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C: Physical Processes in Nanomaterials and Nanostructures

Imaging the Reduction of Electron Trap States in Shelled CIGSe Nanocrystals Using Ultrafast Electron Microscopy Gabriele Meizyte, Riya Bose, Aniruddha Adhikari, Jun Yin, Md Azimul Haque, Manas R. Parida, Mohamed Nejib Hedhili, Tom Wu, Osman M. Bakr, and Omar F. Mohammed J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b03872 • Publication Date (Web): 01 Jun 2018 Downloaded from http://pubs.acs.org on June 1, 2018

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The Journal of Physical Chemistry

Imaging the Reduction of Electron Trap States in Shelled CIGSe Nanocrystals Using Ultrafast Electron Microscopy Gabriele Meizyte,† Riya Bose,† Aniruddha Adhikari,† Jun Yin,† Md Azimul Haque,ǂ Manas R. Parida,† Mohamed N. Hedhili,# Tom Wu,ǂ, Osman M. Bakr,† Omar F. Mohammed†* †

KAUST Solar Center, Physical Sciences and Engineering Division, King Abdullah University

of Science and Technology, Thuwal 23955-6900, Kingdom of Saudi Arabia ǂ

Physical Sciences and Engineering Division, King Abdullah University of Science and

Technology, Thuwal 23955-6900, Kingdom of Saudi Arabia #Core

Labs, King Abdullah University of Science and Technology, Thuwal 23955-6900, Kingdom

of Saudi Arabia

School

of Materials Science and Engineering, University of New South Wales (UNSW), Sydney,

NSW 2052, Australia

ACS Paragon Plus Environment

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ABSTRACT. Insertion of alkali metal ions, especially Na, is a well-established method to significantly increase the power conversion efficiency of copper indium gallium selenide (CIGSe) based photovoltaic devices. However, although it is known that Na ions mostly resides on the surface of CIGSe layer following diffusion, the exact mechanism of how Na affects the carrier dynamics of CIGSe still remains ambiguous. This is mainly due to the unavailability of suitable surface-sensitive techniques. Herein, we employ four dimensional scanning ultrafast microscopy (4D S-UEM), which has the unique capability of mapping the charge carrier dynamics in real time and space selectively on the materials surfaces, to directly observe the effect of Na insertion on the carrier dynamics of shelled CIGSe film. It is found that an additional layer of NaF to the thin film of ZnS shelled CIGSe nanocrystals not only increases the grain size and improves the texture of the film, but more importantly, reduces fast electron trap channels further on the surface of the material, as observed from the secondary electron dynamics in 4D S-UEM. Our density functional theory (DFT) calculations further confirm that Na ions can occupy Cu vacancies and reduce the interfacial charge carrier-defect scatterings. Removal of such undesirable electron trapping channels results in increased photoconductivity of the material, thereby serving as one of the critical parameters that lead to enhancement of the efficiency of CIGSe for light harvesting purposes.


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INTRODUCTION In the search for suitable photoactive materials to harvest solar energy, copper indium gallium selenide (CIGSe) has been one of the most promising candidates yielding power conversion efficiency >20%, because of its excellent light absorption capabilities and long term stability. 1-3 However, fabrication of the high efficiency devices involves the synthesis of CIGSe via vacuum based co-evaporation approach along with processing at high temperature (>500 oC) in presence of Se atmosphere. This expensive manufacturing process serves as the bottleneck impeding widespread commercialization of CIGSe based solar cells. An alternative approach to resolve this issue is the use of CIGSe nanocrystals, which can be synthesized by easy solution based processes and can also offer precise control over the composition and bandgap of the material.4-9 However, CIGSe nanocrystal based devices never reached similar high efficiency as obtained by thin films, mainly because of the presence of a large number of surface traps that serve as undesirable recombination channels for the charge carriers in the nanocrystals.10,11 A widely accepted way to improve the efficiency of CIGSe based photovoltaic devices is the addition of a small amount of alkali metal, typically Na, in the CIGSe layer.5,12-19 Incorporation of Na is achieved either directly via Na diffusion from the sodalime glass substrate, or by adding Nacontaining compounds, e.g., NaF, NaCl, Na2S, Na2Se, before or after the deposition of CIGSe films. This surface treatment improves the open circuit voltage (VOC) and fill factor (FF) of the device.5,12-21 Depending on the deposition technique, several mechanisms have been proposed for the efficiency enhancement, which can be broadly classified into two categories- a) influence on growth kinetics that leads to different grain sizes, smoother texture and suppression of defect formation22-25 and b) improvement of electronic properties by removal of deleterious donor states and introduction of acceptor states, which in turn increases the net hole concentration, and hence


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p-type conductivity.19,21,26 In addition, Na is reported to passivate grain boundary defects and reduce minority carrier traps.13,14 However, the exact reason for such efficiency enhancement remains under debate. One of the reasons behind this is the lack of suitable surface-sensitive characterization techniques that can observe charge carrier dynamics in real time, precisely on the surface, as it is well known that Na is mostly localized onto the surface of CIGSe, rather than in the bulk of the film.13,23,24,27,28 The structural changes on the surface post Na incorporation can be verified following several microscopic and spectroscopic characterization techniques; however, its effects on modifying the electronic properties of CIGSe have only been inferred from the device measurements, or predicted from theoretical calculations.12-28 Spectroscopic studies have also been reported to understand the effect of Na on charge carrier dynamics, however, because of the large penetration depth of the laser beam, these studies provide the dynamical information mainly from the bulk of the material rather than the surface.29,30 Herein, we explore and decipher the impact of Na incorporation on CIGSe nanocrystals thin film by four dimensional ultrafast scanning electron microscopy (4D S-UEM), which has the unique capability to visualize charge carrier dynamics selectively on materials surfaces in real time and space with unprecedented spatial and temporal resolution.11,31-39 The concept of the experimental technique is depicted in Scheme 1. ZnS shelled CIGSe nanocrystals are chosen as the active material instead of CIGSe, as it has already been observed that shelling helps to remove a large fraction of trap states and increases the charge carrier lifetime.11,40 Na is incorporated via the use of an additional NaF layer on the sodalime glass substrate prior to deposition of CIGSe-ZnS nanocrystal layer, followed by annealing at 400 oC for 15 mins. It is observed that this post deposition technique increases the grain size, suggesting passivation of grain boundaries, and enhancement of (112) texture of CIGSe. Interestingly, the time resolved secondary electron (SE)


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images obtained from the 4D S-UEM measurement and the extracted SE kinetics, clearly reveal that NaF treatment removes the fast electron trapping channels and increases charge carrier lifetime at the surface of the CIGSe-ZnS film. Finally, DFT calculations suggest that new charge centers (NaCu) can stabilize the CIGSe-ZnS interfacial structure via filling Cu vacancies with diffused Na ions.

Scheme 1. The pump probe concept employed in 4D S-UEM, where the optical pump excites the specimen and the time delayed electron pulse probes it. The collected secondary electrons help to build the time-resolved images. EXPERIMENTAL METHODS Synthesis of nanocrystals: Materials: Copper (I) chloride (CuCl, 99.99%), Indium (III) chloride (InCl3, 98%), Gallium (III) chloride (GaCl3, 99.99%), Zinc (II) acetate (Zn(OAc)2, tech.), 1-Dodecanethiol (DDT, >98%), Octadecylamine (ODA, 97%), Oleylamine (OAm, tech.), Octadecene (ODE, tech.), Sulfur powder


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(S) were purchased from Aldrich. Se powder (200 mesh, 99.99%) was purchased from Alfa-Aesar. All the chemicals were used as received without further purification. Preparation of Cu/In/Ga stock solution: The Cu, In and Ga stock solutions were prepared by dissolving 0.5 mmol of CuCl (0.049g), 0.25 mmol of InCl3 (0.055g), 0.25 mmol of GaCl3 in 2.5 mL OAm each in three separate vials. The vials were degassed by purging them with N2. Preparation of Se stock solution: The Se stock solution was prepared by dissolving 1 mmol of Se powder (0.079g) in 0.5 mL DDT and 0.5 mL OAm. The vial was degassed by purging them with N2. Synthesis of CIGSe nanorcystals: Synthesis of CIGSe nanocrystals followed a reported procedure.11 Briefly, 7 mL ODE was taken in a three-neck flask and degassed by purging with N2 for 15 min. In a separate vial, 0.5 mL of each cation solution was dissolved in 0.5 mL DDT and 0.5 mL OAm, degassed and heated until a clear yellow solution was formed. Then, the flask was heated to 200 °C and 0.2 mL Se stock solution and the cation solution were injected back into the flask. The color of the solution darkened immediately, indicating the formation of NCs. The reaction mixture was annealed for 10 min. Then, the reaction mixture was cooled down and the synthesized NCs were purified using ethanol as well as acetone as non-solvents and chloroform as a dispersing solvent. Shelling of CIGSe nanocrystals with ZnS: For shelling with ZnS also, literature reported protocol was followed.35 In a typical procedure, a stock solution of 0.05 mmol Zn(OAc)2 (0.009g) and 0.05 mmol S powder (0.0016g) was dissolved in 0.5 mL ODE and 0.2 mL TOP. This solution was drop-wise added at 150 °C to the purified NCs dispersed in 0.5 g ODA and 8 mL of ODE. The reaction was then annealed for 10 min. Finally, the reaction mixture was cooled to room


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temperature, and the NCs were purified using ethanol as well as acetone as nonsolvents and chloroform as a dispersing solvent. Fabrication of the thin film: CIGSe-ZnS nanocrystals films were fabricated by depositing the nanocrystals on indium coated tin oxide (ITO) substrates. First, the ITO substrates were cleaned with DI Water, detergent solution, acetone and isopropanol in an ultrasonic bath and then dried followed by an oxygen plasma treatment for 10 min. Then the chloroform solution of the nanocrystals was deposited on the substrate by doctor-blading method followed by annealing at 400 oC for 15 min in presence of N2 atmosphere. Deposition of NaF Layer: NaF layer of 15 nm thickness was deposited by thermal evaporation on ITO substrates before the deposition of the nanocrystals. Density functional theory calculations: Density functional theory (DFT) calculations were performed with a generalized gradient approximation (GGA)/Perdew-Burke-Ernzerhof (PBE) level using the projector-augmented wave (PAW) method, as implemented in the VASP code.41,42 The electronic plane-wave cutoff energy was set to 450 eV. For the bulk CIGSe (ZnS), the k-points grid of 6×6×2 (6×6×6) over the Brillouin zone was used during the cell optimization and electronic properties calculations. The resulting lattice constant of Cu4In2Ga2Se8 was a = b = 5.776 Å and c = 11.512 Å, and of zinc-blend ZnS was a = 5.447 Å. For the calculations of CIGSe/ZnS heterojunctions, the in-plane lattice parameters of CIGSe (112) and ZnS (112) surfaces were obtained from optimized bulk Cu4In2Ga2Se8 and ZnS. The corresponding CIGSe/ZnS heterojunctions were built by directly linking these two surfaces with small lattice mismatch (~ 1 %). All the geometries were optimized until all forces on all atoms were smaller than 0.02 eV/Å. The hybrid functional Heyd-Scuseria-Ernzerhof (HSE) with the exact exchange fraction of 0.3 was


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used to obtain the projected density of states and charge densities for conduction band minimum and valence band maximum. All of the three-dimensional representations of crystal structures and charge densities were obtained using VESTA. RESULTS AND DISCUSSION

CIGSe-ZnS nanocrystals are synthesized following colloidal two-step synthesis protocol as previously reported.11 Typically, CIGSe nanocrystals are synthesized first, purified and then redissolved in amine for further shelling with ZnS (see Experimental Methods). Details of the characterizations of the nanocrystals have been provided in the Supporting Information (Figure S1). It should be noted that previous high resolution TEM (HR-TEM) studies show that the shape (spherical) and size distribution of the CIGSe nanocrystals remain similar after ZnS shelling, indicating formation of a diffused alloyed structure via cation exchange.11 Thin films are fabricated by doctor-blading the nanocrystal solution onto ITO coated glass substrate and annealing at 400 o

C in presence of N2 atmosphere. For the Na treated samples, a NaF layer of 15 nm thickness is

deposited on ITO prior to doctor-blading of the nanocrystals. It should be noted here that the thickness of the NaF layer was chosen based on the optimum structure of the film obtained from control experiments with varied NaF thickness. Figures 1(a-b) and (c-d) show the top view atomic force microscopy (AFM) images and cross sectional scanning electron microscopy (SEM) images of the as deposited and NaF treated films, respectively. The increase in the grain size can clearly be observed from the AFM images (Figure 1(a-b)), whereas the cross-section SEM images (Figure 1(c-d)) show identical thickness of the film in both cases. From the XRD data (Figure 1e), it is observed that the CIGSe-ZnS film deposited on NaF shows increased (112) peak intensity with respect to that of (220)/(224). Preferred (112) texturing of CIGSe film with Na addition has already


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been reported in literature.14,22,29 XPS study of the films (Figure S2) shows the presence of Na for both the samples, suggesting that Na diffuses from NaF layer, as well as from the sodalime glass substrate. However, the NaF treated film shows significant increase in the signal intensity, suggesting presence of higher concentration of Na on the surface of the CIGSe-ZnS film. It may be noted that as all Na compounds have identical or very similar Na1S binding energy values, it is not possible to distinguish the exact oxidation state of Na present on the surface from the XPS spectra.27 No change in the bandgap of the CIGSe-ZnS is observed after NaF treatment (Figure S3). 10.3 nm

(a)

CIGSe-ZnS

1 µm

13.5 nm

(b) CIGSe-ZnS+NaF

1 µm

-10.3 nm

(e)

-13.5 nm

(d)

(c)

500 nm

500 nm

Figure 1. (a-b) AFM images (scale bar corresponds to 1 m) (c-d) Cross-section SEM images of the as deposited and NaF treated CIGSe-ZnS thin film, respectively. Top view SEM images of the films are provided in supporting information (Figure S4) (e) XRD of the same, and lattice planes of CIGSe corresponding to the peaks are marked.


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To study and decipher the effect of NaF treatment on the surface charge carrier dynamics, we compared the 4D S-UEM images of the thin films with and without NaF treatment. Detailed information on 4D S-UEM setup and its mode of operation has been published elsewhere.31,34 Briefly, S-UEM combines the output of a fs Clark-MXR fiber laser system with a modified Quanta FEI-650 SEM. The fundamental output of the laser centered at 1030 nm (pulse ∼270 fs pulse duration and 200 kHz to 25.4 MHz repetition rate) is split by a 1:1 beam splitter to pump two second/third harmonic generators (HG, Clark-MXR) simultaneously to produce 515 and 343 nm pulses. The 343 nm output is focused onto the Schottky field emitter tip inside the SEM to generate the pulsed electrons, which are then accelerated towards the sample using 30 kV voltage. The 515 nm output (pulse energy 0.062 nJ, excitation fluence 10 J/cm2) enters the microscope through a viewport at 50-degree angle relative to the surface normal and delivers the excitation clocking pulse to the sample. The scanning process of the electron beam takes place across the surface of the sample, including both the laser excited and unexcited regions in raster pattern, and the SEs emitted from the sample are detected by a positively biased Everhart Thornley detector. Although the primary pulsed electrons are accelerated at 30 kV, the SEs emitted are quite low in energy (peaking in the range of 3-5 eV), thus originating exclusively from the first few nanometers of the sample surface being probed. The high acceleration voltage (30 kV) also ensures minimal temporal spreading of the pulsed electrons, leading to better temporal resolution during the surface imaging. The SE images are obtained at different time delays between the electron and optical pulses, as an integration of 64 frames with a dwell time of 300 ns at each pixel to improve the signal-to-noise ratio. Finally, all experiments are conducted at a repetition rate of 8 MHz to ensure full recovery of the specimen before the arrival of the next pulse.


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0.5

CIGSe-ZnS CIGSe-ZnS+NaF Fit

SE Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.0 0.2

0.4 Time delay (ns)

0.6

Figure 2. Upper panel: Time-resolved SE images from CIGSe-ZnS thin film with and without NaF treatment at the indicated time delays. The dashed ellipses indicate the footprint of the clocking optical beam on the specimen (∼40 μm). No observable change in the contrast at far negative time delays indicates the full recovery of the system to the initial state after each pump−probe event. Lower panel: Dynamics of the temporal evolution of the SE intensity at the center of the laser footprint region. Figure 2 (upper panel) shows the time-resolved SE images obtained from the CIGSe-ZnS thin film with and without NaF treatment at indicated time delays after 515 nm excitation pulse. No change in the image contrast is observed at the far negative time delays (−577.5 ps), indicating the


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recovery of the system to the initial state after each pump−probe experiment. As can clearly be seen in both samples, a dark contrast evolves in the laser-illuminated region with time. This can be understood as follows. At positive time delays, the optical pulse precedes the electron pulse and promotes a fraction of the valence-band electrons to the conduction band, which are then scattered by the energetic primary electron beam, leading to greater SE emission and hence an overall gain in the SE signal intensity. This would be reflected in the time-resolved difference images as emergence of a “bright” contrast. However, dark contrast observed at positive delay times can be an indicator for the suppression of the SE emission via scattering processes (SE energy loss). In this case, scattering events with electron-hole pairs generated with the optical pulse is a possibility for the energy loss. As the effective cross-section for the scattering of SEs with conduction electrons is much higher than that with valence electrons, a decrease in SE emission and, subsequently, low contrast are observed.11,33,36,37 Though the SE signal remains dark in both the samples, a closer look at the SE images reveals that the dark signal fades away much faster for the sample without NaF treatment compared to the untreated sample. Since fading away of the SE signal indicates carrier (electron/hole) recombination, we plot the SE intensities at the center of the laser excitation footprint (50 × 50 pixels area) as a function of the time delays between optical and electron pulses to get a clear idea about charge carrier recombination at the surface (Figure 2, lower panel). Though shelling of the nanocrystals with higher bandgap material ZnS is expected to remove a significant amount of the carrier trapping channels, a fast component within the first 200 ps time frame can still be observed for the sample without NaF treatment, which indicates presence of further carrier traps in the sample, which are detrimental to the light harvesting efficiency of any material. Interestingly, this fast component is substantially removed (~57%) in NaF treated sample, thus signifying removal


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of surface trap channels from the thin film. We further compared our imaging results with carrier dynamics obtained for the films from ultrafast pump−probe spectroscopy with fs and broadband capabilities. A slower ground state bleach recovery and reduction of the fast component for the NaF treated sample were observed (Figure S5), which in turn, indicates removal of deleterious electron trapping channels. However, when both the charge carrier dynamics obtained via S-UEM and TA are compared, it is obvious that the change after NaF treatment is more prominent from the SE kinetics. The reason for this stems from the large penetration depth of the laser beam in ultrafast pump−probe spectroscopy, which acquires information mainly from the bulk of the sample, whereas the SEs emitted from the first few nanometers of the top surface are detected by 4D S-UEM, allowing any dynamical changes in the surface brought by Na localization on the surface to be more prominently monitored.


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Figure 3. (a-c) Optimized crystal structures, (d-f) projected density of states (PDOS) and electronic charge distributions for valence band maximum (VBM) and conduction band minimum (CBM)

of

ideal

Cu8In4Ga4Se16/Zn16S16,

Cu7In4Ga4Se16/Zn16S16

(Cu

defect),

Cu7NaIn4Ga4Se16/Zn16S16 (NaCu intersite) heterojunctions calculated at HSE level of theory.


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We further performed density functional theory (DFT) calculations to understand the charge carrier dynamics of CIGSe-ZnS before and after NaF treatment (the computational details can be found in Experimental Methods). It has previously been reported that Cu deficiencies near the CuInSe2/CdS interface are important characteristics of efficient CuInSe2 solar cells,43 and the Na dopants could diffuse into CuInSe2 layer close to the interface with CdS buffer layer.44 Based on this, we focus on the effects of Cu vacancies and Na dopants on the electronic properties of CIGSeZnS interfaces. Here we use crystal Cu4In2Ga2Se8 to represent the experimental Cu0.96In0.56Ga0.48Se2 with similar composition. As shown in Figure S6, the calculated bandgap of Cu4In2Ga2Se8 is 1.52 eV using HSE hybrid functional. After incorporating Na into CIGSe by replacing a Cu atom, the formed Cu3NaIn2Ga2Se8 crystal has an increased bandgap of 1.69 eV. We further constructed CIGSe/ZnS heterojunctions by directly combining the CIGSe (112) and ZnS (112) surfaces with very small lattice mismatch,45 and considered three models of i) ideal CIGSe-ZnS heterojunction ii) with Cu defect near interface, and iii) with Na dopants at Cu defect site, as shown in Figures 3a-c. The ideal CIGSe-ZnS heterojunction has bandgap of 1.10 eV, and the charge densities for both valence band maximum (VBM) and conduction band minimum (CBM) are mostly delocalized at CIGSe side. After introducing a Cu defect near CIGSe-ZnS interface, although charge density distributions for VBM and CBM retain the same electronic feature, the induced trap state (red area in PDOS of Figure 3e) could lead to the charge-defect scattering at the interface, further decreasing the charge carrier lifetime. Once these Cu defects are occupied by Na dopants through ion diffusion after NaF treatment, the formed Cu7In4Ga4Se16/Zn16S16 heterojunction has the band gap of 1.39 eV, matching well with the experimental result. Therefore, the Na dopants could stabilize interfacial structures and eliminate


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the trap state induced by Cu vacancies, leading to the longer charge carrier lifetime as compared to untreated one. 10-3

Dark

(b)

(a) Current (A)

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10-3

Light

10-4 10-5

10-5

10-6 10-7

10-7

10-8 10-9

10-9

10-11

With NaF Without NaF

-2

-1

0

10-10

With NaF Without NaF

1

2

-2

-1

0

1

2

Voltage (V)

Figure 4. (a) Schematic of the photodetectors fabricated with CIGSe-ZnS nanocrystals with and without NaF treatment (b) Comparison of dark current and photocurrent (under white light illumination) of the photodetectors. To further elucidate the role of Na, we fabricated planar devices on glass for electrical characterization (Figure 4a). As expected, the electrical conductivity of NaF treated CIGSe-ZnS film was found to be much higher than that of untreated CIGSe (Figure 4b), which can be attributed to the removal of carrier trapping channels present on the nanocrystal surface. As the presence of surface traps is directly correlated with light harvesting efficiency, we compared the performance of the NaF treated and untreated devices under white light illumination (Figure 4b). The NaF treated CIGSe device indeed exhibited better performance than the untreated device, which can be attributed to the better microstructure as well as lower trap density in the NaF treated sample.


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10-11

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CONCLUSION To summarize, lack of suitable surface sensitive techniques to monitor the charge carrier dynamics selectively at materials surfaces has so far obscured the mechanism behind the improvement of CIGSe based photovoltaic devices upon addition of alkali metal. Here, leveraging on the unique capabilities of 4D S-UEM, we visualized the charge carrier dynamics selectively on the CIGSe-ZnS films before and after NaF treatment. Our results revealed that the ultrafast electron trapping channels on the surface of CIGSe-ZnS thin film are significantly reduced by NaF treatment. Removal of such fast carrier trapping channels on the surface can be a critical factor to enhance the light harvesting efficiency, which is supported by DFT calculations, and improved optoelectronic properties of the NaF treated CIGSe device. Thus, this study provides a fundamental understanding of the effect of NaF treatment on the surface charge carrier dynamics of the CIGSeZnS nanocrystals, while paving the way for further improvement of device efficiency of the material in an economically viable way.

ASSOCIATED CONTENT Supporting Information. Instrumentation and supporting figures. This material is available free of charge via the Internet at http://pubs.acs.org AUTHOR INFORMATION Corresponding Author *Email. [email protected]


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ACKNOWLEDGMENT The research reported in this publication was supported by funding from King Abdullah University of Science and Technology (KAUST).

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