconductor system, spin injection is mainly determined by the electrons [6].
A typical PL spectrum from the studied QMS ensembles at 4K is shown in the right panel of Fig. 1, taking as an example the DQD ensemble with the excitation wavelength fixed at 810 nm (with the excitation photon energy above the bandgap energies of both WL and GaAs). The observed PL band arises from the excitonic recombination of the QMSs associated with the heavy-hole (hh) exciton ground state, under moderate excitation power to avoid saturation and state filling effect that could complicate polarization studies [28,29]. The broadening of the PL emission is caused by inhomogeneous effects, like fluctuations in QD size and random distributions of strain fields and indium composition. A PL excitation (PLE) spectrum, displayed in the left panel of Fig. 1, was obtained by detecting the PL intensity at the peak position of the PL band while scanning excitation wavelength across the bandgap energies of the WL and GaAs. The following two photo-generation and injection paths of the carriers/excitons into the QMSs can clearly be distinguished. The first path involves photo-generation of excitons and uncorrelated electron-hole pairs (referred to below as free carriers) from the InGaAs WL, represented
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by the hh and light-hole (lh) exciton peaks (denoted by XHWL and XLWL) and the continuum on the higher energy side of each exciton peak originating from the band-to-band (BB) optical transition, respectively. The observed large splitting between the XHWL and XLWL lines, with the former lying lower in energy, is consistent with the strong in-plane compressive strain and quantum confinement along the growth direction in the WL [30]. This large splitting has enabled us to focus on the excitonic effect on the spin injection from the XHWL alone. (The photo-excitation and spin injection from the XLWL resonance are much more complicated due to a strong overlap with the hh BB continuum and require a treatment including many-body effects, which is beyond the scope of the present work.) The second photo-generation and injection path of the carriers/excitons into the QMSs is via the GaAs barrier layer when the excitation photon energy is at and above the excitonic peaks of GaAs (denoted by XHGaAs and XLGaAs). As expected, it involves both the excitons (XHGaAs and XLGaAs) and free carriers. Thanks to the well-separated spectral ranges for the photo-excitation between the WL and GaAs barrier layer, and also between the exciton peaks and the BB continuum in each layer, spin generation/injection processes from these various channels can be compared and studied separately by choosing appropriate excitation wavelengths. We should note that spin injection from the GaAs layers into the QDs and QMSs can take place indirectly via the WL if spin-polarized excitons/carriers photo-generated in GaAs pass WL en route to QDs and QMSs, besides a direct injection process from the regions of the GaAs capping layers that are in direct contact with or within immediate surroundings of the QDs and QMSs. The indirect injection can induce additional spin loss, leading to a further reduced spin injection efficiency.
Under the optical orientation experiments, optical excitation of the WL and GaAs barrier layer with a given circular polarization such as will
generate spin-down electrons via the
hh excitons and the hh BB transitions, and spin-up
electrons via the lh excitons and the lh
BB transitions [26]. excitation does the
opposite, namely, generating spin-up electrons via
the hh excitons and the hh BB
transitions, and spin-down electrons via the
lh excitons and the lh BB transitions.
(Here the neutral exciton states are represented by
, where and are the magnetic
quantum numbers of the electron and hole, respectively.) If the electron spin is conserved during spin injection from the WL and GaAs barrier to the QMSs, at least partially, the PL emission from the QMSs becomes circularly polarized. As the ground-state hh excitons in the QMSs are dominated by the neutral excitons and positively charged trions due to residual C-acceptors in the GaAs layer, the -polarized PL are one-to-one associated with the spin-down
electron of the neutral exciton and the
positively charged trions.
Likewise, the -polarized PL is associated with
the spin-up electron of the neutral
excitons and the positively
charged trions. (Here the positively charged
exciton states are represented by .)
The degree of the circular polarization of the QMS PL, defined by
where
refers to the intensity of the or PL
component, thus measures the degree of electron spin polarization in the QMSs as a result of spin injection from the WL and GaAs barrier. In the studied QMSs, the evidence for spin injection from the WL and GaAs barrier is provided by the non-zero of the QMS PL. The degree of
is found to critically depend on the excitation energy, however, especially when the exciton spin injection channels are activated. The middle curves in Fig.2a-b clearly show that spin
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generation and injection via the excitons by resonantly exciting the exciton peaks XHWL consistently gives rise to a lower degree of electron spin polarization in the SQDs and DQDs as compared with that via the hh BB transition of the WL (denoted by ), evident from the dips in the excitation spectral dependent spectra of
. This reduced spin injection efficiency of the excitons is found to be significant even from the GaAs barrier layer in these two QMSs, as shown by the middle curves in Fig.3a-b. Interestingly, this effect becomes much weaker or even negligible in the QR structure, see the middle curves in Fig.2c and Fig3c. It should be noted that the BB transitions from the GaAs barrier layer shown in Fig.3 are contributed from both hh and lh bands, denoted by and . Under excitation with a given circular polarization, these two channels generate electrons of opposite spin orientation as discussed above. This leads to co-polarized and counter-polarized PL of the QRs under the and excitation, respectively, which compensate each other giving rise to the measured . Taking into account that the contribution from the should be three times larger than that from the
due to their different oscillator strengths [26], the actual optical (or spin) polarization created by the spin injection from the alone is twice of the measured value as indicated by the dot-dash curves in Fig.3c. This reveals that the spin injection via the XHGaAs is actually less efficient than that via the even in the QR structure. Following this argument, the same conclusion can also be drawn for the exciton and free carrier injection involving the lh alone in the GaAs layer. However, it is also apparent that the spin loss via the exciton spin injection is markedly reduced in the QR structure as compared with the SQD and DQD structures, comparing the middle curves between Fin.3c and Fig.3a-b. In fact, there seems to be a clear dependence of the exciton spin injection efficiency on the overall geometric arrangement of the QMSs. This indicates an important role of local microstructures surrounding the QMSs on the spin properties of the WL and GaAs barrier, which cannot be described by the ideal picture of a “globally” 2D and 3D character for the WL and the GaAs barrier where the excitons and free carriers
are delocalized and are injected into all QMSs. The deviation of the GaAs barrier from a 3D character is also apparent from the removal of the degeneracy of the hh and lh excitons, i.e. XHGaAs and XLGaAs, as seen from the close-up PLE spectra (see the top curves in Fig.3). This means that the regions of the GaAs barrier layer that are active in the exciton spin injection must experience a strain field imposed by the WL or/and the QMSs. Consistently, the in-plane strain field in the concerned GaAs layer should be predominantly tensile, which is supported by the reverse ordering of the XHGaAs and XLGaAs in energy as compared to that in the WL and the QDs (or QMSs). More specifically, the energy of the lh exciton XLGaAs is lower than that of the hh exciton XHGaAs in the spin-injecting regions of the GaAs barrier layer as seen from the counter- and co-polarized PL of the QRs under the resonant circularly polarized excitation of the XLGaAs and XHGaAs (see the middle curves in Fig.3c).
To closely examine the microscopic nature of the WL and GaAs barrier layer directly participating in the spin injection processes, we carried out PLE measurements on single QDs and QMSs. This was done by focusing the laser beam to an area of ~1 in a confocal µ-PL system. As a lower number of the QDs and QMSs were monitored, the broad PL emission band observed in the ensemble measurements broke down into sharp lines with a typical linewidth of ~20 , as shown in Fig.4a. Each line corresponds to an excitonic transition from a single QD or QMS. Surprisingly, the PLE spectrum of a single QD or QMS close to the XHWL consists of many sharp lines, see Fig.4b. This finding reveals a strongly disordered nature of the WL where the excitons, contrary to being of a delocalized 2D character with a single fixed exciton energy seen by all QDs or QMSs in a given structure, are in fact confined to potential fluctuations arising from strong fluctuations in In composition and strain field. The strongly spatial localizations of the excitons are further confirmed by different PLE lines observed in different QDs and QMSs from the same structure, as shown in Fig.4b. These results conclude that there are no delocalized 2D excitons in the WL that are mobile and can be captured by all QDs or QMSs, at least not within the monitored area of ~1 . The WL excitons directly involved
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in the spin injection must then be localized within the immediate surroundings of the monitored QD or QMS, such that they should experience subtle differences between different QDs or QMSs leading to the observed different localization energies of these WL excitons near different QDs or QMSs.
Since the XHWL is shown to be localized, their electronic structure and spin properties can to a certain extent be compared to neutral excitons in QDs, where local symmetry plays a vital role. Indeed, as it has been shown both theoretically and experimentally for an ideal QD with a cylindrical symmetry, the hh exciton ground state contains a
degenerate bright doublet composed
of the spin-pure eigenstates for the electron
( ) and the hole ( ). Resonant
excitation with circularly polarized light (or
) can selectively generate the (or
) exciton with pure spin states. In a
realistic case, however, the exciton confinement potential in a QD is anisotropic as a result of anisotropy in QD shape, strain fields as well as atomistic asymmetry [31-33]. This activates the so-called anisotropic exchange interaction (AEI), enhanced in QDs due to strong confinement, which becomes the dominant mechanism for spin relaxation in QDs. AEI leads to a fine-structure splitting (FSS) of the bright exciton doublet
into two spin-mixed states
and
, which are
linearly polarized along the orthogonal x- and y-axis defined by the elongation axis of the exciton confinement potential [34,35]. To examine if the AEI plays a role in the XHWL, we performed PLE of the QDs and QMSs by employing linearly polarized excitation across the spectral range of XHWL. When the linear polarization axis of the excitation light is along the x-axis (or y-axis), denoted by (or ), only the (or ) exciton will be created by light. In our experiments
we applied linearly polarized excitation with its polarization axis alternating in time between the two orthogonal axes of [110] to [1-10], because these crystallographic axes are known to be the preferred elongation axes of the studied SQDs and DQDs from our polarized µ-PL studies of single QDs and single QMSs [35]. The resulting total PL intensity of the QDs or QMSs is detected by the lock-in technique in phase with the modulation frequency of the linearly polarized excitation, yielding a linear polarization degree of PLE by
where (or ) refer to the total PL intensity under the
(or ) excitation. As x=[110] and y=[1-10] in our case, a positive (or negative) characterizes the exciton state with its optical dipole along the [110] (or [1-10]) axis. In the presence of a non-vanishing FSS, the positive and negative signals will appear in two different excitation photon energies leading to a derivative shape of as a function of excitation photon energy. [Due to the presence of some background signals in , due to either a small instrumental calibration error or some physical mechanism such as hh-lh mixing, it is the derivative shape (not the absolute value) in the
spectra that reflects the polarization states of the excitons.] Indeed, measured from the DQD and SQD samples as shown by the lowest curves in Fig.2a-b clearly exhibits a derivative shape across the spectral range of the XHWL. In contrast, no clear derivative shape of can be observed across the spectral range of the XHWL in the QR structure. These results provide evidence for a sizable AEI of the XHWL in the SQD and DQD structures, which is suppressed in the QR structure. Correlating with the circular polarization results shown by the middle curves in Fig.2 (discussed above), it becomes apparent that the observed loss of the exciton spin injection from the WL is directly linked to the AEI. In other words, AEI plays a dominant role in the exciton spin relaxation in the WL. In fact, this is found to be a general trend and spin injection loss through the GaAs excitons can be attributed to AEI as well. The experimental evidence for the presence of AEI even in the spin injection regions of the GaAs layer is provided by the derivative features in the spectra, shown by the lowest curves in Fig. 3. They are the
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result of a combined effect of the AEI on XHGaAs and XLGaAs, as each is expected to undergo splitting into linearly polarized exciton sublevels under the influence of the AEI. The measured PLE spectra are the sum of the two contributions, and consequently the PLE line shape critically depends on the extent of the spectral overlap between the two.
We should point out that the effect of the exciton spin relaxation facilitated by the AEI is not only restricted to exciton spin injection. It can indirectly affect spin injection efficiency via free carriers, if trapping of free carriers into exciton states strongly competes with direct spin injection of free carriers into QDs or QMSs. The latter is particularly important at low temperatures.
The observed correlation between FSS and a decrease of identifies the AEI of the WL and GaAs barrier as the dominant mechanism that limits efficiency of the exciton spin injection into the QDs and QMSs. Our earlier studies of AEI for neutral excitons of the QMSs in the same set of structures showed a strong dependence of the AEI on the geometric structures of the QMSs [35], in an identical fashion as what was observed here for the excitons in the spin-injecting WL and GaAs layer. Therefore we envision that the AEI experienced by the excitons both within and in the immediately surroundings of the QMSs is most likely of the same physical origin. Our earlier studies concluded that the anisotropic strain field induced by an anisotropic distribution of In atoms plays an important role in the observed AEI of the excitons confined within the QMSs, which varies from a strong and directional strain field in the “low-symmetry” DQDs to a weak, randomly oriented strain field in the “higher-symmetry” QRs and QCs [35]. [We should note that, though the AEI-induced exciton spin relaxation in the QMSs can lead to an overall reduction in the absolute values of the measured spin polarization, it will not contribute to the variations between exciton and free-carrier spin generation from the WL and GaAs layers seen in the macro-PLE spectra (Figs.2-3) under either linearly or circularly polarized excitation with photon energy over the spectral range of the WL and GaAs layers. The macro-PLE spectra were obtained by monitoring the PL at the peak position of the PL band from the QMSs, which is contributed by all FSS components from many QMSs.] As the anisotropic strain field is
imposed on a QMS by its surroundings, we can reasonably assume that the WL and GaAs barrier regions immediately surrounding the QMS should in return experience a similar anisotropic strain field but with an opposite sign. It is this anisotropic strain field that has led to the AEI and its associated exciton spin relaxation in the WL and GaAs layers. Our results thus show that growing QMSs with a high symmetry is a promising route to achieve not only “ideal” QMSs for spin readout and quantum information storage/processing but also desirable WL and GaAs barrier layer for efficient spin injection. 3 Conclusions In summary, we have studied exciton spin injection in various InAs/GaAs QD and QMS structures. From both WL and GaAs barrier layer, the spin injection via excitons was consistently found to be less efficient as compared with that via free carriers. We uncovered the exact path of the exciton spin injection as being via localized excitons confined to local potential fluctuations in the regions of the injection layers immediately surrounding the QDs or QMSs. This is based on our experimental finding of a group of excitons (with different energies) injected from the WL into a single QD or a single QMS that are unique for each QD or QMS. We showed that the exciton localization facilitated AEI, as evident from the FSS of the WL and GaAs excitons revealed in the polarization-resolved PLE studies of the QD and QMS structures. AEI provided the dominant driving force for exciton spin relaxation leading to the measured lower spin injection efficiency. We suggested that the AEI in the WL and GaAs barrier is predominantly induced by anisotropic strain fields, similar to the AEI of the neutral excitons in the QMSs, as the strain field imposed on a QMS and its immediate surroundings is expected to be mutual. Furthermore, the AEI-induced reduction in the exciton spin injection efficiency was found to be significantly suppressed from the “low-symmetry” DQDs to the “high-symmetry” QRs. Our results provided an insight into the exact spin injection path and, more importantly, shed light on the dominant physical mechanism responsible for the exciton spin injection loss in the self-assembled QD structures. This work also demonstrated that properly arranged QMSs could be a promising
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pathway for optimal spin injection and spin detection. 4 Methods The samples investigated are a set of self-organized InAs/GaAs QDs and QMSs including SQDs, laterally-aligned DQDs, QCs and QRs. Spontaneous formation of the InAs QDs and QMSs in the Stranski-Krastanov (SK) growth mode took place after deposition of an InAs wetting layer by molecular bean epitaxy (MBE). The QMSs were obtained during a recapping and overgrowth process following the growth of QDs under varying growth conditions. The QRs were formed by partially capping of InAs QDs by 6 monolayers (MLs) of GaAs at a substrate temperature of 470 under overpressure. By depositing an additional 0.6 ML of InAs on the QRs while keeping the same growth temperature, InAs DQDs were formed. However, if the substrates were heated to 520 , the QRs were transformed to QCs instead. An additional GaAs capping layer was deposited to protect the QMSs and to terminate the growth process. The protective capping layer and the WL serves as natural barriers adjacent to the QMSs where the studied spin injection was initiated. The details of the growth procedure and a description of the QMS samples including a thorough µ-PL spectroscopy study with a focus on tailoring AEI and exciton FSS can be found in Ref. [36] and Ref. [35].
PL and PLE of the QMS ensembles were measured in a usual back-scattering geometry with samples kept in a continuing flow liquid-helium cryostat at 4K, under optical excitation with a tunable Ti:sapphire laser. The excitation laser beam was directed normal to the sample surface with polarization, controlled using a variable liquid crystal wave plate, to be either circular polarized ( or ) to fulfill the optical orientation condition or modulated between linear polarization along two orthogonal <110> directions ( and ) perpendicular to the growth direction at a frequency of . The resulting PL was spectrally resolved by a single-grating monochromator and collected by a liquid-nitrogen-cooled Ge detector. µ-PL was performed at 5K with a confocal system (Horiba-Jobin HR800) in the backscattering
geometry, under optical excitation with a wavelength-tunable Ti-sapphire laser. The excitation spot was ~1 µm in diameter. The optical excitation power was kept low to avoid contributions from excited states, multi-exciton complexes and negatively charged trions [29]. The µ-PL signal was dispersed by a high-resolution spectrometer, with a spectral resolution of about 5 µeV, and detected by a charge-coupled-device (CCD) detector. µ-PLE was obtained by registering µ-PL spectra while varying excitation wavelength, which was successful for the XHWL and revealed that the broader PLE line of the QMS ensembles consisted of many sharp lines. Due to the limited step resolution of the tunable Ti:sapphire laser and the resulting limited number of data points in a µ-PLE spectrum across the relatively narrow lines of XHGaAs and XLGaAs, it was unfortunately not possible to reliably resolve sharp lines for the GaAs layer.
Acknowledgements
We thank C.W. Tu for providing the samples. This work was supported by Linköping University through the Professor Contracts, the Swedish Research Council (Grant No. 621-2011-4254), the Linköping Linnaeus Initiative for Novel Functional Materials (LiLI-NFM) supported by the Swedish Research Council (contract number 2008-6582).
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[22] Beyer, J.; Buyanova, I. A; Suraprapapich, S.; Tu, C. W.; Chen W. M. The Hanle Effect and Electron Spin Polarization in InAs/GaAs Quantum Dots Up To Room Temperature. Nanotechnology 2012, 23, 135705.
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Figure 1 The right panel: PL spectrum from the DQD sample
measured at 4K with the excitation wavelength of 810 nm
and power of 100 mW. The left panel: PLE from the same
sample by detecting at the peak position of the PL emission
band. XHWL, XLWL, XHGaAs and XLGaAs denote the hh and lh
excitons from the WL and GaAs barrier layer. ,
Figure 2 Excitation-photon energy dependence (around the
XHWL) of the PL intensity (the top curves), the PL circular
polarization degree under the (the red
curves in the middle panel) and (the blue curves in the
middle panel) excitation and the PLE linear polarization
degree (the lowest curves), obtained at 4K from the
SQD (a), DQD (b) and QR (c) structures by monitoring the
Figure 3 Excitation-photon energy dependence (around the
XHGaAs and XLGaAs) of the PL intensity (the top curves), the
PL circular polarization degree under the
(the red curves in the middle panel) and (the blue curves
in the middle panel) excitation and the PLE linear
polarization degree (the lowest curves), obtained at
www.theNanoResearch.com∣www.Springer.com/journal/12274 | Nano Research
10 Nano Res.
Figure 4 (a) µ-PL intensity as a function of both detection
and excitation wavelength from the QC structure. The
excitation spectral range is chosen to be close to the XHWL.
(b) µ-PLE spectra in the vicinity of the XHWL, measured
from two individual QCs with their PL emission lines marked
in (a).
11
Silver Nanowires with Semiconducting Ligands for Low Temperature Transparent Conductors
Brion Bob,1 Ariella Machness,1 Tze-Bin Song,1 Huanping Zhou,1 Choong-Heui Chung,2 and Yang Yang1,*
1 Department of Materials Science and Engineering and California NanoSystems Institute,
University of California Los Angeles, Los Angeles, CA 90025 (USA)
2 Department of Materials Science and Engineering, Hanbat National University, Daejeon
305-719, Korea
Abstract
Metal nanowire networks represent a promising candidate for the rapid fabrication of transparent electrodes with high transmission and low sheet resistance values at very low deposition temperatures. A commonly encountered obstacle in the formation of conductive nanowire electrodes is establishing high quality electronic contact between nanowires in order to facilitate long range current transport through the network. A new system of nanowire ligand removal and replacement with a semiconducting sol-gel tin oxide matrix has enabled the fabrication of high performance transparent electrodes at dramatically reduced temperatures with minimal need for post-deposition treatments of any kind.
1. Introduction. Silver nanowires (AgNWs) are long, thin, and possess conductivity values on the same order of magnitude as bulk silver
(Ag) [1]. Networks of overlapping nanowires allow light to easily pass through the many gaps and spaces between nanowires, while transporting current through the metallic conduction pathways offered by the wires themselves. The high aspect ratios achievable for solution-grown AgNWs has allowed for the fabrication of transparent conductors with very promising sheet resistance and transmission values, often approaching or even surpassing the performance of vacuum-processed materials such as indium tin oxide (ITO) [2-6].
Significant electrical resistance within the metallic nanowire network is encountered only when current is required to pass between nanowires, often forcing it to pass through layers of stabilizing ligands and insulating materials that are typically used to assist with the synthesis and suspension of the nanowires [7, 8]. The resistance introduced by the insulating junctions between nanowires can be reduced through various physical and chemical means, including burning off ligands and partially melting the wires via thermal annealing [9, 10], depositing additional materials on top of the nanowire network [11-14], applying mechanical forces to enhance network morphology [15-17], or using various other post-treatments to improve the contact between adjacent wires [18-21]. Any attempt to remove insulating materials the network must be weighed against the risk of damaging the wires or blocking transmitted light, and so many such treatments must be reined in from their full effectiveness to avoid endangering the performance of the completed electrode.
We report here a process for forming inks with dramatically enhanced electrical contact between AgNWs through the use of a semiconducting ligand system consisting of tin oxide (SnO2) nanoparticles. The polyvinylpyrrolidone (PVP) ligands introduced during AgNW synthesis in order to encourage one-dimensional growth are stripped from the wire surface using ammonium ions, and are replaced with substantially more conductive SnO2, which then fills the space between wires and enhances the contact geometry in the vicinity of wire/wire junctions. The resulting transparent electrodes are highly conductive immediately upon drying, and can be effectively processed in air at virtually any temperature below 300 °C. The capacity for producing high performance transparent electrodes at room temperature may be useful in the fabrication of devices that are damaged upon significant heating or upon the application of harsh chemical or mechanical post-treatments.
2. Results and Discussion
2.1. Ink Formulation and Characterization
Dispersed AgNWs synthesized using copper chloride seeds represent a particularly challenging material system for promoting wire/wire junction formation, and often require thermal annealing at temperatures near or above 200 °C to induce long range electrical conductivity within the deposited network [22, 23]. The difficulties that these wires present regarding junction formation is potentially due to their relatively large diameters compared to nanowires synthesized using other seeding materials, which has the capacity to enhance the thermal stability of individual wires according to the Gibbs-Thomson effect. We have chosen these wires as a demonstration of pre-deposition semiconducting ligand substitution in order to best illustrate the contrast between treated and untreated wires.
Completed nanocomposite inks are formed by mixing AgNWs with SnO2 nanoparticles in the presence of a compound capable of stripping the ligands from the AgNW surface. In this work, we have found that ammonia or ammonium salts act as effective stripping agents that are able to remove the PVP layer from the AgNW surface and allow for a new stabilizing matrix to take its place. Figure 1 shows a schematic of the process, starting from the precursors used in nanowire and nanoparticle synthesis and ending with the deposition of a completed film. The SnO2 nanoparticle solution naturally contains enough ammonium ions from its own synthesis to effectively peel the insulating ligands from the AgNWs and allow the nanoparticles to replace them as a stabilizing agent. If not enough SnO2 nanoparticles are used in the mixture, then the wires will rapidly agglomerate and settle to the bottom as large clusters. Large amounts of SnO2 in the mixture gradually begin to increase the sheet resistance of the nanowire network upon deposition, but greatly enhance the uniformity, durability, and wetting properties of the resulting films. We have found that AgNW:SnO2 weight ratios ranging between 2:1 and 1:1 produce well dispersed inks that are still highly conductive when deposited as films.
The nanowires were synthesized using a polyol method that has been adapted from the recipe described by Lee et al. [22, 23] Silver nitrate dissolved in ethylene glycol via ultrasonication was used as a precursor in the presence of copper chloride and PVP to provide seeds and produce anisotropic morphologies in the reaction products. Synthetic details can be found in the experimental section. Distinct from previous recipes, we have found that repeating the synthesis two times without cooling down the reaction mixture generally produces significantly longer nanowires than a single reaction step. The lengths of nanowires produced using this method fall over a wide range from 15 to 65 microns, with diameters between 125 and 250 nm. This range of diameters is common for wires grown using copper chloride seeds, although the double reaction produces a number of wires with roughly twice their usual diameter. The morphology of the as-deposited AgNWs as determined via SEM is shown in Figure 2(a), higher magnification images are also provided in Figures 2(c) and 2(d).
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The SnO2 nanoparticles were synthesized using a sol-gel method typical for multivalent metal oxide gelation reactions. A large excess of deionized water was added to SnCl4·5H2O dissolved in ethylene glycol along with tetramethylammonium chloride and ammonium acetate to act as surfactants. The reaction was then allowed to progress for at least one hour at near reflux conditions, after which the resulting nanoparticle dispersion can be collected, washed, and dispersed in a polar solvent of choice. The material properties of SnO2 nanoparticles formed using a similar synthesis method have been reported previously [24], although the present recipe uses excess water to ensure that the hydrolysis reaction proceeds nearly to completion.
After mixing with SnO2 nanoparticles, films deposited from AgNW/SnO2 composite inks show a largely continuous nanoparticle layer on the substrate surface with some nanowires partially buried and some sitting more or less on top of the film. Representative scanning electron microscopy (SEM) images of nanocomposite films are shown in Figure 2(b). Regardless of their position relative to the SnO2 film, all nanowires show a distinct shell on their outer surface that gives them a soft and slightly rough appearance, as is visible in the higher magnification images shown in Figure 2(e) and 2(f). The SnO2 nanoparticles do a particularly good job coating the regions near and around junctions between wires, and frequently appear in the SEM images as bulges wrapped around the wire/wire contact points.
The precise morphology of the SnO2 shell that effectively surrounded each AgNW was analyzed in more detail using transmission electron microscopy (TEM) imaging. Figures 3(a) to 3(c) show individual nanowires in the presence of different ligand systems: as-synthesized PVP in Figure 3(a), inactive SnO2 in Figure 3(b), and SnO2 activated with trace amounts of ammonium ions in Figure 3(c). The as-synthesized nanowires show sharp edges, and few surface features. In the presence of inactive SnO2, which is formed by repeatedly washing the SnO2 nanoparticles in ethanol until all traces of ammonium ions are removed, the nanowires coexist with somewhat randomly distributed nanoparticles that deposit all over the surface of the TEM grid. When AgNWs are mixed with activated SnO2, a thick and continuous SnO2 shell is formed along the nanowire surface. In when sufficiently dilute SnO2 solutions are used to form the nanocomposite ink, nearly all of the nanoparticles are consumed during shell formation and effectively no nanoparticles are left to randomly populate the rest of the image.
As the AgNWs acquire their metal oxide coatings in solution, the properties of the mixture change dramatically. Freshly synthesized AgNWs coated with residual PVP ligands slowly settle to the bottom of their vial or flask over a time period of several hours to one day, forming a dense layer at the bottom. The AgNWs with SnO2 shells do not settle to the bottom, but remain partially suspended even after many weeks at concentrations that are dependent on the amount of SnO2 present in the solution.
A comparison of the settling behavior of various AgNW and SnO2 mixtures after 24 hours is shown in Figures 3(d) and 3(e). The ratios 8:4, 8:16, and 8:8 indicate the concentrations of AgNWs and SnO2 (in mg/mL) present in each solution. The 8:8 uncoupled solution, in which the PVP is not removed from the AgNW surface with ammonia, produces a situation in which the nanowires and nanoparticles do not interact with one another, and instead the nanowires settle as in the isolated nanowire solution while the nanoparticles remain well-dispersed as in the solution of pure SnO2. The mixtures of nanowires and nanoparticles in which trace amounts of ammonia are present do not settle to the bottom, but instead concentrate themselves until repulsion between the semiconducting SnO2 clusters is able to prevent further settling.
Our current explanation for the settling behavior of the wire/particle mixtures is that the PVP coating on the surface of the as-synthesized wires is sufficient to prevent interaction with the nanoparticle solution. The addition of ammonia into the solution quickly strips off the PVP surface coating and allowing the nanoparticles to coordinate directly with the nanowire surface. This explanation is in agreement with the effects of ammonia has on a solution of pure AgNWs, which rapidly begin to agglomerate into clusters and sink to the bottom as soon as any significant quantity of ammonia is added to the ink.
We attribute the stripping ability of ammonia in these mixtures to the strong dative interactions that
occur via the lone pair on the nitrogen atom interacting with the partially filled d-orbitals of the Ag atoms
on the nanowire surface. These interactions are evidently strong enough to displace the existing
coordination of the five-membered rings and carbonyl groups contained in the original PVP ligands and
allow the ammonia to attach directly to the nanowire surface. Since ammonia is one of the original
surfactants used to stabilize the surface of the SnO2 nanoparticles, we consider it reasonable that ammonia
coordination on the nanowire surface would provide an appropriate environment for the nanoparticles to
adhere to the AgNWs.
14
Scanning Energy Dispersive X-ray (EDX) Spectroscopy was also conducted on nanoparticle-coated AgNWs in order to image the presence of Sn and Ag in the nanowire and shell layer. The line scan results are shown in Figure 3(f), having been normalized to better compare the widths of the two signals. The visible broadening of the Sn lineshape compared to that of Ag is indicative of a Sn layer along the outside of the wire. The increasing strength of the Sn signal toward the center of the AgNW is likely due to the enhanced interaction between the TEM’s electron beam and the dense AgNW, which then improves the signal originating from the SnO2 shell as well. It is also possible that there is some intermixing between the Ag and Sn x-ray signals, but we consider this to be less likely as the distance between their characteristic peaks should be larger than the detection system’s energy resolution.
2.2. Network Deposition and Device Applications
For the deposition of transparent conducting films, a weight ratio of 2:1 of AgNWs to SnO2 nanoparticles was chosen in order to obtain a balance between the dispersibility of the nanowires, the uniformity of coated films, and the sheet resistance of the resulting conductive networks. Nanocomposite films were deposited on glass by blade coating from an ethanolic solution using a scotch tape spacer, with deposited networks then being allowed to dry naturally in air over several minutes.
The as-dried nanocomposite films are highly conductive, and require only minimal thermal treatment to dry and harden the film. Without the use of activated SnO2 ligands, deposited nanowire networks are highly insulating, and become conductive only after annealing at above 200 °C. The sheet resistance values of representative films are shown in Figure 4(a). The capability to form transparent conductive networks in a single deposition step that remain useful over a wide range of processing temperatures provides a high degree of versatility for designing thin film device fabrication procedures.
Figure 5(a) shows the sheet resistance and transmission of a number of nanocomposite films deposited from inks containing different nanowire concentrations. The deposited films show excellent conductivity at transmission values up to 85%, and then rapidly increase in sheet resistance as the network begins to reach its connectivity limit. The optimum performance of these networks at low to moderate transmission values is a consequence of the relatively large nanowire diameters, which scatter a noticeable amount of light even when the conditions required for current percolation are just barely met. Nonetheless, the sheet resistance and transmission of the completed nanocomposite networks place them within an acceptable range for applications in a variety of optoelectronic devices. Figure 5(b) shows the wavelength dependent transmission spectra of several nanowire networks, which transmit light well out into the infrared region. The presence of high transmission values out to wavelengths well above 1300 nm, where ITO or other conductive oxide layers would typically begin to show parasitic absorption, is due to the use of semiconducting SnO2 ligands, which is complimentary to the broad spectrum transmission of the silver nanowire network itself.
Avoiding the use of highly doped nanoparticles has the potential to provide optical advantages, but can create difficulties when attempting to make electrical contact to neighboring device layers. In order to investigate their functionality in thin film devices, we have incorporated AgNW/SnO2 nanocomposite films as electrodes in amorphous silicon (a-Si) solar cells. Two contact structures were used during fabrication: one with the nanocomposite film directly in contact with the p-i-n absorber structure and one with a 10 nm Al:ZnO (AZO) layer present to assist in forming Ohmic contact with the device. The I-V characteristics of the resulting devices are shown in Figure 6(a).
The thin AZO contact layers typically show sheet resistance values greater than 2.5 kΩ/⧠, and so cannot be responsible for long range lateral current transport within the electrode structure. However, their presence is clearly beneficial in improving contact between the nanocomposite electrode and the absorber material, as the SnO2 matrix material is evidently not conductive enough to form a high quality contact with the p-type side of the a-Si stack. We hope that future modifications to the AgNW/SnO2 composite, or perhaps the use of islands of high conductivity material such as a discontinuous layer of doped nanoparticles will allow for the deposition of completed electrode stacks that provide both rapid fabrication and good performance.
Figure 6(b) contains the top view image of a completed device. The enhanced viscosity of the nanowire/sol-gel composite inks allows for films to be blade coated onto substrates with a variety of surface properties without reductions in network uniformity. In contrast with traditional back electrodes deposited in vacuum environments, the nanocomposite can be blade coated into place in a single pass under atmospheric conditions and dried within moments. We anticipate that the use of sol-gel mixtures to enhance wetting and dispersibility may prove useful in the formulation of other varieties of semiconducting and metallic inks for deposition onto a variety of substrate structures.
3. Conclusions
In summary, we have successfully exchanged the insulating ligands that normally surround as-synthesized AgNWs with shells of substantially more conductive SnO2 nanoparticles. The exchange of one set of ligands for the other is mediated by
15
the presence of ammonia during the mixing process, which appears to be necessary for the effective removal of the PVP ligands that initially cover the nanowire surface. The resulting nanowire/nanoparticle mixtures allow for the deposition of nanocomposite films that require no annealing or other post-treatments to function as high quality transparent conductors with transmission and sheet resistance values of 85% and 10 Ω/⧠, respectively. Networks formed in this manner can be deposited quickly and easily in open air, and have been demonstrated as an effective n-type electrode in a-Si solar cells when a thin interfacial layer is deposited first to ensure good electronic contact with the rest of the device. The ligand management strategy described here could potentially be useful in any number of material systems that presently suffer from highly insulating materials that reside on the surface of otherwise high performance nano and microstructures.
4. Experimental Details
Tin oxide nanoparticle synthesis. Tin chloride pentahydrate was dissolved in ethylene glycol by
stirring for several hours at a concentration of 10 grams per 80 mL to serve as a stock solution. In a typical
synthesis reaction, 10 mL of the SnCl4·5H2O stock solution is added to a 100 mL flask and stirred at room
temperature. Still at room temperature, 250 mg ammonium acetate and 500 mg ammonium acetate were
added in powder form to regulate the solution pH and to serve as coordinating agents for the growing
oxide nanoparticles. 30 ml of water was then added, and the flask was heated to 90 °C for 1 to 2 hours in
an oil bath, during which the solution took on a cloudy white color. The gelled nanoparticles were then
washed twice in ethanol in order to keep trace amounts of ammonia present in the solution. Additional
washing cycles would deactivate the SnO2, and then require the addition of ammonia to coordinate with
as-synthesized AgNWs.
Silver nanowire synthesis. Copper(ii) chloride dihydrate was first dissolved in ethylene glycol at
1 mg/ml to serve as a stock solution for nanowire seed formation. 20 ml of ethylene glycol was then added
into a 100 ml flask, along with 200 µL of copper chloride solution. the mixture was then heated to 150 °C
while stirring at 325 rpm, and .35g of PVP (MW 55,000) was added. In a small separate flask, .25 grams of
silver nitrate was dissolved in 10 ml ethylene glycol by sonicating for approximately 2 minutes, similar to
the method described here.22 The silver nitrate solution was then injected into the larger flask over
approximately 15 minutes, and the reaction was allowed to progress for 2 hours. After the reaction had
reached completion, the various steps were repeated without cooling down. 200 µL of copper chloride
solution and .35g PVP were added in a similar manner to the first reaction cycle, and another .25g silver
nitrate were dissolved via ultrasonics and injected over 15 minutes. The second reaction cycle was allowed
to progress for another 2 hours, before the flask was cooled and the reaction products were collected and
washed three times in ethanol.
Nanocomposite ink formation. After the synthesis of the two types of nanostructures is complete,
16
the double washed SnO2 nanoparticles and triple-washed nanowires can be combined at a variety of weight
ratios to form the completed nanocomposite ink. The dispersibility of the mixture is improved when more
SnO2 is used, although the sheet resistance of the final networks will begin to increase if they contain
excessive SnO2. AgNW agglomeration during mixing is most easily avoided if the SnO2 and AgNW
solutions are first diluted to the range of 10 to 20 mg/ml in ethanol, with the SnO2 solution being added
first to an empty vial and the AgNW solution added afterwards. The dilute mixture was then be allowed to
settle overnight, and the excess solvent removed to concentrate the wires to a concentration that is
appropriate for blade coating.
Film and electrode deposition. The completed nanocomposite ink was deposited onto any desired
substrates using a razor blade and scotch tape spacer. The majority of the substrates used in this study were
Corning soda lime glass, but the combined inks also deposited well on silicon, SiO2, and any other
substrates tested. Electrode deposition onto a-Si substrates was accomplished by masking off the desired
cell area with tape, and then depositing over the entire region. The p-i-n a-Si stacks and 10 nm AZO
contact layers were deposited using PECVD and sputtering, respectively.
ACKNOWLEDGMENTS The authors would like to acknowledge the use of the Electron Imaging Center for Nanomachines
(EICN) located in the California NanoSystems Institute at UCLA.
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