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Distribution of boron and phosphorus and roles of co-doping in colloidal silicon nanocrystals
Acta Materialia 178 (2019) 186e193
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Distribution of boron and phosphorus and roles of co-doping in colloidal silicon nanocrystals Keita Nomoto a, *, Hiroshi Sugimoto b, Xiang-Yuan Cui a, Anna V. Ceguerra a, Minoru Fujii b, **, Simon P. Ringer a, *** a The University of Sydney. Australian Centre for Microscopy & Microanalysis and School of Aerospace Mechanical and Mechatronic Engineering, The University of Sydney, New South Wales, 2006, Australia b Department of Electrical and Electronic Engineering, Kobe University, 657-8501, Kobe, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 February 2019 Received in revised form 21 June 2019 Accepted 8 August 2019 Available online 10 August 2019
Boron (B) and phosphorous (P) co-doped colloidal silicon nanocrystals (Si NCs) have unique sizedependent optical properties, which lead to potential applications in optoelectronic and biomedical applications. However, the microstructure of the B and P co-doped colloidal Si NCs e in particular, the exact location of the dopant atoms in real space, has not been studied. A lack of understanding of this underlying question limits our ability to better control sample fabrication, as well as our ability to further develop the optical properties. To study the microstructure, a process enabling atom probe tomography (APT) of colloidal Si NCs was developed. A dispersion of colloidal Si NCs in a SiO2 sol-gel solution and a low temperature curing are demonstrated as the key sample preparation steps. Our APT results demonstrate that a B-rich region exists at the surface of the Si NCs, while P atoms are distributed within the Si NCs. First principles density functional theory calculations of a Si NC embedded in SiO2 matrix reveal that P atoms, which always prefer to reside inside a Si NC, significantly influence the distribution of B atoms. Specifically, P atoms lower the B diffusion barrier at Si/SiO2 interface and stabilize B atoms to reside within individual Si NCs. We propose that the B-modified surface changes the chemical properties of the Si NCs by (i) offering chemical resistance to attack by HF and (ii) enabling dispersibility in solution without aggregation. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Atom probe tomography Silicon nanocrystals Boron phosphorus co-doping Colloid Density functional theory
1. Introduction Colloidal semiconductor nanocrystals (NCs, alternatively know as quantum dots) have great potential for optoelectronics and biomedical applications [1e3]. As colloidal materials, they are amenable to solution-based processes such as spin-coating and inkjet printing, which is desirable for low-cost devices, and is compatible with a flexible substrate [1]. The unique size-dependent optical properties of colloidal NCs are also applicable for in-vivo/invitro fluorescence imaging and cancer treatments [2,3]. To optimise these applications, it is important to understand the fluorophore's absorption/emission band profiles, the emission lifetime, the
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (K. Nomoto), [email protected] (M. Fujii), [email protected] (S.P. Ringer). https://doi.org/10.1016/j.actamat.2019.08.013 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
fluorescence quantum yield, the molar absorption coefficient, and the thermal/photochemical stabilities etc. [4]. Cadmium, lead and indium based NCs exhibit excellent results in terms of the optical properties and stability [4]. However, a major concern of using such elements is their toxicity and the potential risk of heavy metal leakage [4]. Colloidal silicon (Si) NCs have the advantage that they are non-toxic and their optoelectronic properties have been wellstudied [5e9]. Similar to other semiconductor NCs, the physical size of the Si NCs can be controlled, and their diameters typically range from 2 to 10 nm. Accordingly, the luminescent range can be tuned by exploiting the quantum confinement effect [10]. Our group has developed a methodology to fabricate Si NCs dispersible in polar solvents without organic ligands. In the allinorganic Si NCs, boron (B) and phosphorus (P) are heavily doped and the dopants play a crucial role to attain an appropriate solution dispersibility in Si NCs [11e16]. Furthermore, doping reduces the luminescence energy of the Si NCs; B and P co-doped Si NCs have size-controllable luminescence in a 650e1350 nm range, whereas
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undoped Si NCs luminesce in a 650e1100 nm range. This former wavelength range is appropriate for biomedical applications due to the long penetration depth attainable in biological substances [17,18]. Fluorescence imaging of human osteoblasts by B and P codoped Si NCs has been demonstrated [19], and the cytotoxicity to a variety of human cells has been evaluated as low [20]. Furthermore, various types of Si NC-based composites such as metal-Si NCs hybrid nanoparticles, surface functionalized Si NCs with silane coupling agents, and a Si NC layer on a plasmonic substrate have been developed to enhance the optical property responses [21e25]. Despite the unique and attractive chemical and optical properties of co-doped Si NCs, little is known about the atomic-scale structure, especially the 3D information on the distribution of the dopants. Since dopants play a crucial role to determine the chemical and optical properties of the co-doped Si NCs, a full understanding of their distribution is an essential part of the design and development of these materials. Atom probe tomography (APT) is unique technique for the study of the 3D images of different elements with sub-nanometer spatial resolution [26,27]. However, conducting APT experiments on colloidal nanoparticles (NPs) is not straightforward, and there is no published APT work on the distribution of impurity atoms in colloidal semiconductor NCs. The challenge for APT experiments arises because of the fact that these colloidal NPs are in solution. This makes it very difficult to position an individual NP precisely at the apex of an atom probe tip with sufficient mechanical adhesion to endure the forces associated with a high electric field such as is applied during an APT experiment [28]. In this study, we tackle this challenge and perform APT of B and P co-doped colloidal Si NCs. We successfully re-disperse co-doped Si NCs in SiO2 without agglomeration by mixing the colloidal solution with a sol-gel SiO2 solution and by curing it at very low temperature. The APT results demonstrate that the B atoms are favourably located at the surface of the Si NCs, while the P atoms are more likely to be located inside the Si NCs. The results suggest that B-enrichment at the surface of the Si NCs is responsible for certain characteristic chemical properties such as high resistance to HF. Increasingly, first principles density function theory (DFT) calculations and APT are being used as companion simulation and experimental tools [29]. In this work, our DFT calculations revealed binding energies that supported the observed results of the APT experiments. More specifically, these calculations indicate the phenomenology of the co-doped NCs, revealing that the presence of P atoms effectively reduces the energy barrier for B atoms to diffuse across the Si NC/SiO2 matrix interface. 2. Experimental
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solution (Sie05S, Kojundo Chemical Lab.) and 20 mL of concentrated colloidal Si NCs (~5 mg/mL) were mixed and dropped on a Si wafer. The sample was heated at 100 C for 1 min to remove solvents and then annealed at 200 C for 30 min in air. The calculated diffusion length of B and P is 1.1 1017 cm and 5.3 1018 cm, respectively [30]. Therefore, impurity profiles of Si NCs are considered to be preserved, within the resolution of APT, after the solidifying process. Finally, 10 nm cap layer of SiO2 was deposited by sputtering. 2.3. Specimen preparation and instrumental parameters Sharp needle-shaped tips for atom probe experiments were prepared using an Auriga focused ion beam scanning electron microscope (FIB-SEM) (Zeiss). Approximately 500 nm layer of Pt was deposited onto the surface of the samples to protect the film from Ga ions damage during focused ion beam milling. The samples prepared by FIB-SEM lift-out [31] were analyzed in a LEAP™ 4000XSi (CAMECA) with a pulsed UV laser (l ¼ 355 nm) and a detector efficiency of 0.57 [26]. The chamber pressure was maintained within the range of 1012e1011 Torr. The laser pulse energy was set to 100 pJ at a pulse rate of 250 kHz. The ion detection rate was set to 0.5% and the typical starting and ending voltage was around 2500 V and 5500 V, respectively. In selecting these instrumental parameters, we were mindful of a previous APT study where it was proposed that the detection of B atoms reduces at higher laser power (e.g. B detection reduced to ~65% at a laser power of 500 pJ [32]). The specimen was cooled to ~40 K in these experiments. The APT data was then reconstructed using the commercially available IVAS™ software (version 3.6.6). To create Si NC volume, the iso-concentration surface that contains at least 57 at. % Si was used with the same reconstruction parameters such as a voxel size of 1 nm and a delocalization value of 3 nm for both samples of Si NCs in BPSG and colloidal Si NCs, as in the previous report [26,33]. The transmission electron microscope (TEM) observation of Si NCs was performed using a JEM-2100F (JEOL) electron microscope operated at 200 kV. High resolution TEM in Fig. 1c was conducted on Themis Z double-corrected S/TEM (Thermo Fisher Scientific) operated at accelerating voltages of 300 kV. For TEM observations, Si NCs solutions were drop-casted on a TEM mesh. The PL spectra of the colloidal solutions and solid samples were obtained using a single spectrometer equipped with a liquid-N2 cooled InGaAs diode array (OMA-V-SE, Roper Scientific). The PL excitation was performed with a 405 nm laser diode (CUBE 405100C, Coherent). The spectral response of the detection system was corrected with the reference spectrum of a standard halogen lamp. All the measurements were carried out at room temperature.
2.1. Preparation of B and P co-doped Si NCs 2.4. DFT calculation The B and P co-doped colloidal Si NCs were prepared by a cosputtering method as described in our previous papers [14]. Sirich borophosphosilicate glass (BPSG) films were first prepared by simultaneous sputtering of Si, SiO2, B2O3 and P2O5. The film deposited on thin stainless-steel plate was peeled off and crushed using a mortar. The powder was annealed in a N2 gas atmosphere at 1150 C for 30 min to grow Si NCs in BPSG matrices. The Si NCs were liberated from matrices by HF etching (46 wt. %) for 1 h in an ultrasonic bath. Precipitates of Si NCs were removed by centrifugation (4000 rpm, 1 min) and dispersed in methanol. 2.2. Preparation of co-doped Si NCs re-dispersed in SiO2 thin film The co-doped Si NCs dispersed in SiO2 thin film was developed by a new approach shown in Fig. 2. A 10 mL volume of SiO2 sol-gel
DFT calculations were performed using the generalised gradient approximation (GGA) [34] in VASP [35]. The plane wave basis set cutoff energy was 400 eV. The Monkhorst-Pack grids of (1 1 1) were used for the large 648-atom supercells in reciprocal space. The energy convergence criterion between two electronic steps was 104 eV. The convergence criteria for the forces on the atoms are less than 0.01 eV/Å. 3. Results and discussion Fig. 1a is a photograph of the sample in methanol used in this study. The clear yellow-ish solution indicates that the sample is well-dispersed in methanol. As shown in Fig. 1b and c, low and high magnification TEM images of the colloidal Si NCs demonstrate that
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Fig. 1. (a) Photograph of the B and P co-doped colloidal Si NCs in methanol. TEM micrographs of the colloidal Si NCs. (b) at low magnification, and (c) at high magnification.
Fig. 2. Diagram of the specimen preparation procedure developed for atom probe measurement. (a) solid phase nucleation of Si NCs in BPSG. (b) colloidal Si NCs in methanol. (c) mixing of colloidal Si NCs in methanol and SiO2 sol-gel solution. (d) dropping the mixture onto a Si wafer. (e) spin-coating. (f) baking to remove solvents. (g) annealing to solidify the film and so create a stable dispersion of Si NCs in the SiO2 matrix (the photograph is of the solidified film on the Si wafer). (h) preparing the atom probe tip by FIB-SEM (the SEM image was taken immediately after the final annular milling step).
the Si NCs are well-separated each other without agglomeration. The methodology to fabricate the well-dispersed colloidal Si NCs in solution is to be found in our previous work [14]. Fig. 2 illustrates the procedure of the specimen preparation for atom probe measurements. First, B and P co-doped Si NCs were grown in BPSG matrices by the procedure described in the experimental section. In this article, we regard the “Si NCs in BPSG” (Fig. 2a) as a reference to compare the data with those of colloidal Si NCs, which were prepared by etching out BPSG matrices by a HF solution and by transferring the free-standing Si NCs in methanol (Fig. 2b). The Si NCs in the methanol solution were re-dispersed in the SiO2 film (Fig. 2g) for atom probe experiment. Fig. 3a provides the transmittance spectra of the three different Si NCs types (i.e. the Si NCs in the initial BPSG matrix (black curve), the colloidal Si NCs suspended in methanol (red curve), and the Si NCs in the fixed dispersion within the SiO2 matrix (green curve)).
The absorption in the blue-to-UV region is due to the Si NCs. The very high transmittance in other regions demonstrates the formation of a flat and smooth film. Note that the number density of the final dispersion of Si NCs embedded in the SiO2 matrix is much lower than in the original BPSG matrix, and in the methanol suspension. Fig. 3b provides the corresponding photoluminescent (PL) spectra for the different types of the samples. Overall, the PL peak position is observed at around 1100e1150 nm for all samples. The PL peak of B and P co-doped Si NCs typically shifts to longer wavelength compared to undoped Si NCs and this would be due to the donor-to-acceptor transition [36]. The PL spectrum of co-doped colloidal Si NCs suspended in methanol (red curve) is shifted to slightly shorter wavelengths compared to those in the BPSG matrix (black curve). This is consistent with our previous work and is thought to be due to the change in the size and surface termination of Si NCs by etching [13]. The PL spectrum of the Si NCs dispersed in
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Fig. 3. (a) Transmission spectra and (b) Photoluminescent spectra of colloidal Si NCs dispersed in SiO2 (light green) compared to the co-doped Si NCs in BPSG (black) and as suspended in methanol (red). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
the SiO2 (green curve) exhibits a very similar profile to that of the Si NCs in methanol, confirming that there is no significant change in the nature of the size and dopant distribution caused by the sol-gel embedding process used in this study. Fig. 4 provides the results of atom probe tomography from the B and P co-doped Si NCs dispersed in SiO2. The results are provided in the form of tomographic atom maps. For comparison purposes, tomographic atom maps and mass spectra of solid phase co-doped Si NCs in BPSG [33] are provided in S1 and S2 in the supplementary. The 2D tomographic atom maps show that the distribution of Si, P, and B atoms are concentrated almost at the same area. It was difficult to disperse Si NCs uniformly on the Si substrate during the solidifying process. As a result, only 3 out of 25 atom probe tips contained Si NCs and no Si NCs were in the atom probe tip for the rest of the samples. The measured mean diameter of the Si NCs dispersed in SiO2 was 3.3 nm (95% confidence interval, 2.9 nme3.7 nm). The size distribution is provided in S3 in the supplementary. This compares with the average diameter of the Si NCs in the BPSG of 4.2 nm (95% confidence interval, 4.1 nme4.4 nm). Although the measured diameter is probably not absolute value due to an artifact such as local magnification effects (LME) [37], the trend to smaller co-doped Si NCs dispersed in the SiO2 matrix compared to the average size of those in the original BPSG is consistent with our PL measurement results provided in
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Fig. 3. Moreover, a decrease in the average dimeter of the NCs may be expected as a result of the HF etching and is, in fact, consistent with a detailed study on the effect of HF etching resistance of undoped and doped Si NCs [12]. Fig. 5 provides an analysis of our atom probe data using a proxigram analysis [38] to compare the distribution of Si, O, P, and B atoms between the Si NCs dispersed in (i) the SiO2 matrix, and (ii) the BPSG (data included for ease of comparison from our previous work [33]). The proxigrams were averaged for all Si NCs. The same experimental and reconstruction parameters were used for experiments on both sets of samples in the Si NCs/SiO2 system and in the Si NCs/BPSG system. This enables us to compare the atomic scale distribution of elements in the Si NCs before and after HF etching. The results in Fig. 5 demonstrate that the elemental distribution profiles obtained between these samples are generally similar save for one particular difference which will be discussed in detail. Note that there would be LME [37] due to the differences in the field evaporation behavior of the different matrix materials (Si NCs and SiO2 matrix in this case) [39]. This typically occurs at precipitate interfaces over sub-nm distances [40]. The element concentration in this study is not to be taken literally since there are uncertainties of the atom probe measurement such as the noise level of the analysis, the effect of thermal tails in the mass spectra [41], and the absence of LME correction [42,43]. Although it is difficult to clearly define the boundary between Si NCs and SiO2 matrix due to the LME, there are several methods to generate Si NCs such as using Si iso-concentration and cluster analysis [44,45]. As mentioned in the experimental section, the proxigram profile in Fig. 5 is made using Si iso-surface concentration of 57 at.% as shown in our previous study [33]. With this value, the Si NCs are generated in sections where the Si atoms are densely gathered as shown in S4 in the supplementary. In S5, we have also conducted the proxigram analysis using the cluster algorithm (see the supplementary). This proxigram profiles are very similar to the one in Fig. 5, even though the method to generate Si NCs is different. Therefore, the discussion below should be compatible regardless of the ways to define the boundary between Si NC and SiO2 matrix. Firstly, comparing the profiles in Fig. 5a, we note very similar Si and O concentration profiles across the NCs embedded in both the SiO2 matrix and the BPSG matrix. These proxigram traces within the NCs reflect the decrease in their size in the case of the SiO2 matrix. Secondly, the distribution of the P atoms in Fig. 5b reveals that the distribution of P atoms for both samples is almost identical. This observation holds notwithstanding the significant error bars for the concentrations within the NCs due to the low number of particles analyzed, and the reduced statistics (atom counts) as we move towards the center of the Si NCs. For example, the P concentration rapidly increases toward the interface of the Si NCs and remains at around 4 at. % in the core of the Si NCs. This P distribution profile which shows the incorporation of P atoms into Si NCs is consistent with several other APT studies reported elsewhere [46e48], whether or not the P atoms are solely-doped or co-doped with B. Finally, the B concentration profile is provided in Fig. 5c. Both the SiO2 matrix and the BPSG matrix samples exhibit an accumulation of B at the surface of the Si NCs, and the concentration gradually decreases toward the core of the NC. The concentration gradients for B are considerably less than those for P. The proxigrams were also calculated for a single and/or a few Si NCs and we found that the B-enrichment is clearer for larger NCs compared to smaller NCs (S6 in the supplementary). Interestingly, our previous work [33] revealed that such B-enrichment at the surface of the Si NCs embedded in the BPSG matrix was not observed for the case of solely-B-doped Si NCs. In fact, the B concentration in Si NCs of solely-B-doped Si NCs was less than that of the B and P co-doped Si NCs. Therefore, it is proposed that co-doping is an effective
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Fig. 4. Atom probe tomography results. (a) Tomographic atom maps of the B and P co-doped colloidal Si NCs dispersed in the SiO2 matrix. Companion projections are provided from thin sliced regions (the boxes drawn in grey dots) below each tomogram. (b) A high resolution atom map of a single B and P co-doped colloidal Si NC.
Fig. 5. Proxigram analysis on B and P co-doped Si NCs in BPSG and Si NCs dispersed in SiO2. (a) silicon and oxygen profiles, (b) phosphorus profiles, and (c) boron profiles in the Si NCs, in the matrix and at the interface. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
means of promoting segregation of both B and P species to the interface of Si NCs, and to achieve partitioning of both species inside (within) the NCs. To shed light on the underlying mechanism of B and P distribution, we have performed extensive first principles DFT calculations on Si NCs in a SiO2 matrix. The goal of our DFT calculations was to explore the preferential location of B and P atoms during growth of the Si NCs in a BPSG matrix. As with Ref. [49], we simulated a Si NC embedded in a large 648-atom supercell betacristobalite-SiO2 matrix by removing 80 central O atoms with inverse symmetry. The size of Si NC in this model is approximately 1.4 nm in diameter. Previous DFT studies suggest that the simulated elemental distribution and electronic structures are well qualitatively applied to larger Si NCs, as fabricated in experiments [50e52]. We then performed full relaxation (both volume and internal atomic positions). This resulted in a contraction of the lattice constant from 21.48 Å to 21.21 Å, accompanied by a large energy gain of 1.54 eV. The relaxed atomic structure prior to doping is shown in Fig. 6a. Note in Ref. [49], this relaxation step and the associated volume contraction was neglected. While this volume contraction has only a minor effect on the SieO bond length far from the interface in the matrix
(changing from 1.618 Å to 1.615 Å), it has a major influence on the atomic structure of the Si NC at the interface region. The nanoscale structure of Si NC embedded in the SiO2 matrix can be divided into three regions: (I) the SiO2 matrix, (II) the interface region, and (III) the Si NC. In each region, the Si atoms have distinctly different chemical coordination environments, as shown in Fig. 6b. To energetically map the dopant distribution profile, one substitutional B or P atom was placed at one of the eight doping sites so as to explore each of the three different regions. The relative energy curves as a function of distance to the center of the Si NC are provided in Fig. 7. It is clearly evident that P atoms prefer to reside inside of the Si NC, due to the large energy gain (~4 eV/atom) and negligible energy barrier at the interface region. In sharp contrast, the B atoms have two energetic minima: one at the interface region, and the other within the Si NC, with the former slightly favorable by 0.18 eV/atom. Significantly, there exists a large energy barrier (~2.5 eV/atom) for B atoms to diffuse from the SiO2 matrix into Si NC. This suggests that B dopants will prefer to reside at the NC/ matrix interface region. The similar trend can be observed from previous DFT calculations, even though the authors used a different structure model [50,53]. Having revealed the behavior associated
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Fig. 6. (a) DFT fully relaxed atomic structure of supercell for Si NC embedded in SiO2 matrix. (b) the O atoms (in red) are omitted for clarity. Various doping sites are labelled with distances to the center of Si NC. Blue spheres denote Si bonded with four O atoms, yellow spheres are Si atoms bonded with three O atoms and one Si atom, green sphere are Si atoms bonded with four Si atoms. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 7. DFT calculated relative energy profiles for different dopants and O vacancy as a function of various doping sites. The distance is from the doping site to the center of the NC-Si.
with single doping, we sought to understand the effects of codoping. Therefore, we first fixed one P atom at one of the favorable sites (close to the C0 and C1 sites in Fig. 6b), and then introduced one B atom at a variety of different doping sites. The energy profile, called “B-after-P”, is provided in Fig. 7. By comparing this with the single B doping energy profile, we can conclude that there are two important effects for B þ P co-doping. Firstly, from an energetic viewpoint, the presence of P atoms significantly enhances the possibility for the Si NC to accommodate B atoms at the interface region, by no less than ~3 eV/atom. The similar preferential B distribution with the assistance of P atoms has been also demonstrated by Guerra et al. and Ni et al. [49,50]. Secondly, from a kinetic perspective, the introduction of P atoms dramatically reduces the energy barriers for B diffusion from 2.5 eV/atom to ~0.7 eV/atom, thus facilitating B species migration into the Si NC. Dynamically, we anticipate that O vacancies will also play an important role in assisting the dopant diffusion, especially at elevated temperature. Thus, we have also evaluated the energy profile for an O vacancy in both the SiO2 matrix and the interface region in Fig. 7. We find that, energetically, the O vacancy exhibits a slight preference to be at the interface over the SiO2 matrix, by ~0.2 eV/vacancy. Nevertheless, we believe that, kinetically, the O vacancy plays an important role in affecting the
doping and the elemental distribution of the dopants in Si NCs as the presence of vacancy will significantly facilitate the diffusion process [54]. We now return to further discuss the APT results. It is worth mentioning that the B concentration at the surface of the Si NCs dispersed in the SiO2 matrix was slightly higher than that of the Si NCs in BPSG (Fig. 5c). This result implies that more B atoms remain inside or near the surface of the colloidal Si NCs, which is in a good agreement with the previous Raman and X-ray photoelectron spectroscopy (XPS) results. Therefore, we propose that the HF etching process preferentially takes place at SieO, SieSi, SieH, SieP bonds rather than at the SieB bond at the surface of the Si NCs. Similar behavior has been demonstrated experimentally for the case of HF etching of porous Si [55]. As shown in our previous paper [12], the resistance to HF solution for doped Si NCs is much stronger than undoped Si NCs, showing that the etching time becomes longer in the order of the undoped Si NCs, the solely-P-doped Si NCs, the solely-B-doped Si NCs, and the B and P co-doped Si NCs. The B and P co-doped Si NCs exhibits highest resistance to the HF solution and we suggest that this is an important indicator as to why the B and P co-doping is essential to fabricate colloidal Si NCs with strong photonic performance. Finally, we discuss the mechanism of dispersibility of our colloidal Si NCs in the methanol, by summarizing the various microstructures observed in the literature. From previous experiments, it has been shown that the z-potential of our co-doped colloidal Si NCs is in the range 30 mV < z < 40 mV [13], indicating that our colloidal Si NC surface is likely negatively charged, and that B and P co-doped Si NCs are stabilized in solution by electrostatically repulsion of each other. On the other hand, undoped, solely-B-doped, and solely-P-doped Si NCs did not form colloids, as demonstrated in a previous study [12]. We suggest that there are two important reasons for this: (i) that the undoped and solely-doped Si NCs are dissolved by HF etching faster than the codoped Si NCs [12,14], and (ii) even if the Si NCs are not fully dissolved during the short time exposure to the HF (e.g. < 10 min), they will aggregate in solution, resulting in their being filtered out by the centrifugation process. By combining our previous APT results [33] with the Raman and XPS results [15], and the DFT calculations here, it is proposed that B atoms are preferentially located around the interface of Si NCs in BPSG, which works more efficiently with the help of P atoms. As shown in Fig. 5, after the HF etching process to fabricate our colloidal Si NCs, the B-enrichment at the surface of the Si NCs was confirmed using APT, indicating that the B-enrichment at the surface of the Si NCs is chemically resistant to HF etching. Considering that the solely-B-doped Si NCs did not effectively form a colloid, B þ P co-doping is needed to enhance the
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B-enrichment at the surface of the Si NCs so as to efficiently protect them from HF etching. In addition to providing HF resistance, B atoms at the Si NC surface are important for their dispersibility in solution. Wheeler et al. has demonstrated that a molecular interaction at the Si NC surface stabilizes Si NCs in solution in the form of colloid [56,57], in an action whereby B atoms behave as a Lewis acid that is attracted to a Lewis base, resulting in dissolution in different kinds of solutions [56,57]. This study also suggests that the existence of B at the Si NCs surface induces a negative surface charge that interacts with a Lewis base such as OH, which enables the Si NCs to be dissolved in methanol or water. 4. Conclusions In conclusion, the atomic-scale microstructure of B and P codoped colloidal Si NCs has been studied using APT experiments and DFT calculations. APT of the colloidal Si NCs was firstly demonstrated and the B and P distribution revealed. Experimentally, it was found that there is a B-rich region at the surface of Si NCs, while the P atoms are more likely to be located inside the Si NCs. Our DFT simulations of Si NCs in a SiO2 matrix indicate that, energetically, P atoms prefer to reside within the Si NCs, and that the presence of P lowers the B diffusion barrier, which then attracts and stabilizes the occurrence of B atoms at the surface of the Si NCs. We suggest that the observed enrichment of B atoms at the surface of Si NCs changes their chemical properties. Firstly, this B-modified surface provides the observed resistance to HF etching. Secondly, the B-modified surface provides a mechanism for the dispersibility of colloidal Si NCs in solutions, enabling the Si NCs to disperse in solution without aggregation. Acknowledgements The authors acknowledge the facilities and the scientific and technical assistance of the Microscopy Australia node at the University of Sydney (Sydney Microscopy & Microanalysis). This research was partially supported by the Faculty of Engineering & Information Technologies at the University of Sydney, under the Research Cluster Program. This research was also undertaken with the assistance of resources from the National Computational Infrastructure (NCI). The NCI and the AMMRF are supported by the Australian Government under the NCRIS program. This collaborative research was also partly supported by the Australian Research Council (grant DP160101713), and the 2018 JSPS Bilateral Joint Research Projects (Japan-Australia), JSPS KAKENHI Grant Number 16H03828 and 18K14092, and the 2015 JST Visegrad Group (V4) Japan Joint Research Project on Advanced Materials.
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Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.actamat.2019.08.013.
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Distribution of boron and phosphorus and roles of co-doping in colloidal silicon nanocrystals Keita Nomoto a, *, Hiroshi Sugimoto b, Xiang-Yuan Cui a, Anna V. Ceguerra a, Minoru Fujii b, **, Simon P. Ringer a, *** a The University of Sydney. Australian Centre for Microscopy & Microanalysis and School of Aerospace Mechanical and Mechatronic Engineering, The University of Sydney, New South Wales, 2006, Australia b Department of Electrical and Electronic Engineering, Kobe University, 657-8501, Kobe, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 February 2019 Received in revised form 21 June 2019 Accepted 8 August 2019 Available online 10 August 2019
Boron (B) and phosphorous (P) co-doped colloidal silicon nanocrystals (Si NCs) have unique sizedependent optical properties, which lead to potential applications in optoelectronic and biomedical applications. However, the microstructure of the B and P co-doped colloidal Si NCs e in particular, the exact location of the dopant atoms in real space, has not been studied. A lack of understanding of this underlying question limits our ability to better control sample fabrication, as well as our ability to further develop the optical properties. To study the microstructure, a process enabling atom probe tomography (APT) of colloidal Si NCs was developed. A dispersion of colloidal Si NCs in a SiO2 sol-gel solution and a low temperature curing are demonstrated as the key sample preparation steps. Our APT results demonstrate that a B-rich region exists at the surface of the Si NCs, while P atoms are distributed within the Si NCs. First principles density functional theory calculations of a Si NC embedded in SiO2 matrix reveal that P atoms, which always prefer to reside inside a Si NC, significantly influence the distribution of B atoms. Specifically, P atoms lower the B diffusion barrier at Si/SiO2 interface and stabilize B atoms to reside within individual Si NCs. We propose that the B-modified surface changes the chemical properties of the Si NCs by (i) offering chemical resistance to attack by HF and (ii) enabling dispersibility in solution without aggregation. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Atom probe tomography Silicon nanocrystals Boron phosphorus co-doping Colloid Density functional theory
1. Introduction Colloidal semiconductor nanocrystals (NCs, alternatively know as quantum dots) have great potential for optoelectronics and biomedical applications [1e3]. As colloidal materials, they are amenable to solution-based processes such as spin-coating and inkjet printing, which is desirable for low-cost devices, and is compatible with a flexible substrate [1]. The unique size-dependent optical properties of colloidal NCs are also applicable for in-vivo/invitro fluorescence imaging and cancer treatments [2,3]. To optimise these applications, it is important to understand the fluorophore's absorption/emission band profiles, the emission lifetime, the
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (K. Nomoto), [email protected] (M. Fujii), [email protected] (S.P. Ringer). https://doi.org/10.1016/j.actamat.2019.08.013 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
fluorescence quantum yield, the molar absorption coefficient, and the thermal/photochemical stabilities etc. [4]. Cadmium, lead and indium based NCs exhibit excellent results in terms of the optical properties and stability [4]. However, a major concern of using such elements is their toxicity and the potential risk of heavy metal leakage [4]. Colloidal silicon (Si) NCs have the advantage that they are non-toxic and their optoelectronic properties have been wellstudied [5e9]. Similar to other semiconductor NCs, the physical size of the Si NCs can be controlled, and their diameters typically range from 2 to 10 nm. Accordingly, the luminescent range can be tuned by exploiting the quantum confinement effect [10]. Our group has developed a methodology to fabricate Si NCs dispersible in polar solvents without organic ligands. In the allinorganic Si NCs, boron (B) and phosphorus (P) are heavily doped and the dopants play a crucial role to attain an appropriate solution dispersibility in Si NCs [11e16]. Furthermore, doping reduces the luminescence energy of the Si NCs; B and P co-doped Si NCs have size-controllable luminescence in a 650e1350 nm range, whereas
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undoped Si NCs luminesce in a 650e1100 nm range. This former wavelength range is appropriate for biomedical applications due to the long penetration depth attainable in biological substances [17,18]. Fluorescence imaging of human osteoblasts by B and P codoped Si NCs has been demonstrated [19], and the cytotoxicity to a variety of human cells has been evaluated as low [20]. Furthermore, various types of Si NC-based composites such as metal-Si NCs hybrid nanoparticles, surface functionalized Si NCs with silane coupling agents, and a Si NC layer on a plasmonic substrate have been developed to enhance the optical property responses [21e25]. Despite the unique and attractive chemical and optical properties of co-doped Si NCs, little is known about the atomic-scale structure, especially the 3D information on the distribution of the dopants. Since dopants play a crucial role to determine the chemical and optical properties of the co-doped Si NCs, a full understanding of their distribution is an essential part of the design and development of these materials. Atom probe tomography (APT) is unique technique for the study of the 3D images of different elements with sub-nanometer spatial resolution [26,27]. However, conducting APT experiments on colloidal nanoparticles (NPs) is not straightforward, and there is no published APT work on the distribution of impurity atoms in colloidal semiconductor NCs. The challenge for APT experiments arises because of the fact that these colloidal NPs are in solution. This makes it very difficult to position an individual NP precisely at the apex of an atom probe tip with sufficient mechanical adhesion to endure the forces associated with a high electric field such as is applied during an APT experiment [28]. In this study, we tackle this challenge and perform APT of B and P co-doped colloidal Si NCs. We successfully re-disperse co-doped Si NCs in SiO2 without agglomeration by mixing the colloidal solution with a sol-gel SiO2 solution and by curing it at very low temperature. The APT results demonstrate that the B atoms are favourably located at the surface of the Si NCs, while the P atoms are more likely to be located inside the Si NCs. The results suggest that B-enrichment at the surface of the Si NCs is responsible for certain characteristic chemical properties such as high resistance to HF. Increasingly, first principles density function theory (DFT) calculations and APT are being used as companion simulation and experimental tools [29]. In this work, our DFT calculations revealed binding energies that supported the observed results of the APT experiments. More specifically, these calculations indicate the phenomenology of the co-doped NCs, revealing that the presence of P atoms effectively reduces the energy barrier for B atoms to diffuse across the Si NC/SiO2 matrix interface. 2. Experimental
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solution (Sie05S, Kojundo Chemical Lab.) and 20 mL of concentrated colloidal Si NCs (~5 mg/mL) were mixed and dropped on a Si wafer. The sample was heated at 100 C for 1 min to remove solvents and then annealed at 200 C for 30 min in air. The calculated diffusion length of B and P is 1.1 1017 cm and 5.3 1018 cm, respectively [30]. Therefore, impurity profiles of Si NCs are considered to be preserved, within the resolution of APT, after the solidifying process. Finally, 10 nm cap layer of SiO2 was deposited by sputtering. 2.3. Specimen preparation and instrumental parameters Sharp needle-shaped tips for atom probe experiments were prepared using an Auriga focused ion beam scanning electron microscope (FIB-SEM) (Zeiss). Approximately 500 nm layer of Pt was deposited onto the surface of the samples to protect the film from Ga ions damage during focused ion beam milling. The samples prepared by FIB-SEM lift-out [31] were analyzed in a LEAP™ 4000XSi (CAMECA) with a pulsed UV laser (l ¼ 355 nm) and a detector efficiency of 0.57 [26]. The chamber pressure was maintained within the range of 1012e1011 Torr. The laser pulse energy was set to 100 pJ at a pulse rate of 250 kHz. The ion detection rate was set to 0.5% and the typical starting and ending voltage was around 2500 V and 5500 V, respectively. In selecting these instrumental parameters, we were mindful of a previous APT study where it was proposed that the detection of B atoms reduces at higher laser power (e.g. B detection reduced to ~65% at a laser power of 500 pJ [32]). The specimen was cooled to ~40 K in these experiments. The APT data was then reconstructed using the commercially available IVAS™ software (version 3.6.6). To create Si NC volume, the iso-concentration surface that contains at least 57 at. % Si was used with the same reconstruction parameters such as a voxel size of 1 nm and a delocalization value of 3 nm for both samples of Si NCs in BPSG and colloidal Si NCs, as in the previous report [26,33]. The transmission electron microscope (TEM) observation of Si NCs was performed using a JEM-2100F (JEOL) electron microscope operated at 200 kV. High resolution TEM in Fig. 1c was conducted on Themis Z double-corrected S/TEM (Thermo Fisher Scientific) operated at accelerating voltages of 300 kV. For TEM observations, Si NCs solutions were drop-casted on a TEM mesh. The PL spectra of the colloidal solutions and solid samples were obtained using a single spectrometer equipped with a liquid-N2 cooled InGaAs diode array (OMA-V-SE, Roper Scientific). The PL excitation was performed with a 405 nm laser diode (CUBE 405100C, Coherent). The spectral response of the detection system was corrected with the reference spectrum of a standard halogen lamp. All the measurements were carried out at room temperature.
2.1. Preparation of B and P co-doped Si NCs 2.4. DFT calculation The B and P co-doped colloidal Si NCs were prepared by a cosputtering method as described in our previous papers [14]. Sirich borophosphosilicate glass (BPSG) films were first prepared by simultaneous sputtering of Si, SiO2, B2O3 and P2O5. The film deposited on thin stainless-steel plate was peeled off and crushed using a mortar. The powder was annealed in a N2 gas atmosphere at 1150 C for 30 min to grow Si NCs in BPSG matrices. The Si NCs were liberated from matrices by HF etching (46 wt. %) for 1 h in an ultrasonic bath. Precipitates of Si NCs were removed by centrifugation (4000 rpm, 1 min) and dispersed in methanol. 2.2. Preparation of co-doped Si NCs re-dispersed in SiO2 thin film The co-doped Si NCs dispersed in SiO2 thin film was developed by a new approach shown in Fig. 2. A 10 mL volume of SiO2 sol-gel
DFT calculations were performed using the generalised gradient approximation (GGA) [34] in VASP [35]. The plane wave basis set cutoff energy was 400 eV. The Monkhorst-Pack grids of (1 1 1) were used for the large 648-atom supercells in reciprocal space. The energy convergence criterion between two electronic steps was 104 eV. The convergence criteria for the forces on the atoms are less than 0.01 eV/Å. 3. Results and discussion Fig. 1a is a photograph of the sample in methanol used in this study. The clear yellow-ish solution indicates that the sample is well-dispersed in methanol. As shown in Fig. 1b and c, low and high magnification TEM images of the colloidal Si NCs demonstrate that
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Fig. 1. (a) Photograph of the B and P co-doped colloidal Si NCs in methanol. TEM micrographs of the colloidal Si NCs. (b) at low magnification, and (c) at high magnification.
Fig. 2. Diagram of the specimen preparation procedure developed for atom probe measurement. (a) solid phase nucleation of Si NCs in BPSG. (b) colloidal Si NCs in methanol. (c) mixing of colloidal Si NCs in methanol and SiO2 sol-gel solution. (d) dropping the mixture onto a Si wafer. (e) spin-coating. (f) baking to remove solvents. (g) annealing to solidify the film and so create a stable dispersion of Si NCs in the SiO2 matrix (the photograph is of the solidified film on the Si wafer). (h) preparing the atom probe tip by FIB-SEM (the SEM image was taken immediately after the final annular milling step).
the Si NCs are well-separated each other without agglomeration. The methodology to fabricate the well-dispersed colloidal Si NCs in solution is to be found in our previous work [14]. Fig. 2 illustrates the procedure of the specimen preparation for atom probe measurements. First, B and P co-doped Si NCs were grown in BPSG matrices by the procedure described in the experimental section. In this article, we regard the “Si NCs in BPSG” (Fig. 2a) as a reference to compare the data with those of colloidal Si NCs, which were prepared by etching out BPSG matrices by a HF solution and by transferring the free-standing Si NCs in methanol (Fig. 2b). The Si NCs in the methanol solution were re-dispersed in the SiO2 film (Fig. 2g) for atom probe experiment. Fig. 3a provides the transmittance spectra of the three different Si NCs types (i.e. the Si NCs in the initial BPSG matrix (black curve), the colloidal Si NCs suspended in methanol (red curve), and the Si NCs in the fixed dispersion within the SiO2 matrix (green curve)).
The absorption in the blue-to-UV region is due to the Si NCs. The very high transmittance in other regions demonstrates the formation of a flat and smooth film. Note that the number density of the final dispersion of Si NCs embedded in the SiO2 matrix is much lower than in the original BPSG matrix, and in the methanol suspension. Fig. 3b provides the corresponding photoluminescent (PL) spectra for the different types of the samples. Overall, the PL peak position is observed at around 1100e1150 nm for all samples. The PL peak of B and P co-doped Si NCs typically shifts to longer wavelength compared to undoped Si NCs and this would be due to the donor-to-acceptor transition [36]. The PL spectrum of co-doped colloidal Si NCs suspended in methanol (red curve) is shifted to slightly shorter wavelengths compared to those in the BPSG matrix (black curve). This is consistent with our previous work and is thought to be due to the change in the size and surface termination of Si NCs by etching [13]. The PL spectrum of the Si NCs dispersed in
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Fig. 3. (a) Transmission spectra and (b) Photoluminescent spectra of colloidal Si NCs dispersed in SiO2 (light green) compared to the co-doped Si NCs in BPSG (black) and as suspended in methanol (red). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
the SiO2 (green curve) exhibits a very similar profile to that of the Si NCs in methanol, confirming that there is no significant change in the nature of the size and dopant distribution caused by the sol-gel embedding process used in this study. Fig. 4 provides the results of atom probe tomography from the B and P co-doped Si NCs dispersed in SiO2. The results are provided in the form of tomographic atom maps. For comparison purposes, tomographic atom maps and mass spectra of solid phase co-doped Si NCs in BPSG [33] are provided in S1 and S2 in the supplementary. The 2D tomographic atom maps show that the distribution of Si, P, and B atoms are concentrated almost at the same area. It was difficult to disperse Si NCs uniformly on the Si substrate during the solidifying process. As a result, only 3 out of 25 atom probe tips contained Si NCs and no Si NCs were in the atom probe tip for the rest of the samples. The measured mean diameter of the Si NCs dispersed in SiO2 was 3.3 nm (95% confidence interval, 2.9 nme3.7 nm). The size distribution is provided in S3 in the supplementary. This compares with the average diameter of the Si NCs in the BPSG of 4.2 nm (95% confidence interval, 4.1 nme4.4 nm). Although the measured diameter is probably not absolute value due to an artifact such as local magnification effects (LME) [37], the trend to smaller co-doped Si NCs dispersed in the SiO2 matrix compared to the average size of those in the original BPSG is consistent with our PL measurement results provided in
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Fig. 3. Moreover, a decrease in the average dimeter of the NCs may be expected as a result of the HF etching and is, in fact, consistent with a detailed study on the effect of HF etching resistance of undoped and doped Si NCs [12]. Fig. 5 provides an analysis of our atom probe data using a proxigram analysis [38] to compare the distribution of Si, O, P, and B atoms between the Si NCs dispersed in (i) the SiO2 matrix, and (ii) the BPSG (data included for ease of comparison from our previous work [33]). The proxigrams were averaged for all Si NCs. The same experimental and reconstruction parameters were used for experiments on both sets of samples in the Si NCs/SiO2 system and in the Si NCs/BPSG system. This enables us to compare the atomic scale distribution of elements in the Si NCs before and after HF etching. The results in Fig. 5 demonstrate that the elemental distribution profiles obtained between these samples are generally similar save for one particular difference which will be discussed in detail. Note that there would be LME [37] due to the differences in the field evaporation behavior of the different matrix materials (Si NCs and SiO2 matrix in this case) [39]. This typically occurs at precipitate interfaces over sub-nm distances [40]. The element concentration in this study is not to be taken literally since there are uncertainties of the atom probe measurement such as the noise level of the analysis, the effect of thermal tails in the mass spectra [41], and the absence of LME correction [42,43]. Although it is difficult to clearly define the boundary between Si NCs and SiO2 matrix due to the LME, there are several methods to generate Si NCs such as using Si iso-concentration and cluster analysis [44,45]. As mentioned in the experimental section, the proxigram profile in Fig. 5 is made using Si iso-surface concentration of 57 at.% as shown in our previous study [33]. With this value, the Si NCs are generated in sections where the Si atoms are densely gathered as shown in S4 in the supplementary. In S5, we have also conducted the proxigram analysis using the cluster algorithm (see the supplementary). This proxigram profiles are very similar to the one in Fig. 5, even though the method to generate Si NCs is different. Therefore, the discussion below should be compatible regardless of the ways to define the boundary between Si NC and SiO2 matrix. Firstly, comparing the profiles in Fig. 5a, we note very similar Si and O concentration profiles across the NCs embedded in both the SiO2 matrix and the BPSG matrix. These proxigram traces within the NCs reflect the decrease in their size in the case of the SiO2 matrix. Secondly, the distribution of the P atoms in Fig. 5b reveals that the distribution of P atoms for both samples is almost identical. This observation holds notwithstanding the significant error bars for the concentrations within the NCs due to the low number of particles analyzed, and the reduced statistics (atom counts) as we move towards the center of the Si NCs. For example, the P concentration rapidly increases toward the interface of the Si NCs and remains at around 4 at. % in the core of the Si NCs. This P distribution profile which shows the incorporation of P atoms into Si NCs is consistent with several other APT studies reported elsewhere [46e48], whether or not the P atoms are solely-doped or co-doped with B. Finally, the B concentration profile is provided in Fig. 5c. Both the SiO2 matrix and the BPSG matrix samples exhibit an accumulation of B at the surface of the Si NCs, and the concentration gradually decreases toward the core of the NC. The concentration gradients for B are considerably less than those for P. The proxigrams were also calculated for a single and/or a few Si NCs and we found that the B-enrichment is clearer for larger NCs compared to smaller NCs (S6 in the supplementary). Interestingly, our previous work [33] revealed that such B-enrichment at the surface of the Si NCs embedded in the BPSG matrix was not observed for the case of solely-B-doped Si NCs. In fact, the B concentration in Si NCs of solely-B-doped Si NCs was less than that of the B and P co-doped Si NCs. Therefore, it is proposed that co-doping is an effective
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Fig. 4. Atom probe tomography results. (a) Tomographic atom maps of the B and P co-doped colloidal Si NCs dispersed in the SiO2 matrix. Companion projections are provided from thin sliced regions (the boxes drawn in grey dots) below each tomogram. (b) A high resolution atom map of a single B and P co-doped colloidal Si NC.
Fig. 5. Proxigram analysis on B and P co-doped Si NCs in BPSG and Si NCs dispersed in SiO2. (a) silicon and oxygen profiles, (b) phosphorus profiles, and (c) boron profiles in the Si NCs, in the matrix and at the interface. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
means of promoting segregation of both B and P species to the interface of Si NCs, and to achieve partitioning of both species inside (within) the NCs. To shed light on the underlying mechanism of B and P distribution, we have performed extensive first principles DFT calculations on Si NCs in a SiO2 matrix. The goal of our DFT calculations was to explore the preferential location of B and P atoms during growth of the Si NCs in a BPSG matrix. As with Ref. [49], we simulated a Si NC embedded in a large 648-atom supercell betacristobalite-SiO2 matrix by removing 80 central O atoms with inverse symmetry. The size of Si NC in this model is approximately 1.4 nm in diameter. Previous DFT studies suggest that the simulated elemental distribution and electronic structures are well qualitatively applied to larger Si NCs, as fabricated in experiments [50e52]. We then performed full relaxation (both volume and internal atomic positions). This resulted in a contraction of the lattice constant from 21.48 Å to 21.21 Å, accompanied by a large energy gain of 1.54 eV. The relaxed atomic structure prior to doping is shown in Fig. 6a. Note in Ref. [49], this relaxation step and the associated volume contraction was neglected. While this volume contraction has only a minor effect on the SieO bond length far from the interface in the matrix
(changing from 1.618 Å to 1.615 Å), it has a major influence on the atomic structure of the Si NC at the interface region. The nanoscale structure of Si NC embedded in the SiO2 matrix can be divided into three regions: (I) the SiO2 matrix, (II) the interface region, and (III) the Si NC. In each region, the Si atoms have distinctly different chemical coordination environments, as shown in Fig. 6b. To energetically map the dopant distribution profile, one substitutional B or P atom was placed at one of the eight doping sites so as to explore each of the three different regions. The relative energy curves as a function of distance to the center of the Si NC are provided in Fig. 7. It is clearly evident that P atoms prefer to reside inside of the Si NC, due to the large energy gain (~4 eV/atom) and negligible energy barrier at the interface region. In sharp contrast, the B atoms have two energetic minima: one at the interface region, and the other within the Si NC, with the former slightly favorable by 0.18 eV/atom. Significantly, there exists a large energy barrier (~2.5 eV/atom) for B atoms to diffuse from the SiO2 matrix into Si NC. This suggests that B dopants will prefer to reside at the NC/ matrix interface region. The similar trend can be observed from previous DFT calculations, even though the authors used a different structure model [50,53]. Having revealed the behavior associated
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Fig. 6. (a) DFT fully relaxed atomic structure of supercell for Si NC embedded in SiO2 matrix. (b) the O atoms (in red) are omitted for clarity. Various doping sites are labelled with distances to the center of Si NC. Blue spheres denote Si bonded with four O atoms, yellow spheres are Si atoms bonded with three O atoms and one Si atom, green sphere are Si atoms bonded with four Si atoms. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 7. DFT calculated relative energy profiles for different dopants and O vacancy as a function of various doping sites. The distance is from the doping site to the center of the NC-Si.
with single doping, we sought to understand the effects of codoping. Therefore, we first fixed one P atom at one of the favorable sites (close to the C0 and C1 sites in Fig. 6b), and then introduced one B atom at a variety of different doping sites. The energy profile, called “B-after-P”, is provided in Fig. 7. By comparing this with the single B doping energy profile, we can conclude that there are two important effects for B þ P co-doping. Firstly, from an energetic viewpoint, the presence of P atoms significantly enhances the possibility for the Si NC to accommodate B atoms at the interface region, by no less than ~3 eV/atom. The similar preferential B distribution with the assistance of P atoms has been also demonstrated by Guerra et al. and Ni et al. [49,50]. Secondly, from a kinetic perspective, the introduction of P atoms dramatically reduces the energy barriers for B diffusion from 2.5 eV/atom to ~0.7 eV/atom, thus facilitating B species migration into the Si NC. Dynamically, we anticipate that O vacancies will also play an important role in assisting the dopant diffusion, especially at elevated temperature. Thus, we have also evaluated the energy profile for an O vacancy in both the SiO2 matrix and the interface region in Fig. 7. We find that, energetically, the O vacancy exhibits a slight preference to be at the interface over the SiO2 matrix, by ~0.2 eV/vacancy. Nevertheless, we believe that, kinetically, the O vacancy plays an important role in affecting the
doping and the elemental distribution of the dopants in Si NCs as the presence of vacancy will significantly facilitate the diffusion process [54]. We now return to further discuss the APT results. It is worth mentioning that the B concentration at the surface of the Si NCs dispersed in the SiO2 matrix was slightly higher than that of the Si NCs in BPSG (Fig. 5c). This result implies that more B atoms remain inside or near the surface of the colloidal Si NCs, which is in a good agreement with the previous Raman and X-ray photoelectron spectroscopy (XPS) results. Therefore, we propose that the HF etching process preferentially takes place at SieO, SieSi, SieH, SieP bonds rather than at the SieB bond at the surface of the Si NCs. Similar behavior has been demonstrated experimentally for the case of HF etching of porous Si [55]. As shown in our previous paper [12], the resistance to HF solution for doped Si NCs is much stronger than undoped Si NCs, showing that the etching time becomes longer in the order of the undoped Si NCs, the solely-P-doped Si NCs, the solely-B-doped Si NCs, and the B and P co-doped Si NCs. The B and P co-doped Si NCs exhibits highest resistance to the HF solution and we suggest that this is an important indicator as to why the B and P co-doping is essential to fabricate colloidal Si NCs with strong photonic performance. Finally, we discuss the mechanism of dispersibility of our colloidal Si NCs in the methanol, by summarizing the various microstructures observed in the literature. From previous experiments, it has been shown that the z-potential of our co-doped colloidal Si NCs is in the range 30 mV < z < 40 mV [13], indicating that our colloidal Si NC surface is likely negatively charged, and that B and P co-doped Si NCs are stabilized in solution by electrostatically repulsion of each other. On the other hand, undoped, solely-B-doped, and solely-P-doped Si NCs did not form colloids, as demonstrated in a previous study [12]. We suggest that there are two important reasons for this: (i) that the undoped and solely-doped Si NCs are dissolved by HF etching faster than the codoped Si NCs [12,14], and (ii) even if the Si NCs are not fully dissolved during the short time exposure to the HF (e.g. < 10 min), they will aggregate in solution, resulting in their being filtered out by the centrifugation process. By combining our previous APT results [33] with the Raman and XPS results [15], and the DFT calculations here, it is proposed that B atoms are preferentially located around the interface of Si NCs in BPSG, which works more efficiently with the help of P atoms. As shown in Fig. 5, after the HF etching process to fabricate our colloidal Si NCs, the B-enrichment at the surface of the Si NCs was confirmed using APT, indicating that the B-enrichment at the surface of the Si NCs is chemically resistant to HF etching. Considering that the solely-B-doped Si NCs did not effectively form a colloid, B þ P co-doping is needed to enhance the
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B-enrichment at the surface of the Si NCs so as to efficiently protect them from HF etching. In addition to providing HF resistance, B atoms at the Si NC surface are important for their dispersibility in solution. Wheeler et al. has demonstrated that a molecular interaction at the Si NC surface stabilizes Si NCs in solution in the form of colloid [56,57], in an action whereby B atoms behave as a Lewis acid that is attracted to a Lewis base, resulting in dissolution in different kinds of solutions [56,57]. This study also suggests that the existence of B at the Si NCs surface induces a negative surface charge that interacts with a Lewis base such as OH, which enables the Si NCs to be dissolved in methanol or water. 4. Conclusions In conclusion, the atomic-scale microstructure of B and P codoped colloidal Si NCs has been studied using APT experiments and DFT calculations. APT of the colloidal Si NCs was firstly demonstrated and the B and P distribution revealed. Experimentally, it was found that there is a B-rich region at the surface of Si NCs, while the P atoms are more likely to be located inside the Si NCs. Our DFT simulations of Si NCs in a SiO2 matrix indicate that, energetically, P atoms prefer to reside within the Si NCs, and that the presence of P lowers the B diffusion barrier, which then attracts and stabilizes the occurrence of B atoms at the surface of the Si NCs. We suggest that the observed enrichment of B atoms at the surface of Si NCs changes their chemical properties. Firstly, this B-modified surface provides the observed resistance to HF etching. Secondly, the B-modified surface provides a mechanism for the dispersibility of colloidal Si NCs in solutions, enabling the Si NCs to disperse in solution without aggregation. Acknowledgements The authors acknowledge the facilities and the scientific and technical assistance of the Microscopy Australia node at the University of Sydney (Sydney Microscopy & Microanalysis). This research was partially supported by the Faculty of Engineering & Information Technologies at the University of Sydney, under the Research Cluster Program. This research was also undertaken with the assistance of resources from the National Computational Infrastructure (NCI). The NCI and the AMMRF are supported by the Australian Government under the NCRIS program. This collaborative research was also partly supported by the Australian Research Council (grant DP160101713), and the 2018 JSPS Bilateral Joint Research Projects (Japan-Australia), JSPS KAKENHI Grant Number 16H03828 and 18K14092, and the 2015 JST Visegrad Group (V4) Japan Joint Research Project on Advanced Materials.
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Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.actamat.2019.08.013.
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