Colloidal Rh nanocrystals were synthesized through a polyol process involving sequential reactions at two different temperatures (Biacchi and Schaak, 2015). In a typical synthesis, 0.238 g of potassium bromide (KBr, Alfa Aesa), 0.088 g of PVP, 0.024 g of sodium hexachlororhodate (III) dodecahydrate (Na3RhCl6·12 H2O, Alfa Aesa) were mixed with 7 mL of ethylene glycol (EG, Fisher Scientific) in a 20-mL glass vial. Warming up the solution at 40◦C for 1 h dissolved the reagent powders completely. The solution was then heated up to 90◦C and this temperature was maintained for 15 min to initiate nucleation. The reaction was heated to 150◦C and maintained for 1 h, facilitating the growth of colloidal Rh nanocrystals. The resulting dispersion of Rh nanocrystals was mixed with 3 mL of acetone/water (9/1 in V/V), followed by centrifugation at 13,400 rpm for 10 min. The settled powders were then redispersed with 3 mL of acetone/water (9/1 in V/V). Repeating the centrifugation/re-dispersion cycles for 5 times removed PVP and ions from the dispersion of the Rh nanocrystals.

### Functionalization of Silica (SiOx) Nanoparticles

The synthesized SiOx nanoparticles were functionalized with (3 amniopropyl) triethoxysilane (APTES, 98%, Acros Organics) to introduce positively charged surfaces. An ethanolic dispersion of 2 mg mL**–**1 of silica nanoparticles was first prepared with assistance of ultrasonication. To 10 mL of the silica nanoparticle dispersion was dropwise added 0.1 mL of APTES while the temperature of the dispersion was maintained at 60◦C. It took 30 s to complete the addition of APTES. The dispersion was continuously stirred for 8 h, leading to the conjugation of APTES to the surface of the silica nanoparticles. The functionalized silica nanoparticles were collected through centrifugation and washing with ethanol, followed by drying in an oven set at 60◦C. As for r-SiOx NPs, the corresponding powders were calcined at 500◦C for 3 h to burn off PVP from the r-SiOx NPs. An appropriate amount of the calcinated r-SiOx NPs were then added to ethanol, forming a dispersion with a silica concentration of 2 mg mL**–**1 . To 12 mL of the r-SiOx NP dispersion was added 7 mL of aqueous solution of hydrochloric acid (37 wt.%, Fisher Scientific). The dispersion was constantly stirred overnight at room temperature. The pre-treated r-SiOx NPs were then collected by centrifugation and washing with ethanol for further surface modification with APTES.

The APTES-modified silica nanoparticles exhibited positively charged surfaces, to which metal nanocrystals with negatively charged surfaces could be attached through strong electrostatic attractions (Zhang et al., 2016; Rasamani et al., 2017, 2018). In a process of loading 1 wt.% Pt nanocrystals to the silica NPs, 4.21 mL of the as-synthesized Pt nanocrystal solution was slowly added to 10 mL aqueous dispersion of the silica nanoparticles with a concentration of 2 mg mL−1 . Constantly stirring the dispersion of two types of nanoparticles at 600 rpm for 15 min resulted in the attachment of the Pt nanocrystals to the surface of the silica nanoparticles, forming composite SiOx-NP/Pt particles. The obtained SiOx-NP/Pt particles were collected via centrifugation for 10 min at 6,000 rpm, followed by drying in an oven set at 60◦C for 2 h. The same method was also used to load the synthesized Rh nanocrystals to the APTESmodified silica nanoparticles. The loading of metal nanocrystals was tuned by controlling the amount of metal nanocrystals and silica nanoparticles used in the synthesis.

#### MATERIALS CHARACTERIZATION

Transmission electron microscopy (TEM) images of metal nanocrystals were recorded on a microscope (JEOL JEM-1400). Scanning electron microscopy (SEM) images of silica nanoparticles and silica/metal composite particles were characterized with a field-emission microscope (FEI Quanta FEG 450) that was operated at an acceleration voltage of 20 kV. Inductively coupled plasma atomic emission spectroscopy (ICP-OES, Thermo Scientific 7000 Series) was used to determine the loading contents of metal nanocrystals on the silica/metal composite particles. Diffuse reflectance spectroscopy (DRS) was analyzed by an ultraviolet–visible (UV–vis) spectrophotometer (Thermo Scientific, Evolution 220) equipped with an integration sphere.

#### RESULTS AND DISCUSSION

Silica nanoparticles with large enough sizes can always produce random scattering in the visible spectral region regardless of their morphology, but generating resonant scattering strongly depends on their morphology. For instance, theoretical modeling and calculations have shown that spherical SiO2 particles of several microns supports Fabry-Perot or Whispering Gallery resonances, forming electrical fields near the SiO2 surfaces much stronger than that of the incident light (Ausman and Schatz, 2008). Varying the size of the SiO2 particles tunes the resonant wavelengths and the enhancement of electrical fields near the particle surfaces. Decreasing the size of silica spheres down to the sub-micrometer scale still supports surface scattering resonances despite the broadness of the resonance peaks. The light scattering resonances on the surfaces of silica nanospheres (SiOx NSs) significantly enhance the optical absorbance and change the spectral profile of absorption in quantum-sized Pt nanocrystals (with size <10 nm) that are attached to the surface of the SiOx NSs (Zhang et al., 2016). An aqueous dispersion of well-dispersed Pt nanocrystals with an average size of 3.2 nm exhibits a brown color and a peakless absorption spectrum in the range of 300– 800 nm (**Figure 1A**). The Pt nanocrystals have negatively charged surfaces, which allows them to attach to the surfaces of positively charged SiOx NSs through the strong electrostatic attraction. This hybridization process results in a uniform distribution of the Pt nanocrystals on the SiOx NSs, forming SiOx-NS/Pt composite particles (**Figure S1**). A powder of the SiOx-NS/Pt particles is greenish (inset, **Figure 1B**), corresponding to the strong light absorption around 530 nm. **Figure 1B** compares the SiOx-NS/Pt particles and the SiOx NSs with regard to the DRS spectra, which exclude the contribution of scattering and are only sensitive to the absorption. The negligible signal of the SiOx NSs indicates that the absorption signal of the SiOx-NS/Pt particles originates only from the Pt nanocrystals. Different from the freestanding Pt nanocrystals dispersed in an aqueous solution, the SiOx-NS/Pt particles exhibit well-defined peaks, which are consistent with the light scattering resonances on the surface of the SiOx NSs with the highest geometric symmetry (R3) (Zhang et al., 2016). The DRS spectrum of the SiOx-NS/Pt particles also exhibits a nonzero baseline, more likely independent of the wavelength, which agrees with the feature of random scattering. The comparisons indicate that both resonant scattering and random scattering of the dielectric SiOx NSs are responsible for enhancing the light absorption in the quantum-sized Pt nanocrystals. The respective contribution of resonant scattering and random scattering to the enhanced light absorption is not distinguished.

Theoretical modeling and calculations indicate that the spherical geometry of dielectric particles favors resonant scattering and reducing the geometrical symmetry of the dielectric particles weakens resonant scattering (Gupta et al., 2018). **Figure 2A** highlights that the geometrical symmetry of silica nanoparticles can be lowered by tuning the concentration of CTAB in the sol-gel reaction as described elsewhere (**Figure S2**) (Hao et al., 2015; Rahmani et al., 2017). Highly symmetrical SiOx NSs are formed by using the modified Stöber method (**Figure S1**). Adopting a different synthetic procedure with addition of CTAB yields SiOx with varied morphologies. Addition of a low concentration (4.9 mM) of CTAB to the reaction solution slightly truncates the two poles of SiOx NSs, forming silica nanoparticles (labeled as t-SiOx NSs shown in **Figure 2B**) with a lower geometric symmetry than the SiOx NSs. Adding more CTAB further truncates the silica nanoparticles and silica nanoparticles with an ellipsoidal shape (labeled as e-SiOx NPs, **Figure 2C**) are formed with an extremely high concentration (9.8 mM) of CTAB. Same as the SiOx NSs, both the t-SiOx NSs and e-SiOx NPs also exhibit negligible absorption signal in the corresponding DRS spectra (purple and pink curves, **Figure 2D**). In contrast, the Pt nanocrystals on the t-SiOx NSs and e-SiOx NPs exhibit strong light absorption (red and blue curves, **Figure 2D**), indicating that both the t-SiOx NSs and e-SiOx NPs are still capable of enhancing light absorption of the Pt nanocrystals due to light scattering of the silica nanoparticles. The corresponding DRS spectra become essentially peakless, implying the absence of resonant light scattering on the t-SiOx NSs and e-SiOx NPs with lowered geometric symmetry of D∞h. The spectral difference of the Pt nanocrystals on the differently shaped silica nanoparticles highlights the importance of geometric symmetry of the silica nanoparticles on determining their light scattering mode. Resonant scattering, which is responsible for the appearance of well-defined intense absorption peaks in the DSR spectra, is very sensitive to the geometric symmetry of silica nanoparticles. Only the SiOx NSs with the

highest geometrical symmetry of R3 supports strong surface scattering resonances. Lowering the geometric symmetry of the silica nanoparticles drastically suppresses the resonant scattering while random scattering is barely influenced (**Figure 2D**). The difference in the DRS intensity shows that the resonant scattering and random scattering of SiOx NSs make approximately equal contributions to enhance the light absorption in the Pt nanocrystals at the resonance frequencies. At the non-resonance frequencies, the random scattering dominates the enhancement of light absorption in the Pt nanocrystals.

Light scattering on the SiOx NSs can enhance optical absorption in any nanoparticles attached to the SiOx NSs (Zhang et al., 2016; Rasamani et al., 2017, 2018; Dai et al., 2018b). For example, quantum-sized Rh nanocrystals with sizes of ∼3 nm exhibit a peakless absorption spectrum while strong light absorption with well-defined peaks is observed from the Rh nanocrystals loaded to the SiOx NSs (with an average

diameter of 356 nm, **Figure S3**). To further verify the importance of spherical symmetry of the silica nanoparticles on resonant scattering, another set of silica nanoparticles with reduced geometrical symmetry are synthesized by controlling the aspect ratio of the rodlike particles. The silica nanoparticles start nucleation at nanosized water-rich emulsion droplets followed by anisotropic growth along the direction away from emulsion droplets, forming rodlike nanoparticles with one flat end and one rounded end (**Figure 3A**, **Figure S4**). These rodlike silica nanoparticles (r-SiOx NPs) exhibit a geometric symmetry of C∞V and their geometric aspect ratio is determined by the lateral dimensions shown in **Figure 3A**. Similar to the SiOx NSs, these r-SiOx NPs are also feasible to attract the quantum-sized Rh nanocrystals to their surfaces, forming SiOx/Rh composites. **Figures 3B,C** show the SEM images of bare silica and composite samples (insets) formed from the r-SiOx NPs with aspect ratios

of 0.8 and 1.4, respectively. The corresponding powders of these composite particles exhibit strong optical absorption although their DSR spectra are different from that of the SiOx-NS/Rh composite particles (**Figure 3D**). The spectral difference shows that the intensity of the absorption peaks decreases as the aspect ratio of the r-SiOx NPs deviates from the unity (i.e., 1) that corresponds to the aspect ratio of the SiOx NSs. This relationship can be quantitatively compared by integrating these DRS spectra in the range of 300–800 nm, and the integrated values are shown in **Figure S5** as a function of aspect ratio of the supporting SiOx NPs. The volcano shape with a maximum at the aspect ratio of 1 (corresponding to spherical SiOx NPs) again highlights that the SiOx NSs are more effective in enhancing light absorption in the Rh nanocrystals than the rodlike silica nanoparticles regardless of the aspect ratios >1 or <1. A more deviation of the aspect ratio weakens the absorption peaks more, further confirming that high geometric symmetry of the SiOx

NSs is crucial to support strong resonant scattering on the dielectric silica nanoparticles. Regardless of the aspect ratio of the silica nanoparticles, the featureless baselines of these DRS spectra remain essentially consistent, indicating that the random scattering of the silica nanoparticles makes an approximately constant contribution to enhance the light absorption in the Rh nanocrystals. At the resonance frequencies, the light absorption in the Rh nanocrystals enhanced by the resonant scattering of the SiOx NSs is comparable to the absorption enhanced by the random scattering. The light absorption is dominated by the enhancement originated from the random scattering at the nonresonance frequencies. The DRS spectra shown in **Figures 2D**, **3D** consistently highlight that both resonant scattering and random scattering of the silica nanoparticles are capable of enhancing the light absorption in the quantum-sized metal nanocrystals attached to the silica nanoparticles. The occurrence

of resonant scattering strongly depends on the geometric symmetry of the silica nanoparticles. The SiOx NSs with the highest geometric symmetry supports the strongest resonant scattering while the resonant scattering weakens with decrease of their geometric symmetry. In contrast, the random scattering is independent of the geometry of the silica nanoparticles.

The light scattering efficiency of the dielectric silica nanoparticles also relies on the refractive index (n) of the surrounding environment (Mishchenko et al., 2002; Zhu et al., 2010). A large difference of refractive indexes between the surround medium and the silica nanoparticles (nsilica ≈1.4–1.45 for the sol-gel silica) promotes light scattering. Therefore, the silica nanoparticles dispersed in water (nwater = 1.33 at room temperature) exhibit a much lower light scattering efficiency compared to the dry powder of the silica nanoparticles in air (nair = 1). The variation of refractive index of the surround medium also influences the scattering resonance frequencies of the SiOx NSs. **Figure 4A** compares the DRS spectra of aqueous dispersions of freestanding Rh nanocrystals and SiOx/Rh composite particles formed with varying silica nanoparticles, which exhibit profiles different from the DRS spectra of the corresponding SiOx/Rh dry powders (**Figure 3D**). However, the light absorption in the Rh nanocrystals is always enhanced when the Rh nanocrystals are attached to the silica nanoparticles. The good dispersion of the Rh nanocrystals in water and on the surfaces of the silica nanoparticles ensures that the measured DRS signals represent the accumulation of optical absorption of individual Rh nanocrystals with exclusion of possible interparticular coupling between the adjacent Rh nanocrystals. As a result, the enhancement spectrum, i.e., enhancement factor as a function of wavelength, enabled by the silica nanoparticles can be calculated by dividing the DRS spectrum of a dispersion of the corresponding SiOx/Rh composite particles against the DRS spectrum of the dispersion of the freestanding Rh nanocrystals (green dotted curve, **Figure 4A**). **Figure 4B** presents the calculated enhancement spectra enabled by the SiOx NSs and the r-SiOx NPs with aspect ratios of 0.8 and 1.4, showing that the maximum enhancement factor can reach ∼12 around 750 nm for the SiOx NSs. In the visible spectral region (i.e., at λ > 450 nm), the enhancement factor is always higher than 7 for the SiOx NSs although it is lower for the r-SiOx NPs. Since the light scattering efficiency of the dry SiOx NSs in air is much higher than that of the wet SiOx NSs in solvents (e.g., ethanol, water, etc.), the optical absorption power of the Rh nanocrystals attached to the SiOx NSs could be enhanced by a higher factor when SiOx-NS/Rh composite particles are used in gas atmospheres. In order to achieve the maximum scattering-enhanced light absorption in quantum-sized metal nanoparticles, the synthesis of SiOx NSs has to be carefully controlled to ensure the refractive index of the SiOx NSs to reach the possible maximum value (Li et al., 2014; Zulfiqar et al., 2016; Dai et al., 2018a).

#### CONCLUSION

In conclusion, silica nanoparticles with lateral dimensions of hundreds of nanometers and larger represent a class of dielectric antenna that does not absorb visible light, exhibiting strong light scattering in the visible spectral region with a minimum energy loss. Regardless of the geometry of the silica nanoparticle antenna, random scattering of the incident light is always observed while the wave function is not altered. When the geometry of the silica nanoparticles exhibits a high enough symmetry, the nanoparticle antenna can also support surface scattering resonances that generate new standing waves with wave functions different from the incident light. For example, light scattering resonances on the surface of SiOx NSs with the highest geometric symmetry result in the formation surface "hot spots," at which the new electromagnetic waves exhibit higher power density than the incident light. **Figure 5** schematically highlights the dependence of scattering mode on the geometry of the silica nanoparticles. The highly symmetric SiOx NSs support both resonant scattering and random scattering while the silica nanoparticles with low geometric symmetry (e.g., r-SiOx NPs with large aspect ratios shown in **Figure S4**) support merely random scattering. When light-absorbing species are attached to the surface of the silica nanoparticles, their optical absorption power can be enhanced by both random scattering and resonant scattering. The DRS spectra of the SiOx/Pt and SiOx/Rh composite particles shown in **Figures 2D**, **3D** represent the typical examples highlighting that the optical absorption of quantum-sized metal nanocrystals is significantly enhanced by light scattering (including both random scattering and resonant scattering) on the silica nanoparticles. The random scattering of light from the silica nanoparticles propagates in directions different from that of the incident light, elongating the propagation pathway of light in a material block made of silica/metal composite particles. The longer light pathway allows more incident light to be absorbed by the metal nanocrystals attached on the silica nanoparticles, resulting in an enhancement of overall light absorption. It is worth pointing out that the random scattering does not change the spatial power density of the scattered light. In contrast, resonant scattering of the incident light creates "hot spots" on the surface of SiOx NSs to significantly increase the local electrical fields (corresponding to power density), which enhance the optical (energy) absorption in the metal nanocrystals attached to the surface of the SiOx NSs. The mechanistic difference in enhancing light absorption in quantum-sized metal nanocrystals for random scattering and resonant scattering might be more influential toward non-linear optical properties, for example, photo-excited hot electron generation in the quantum-sized metal nanocrystals (Wei et al., 2018). The understanding sheds light on designing composite materials with the dielectric silica nanoparticle antenna to promote performance of applications.

#### AUTHOR CONTRIBUTIONS

YS, XD, and KR: experimental design, data analysis and interpretation, manuscript writing, and manuscript revision; XD, KR, GH, and RM: material synthesis and characterizations, data acquisition.

# ACKNOWLEDGMENTS

This work is supported by the startup from Temple University.

#### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2018.00494/full#supplementary-material

# REFERENCES

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2018 Dai, Rasamani, Hall, Makrypodi and Sun. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Encapsulation of Metal Nanoparticle Catalysts Within Mesoporous Zeolites and Their Enhanced Catalytic Performances: A Review

Dongdong Xu, Hao Lv and Ben Liu\*

Jiangsu Key Laboratory of New Power Batteries, Jiangsu Collaborative Innovation Center of Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal University, Nanjing, China

#### Edited by:

Jie He, University of Connecticut, United States

#### Reviewed by:

Gonghu Li, University of New Hampshire, United States Haoquan Zheng, Shaanxi Normal University, China Heriberto Pfeiffer, Universidad Nacional Autónoma de México, Mexico

> \*Correspondence: Ben Liu<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Nanoscience, a section of the journal Frontiers in Chemistry

Received: 31 July 2018 Accepted: 23 October 2018 Published: 09 November 2018

#### Citation:

Xu D, Lv H and Liu B (2018) Encapsulation of Metal Nanoparticle Catalysts Within Mesoporous Zeolites and Their Enhanced Catalytic Performances: A Review. Front. Chem. 6:550. doi: 10.3389/fchem.2018.00550 Metal nanoparticles (NPs) exhibit desired activities in various catalytic reactions. However, the aggregation and sintering of metal NPs usually cause the loss of catalytic performance in practical reaction processes. Encapsulation of catalytically active metal NPs on/within a high-surface-area inorganic support partially resolve such concerns. Microporous zeolites, owing to their rigid frameworks and porous structural features, have been considered as one of ideal inorganic supports. Metal NPs can be easily encapsulated and stabilized within zeolitic frameworks to prevent unwished aggregation during the catalysis. Unfortunately, sole microporous nanochannels (generally <1 nm) in conventional zeolites are not easy to be accessed. The introduction of another set of nanochannel (e.g., mesopore), known as mesoporous zeolites, can greatly improve the mass-transfer efficiency, which is structurally beneficial for most catalytic reactions. The coexistence of micropores and mesopores in inorganic supports provides the synergetic advantages of both fine confinement effect for metal NPs and easy diffusion for organic reactants/intermediates/products. This review focuses on the recent advances in the design and synthesis of mesoporous zeolites-encapsulated metal NP catalysts as well as their desired catalytic performances (activity and stability) in organic reactions. We first discuss the advantages of mesoporous zeolites as the supports and present general strategies for the construction of mesoporous zeolites. Then, the preparation methods on how to encapsulate NP catalysts within both microporous and mesoporous zeolites are clearly demonstrated. Third, some recent important cases on catalytic applications are presented to verify structural advantages of mesoporous zeolite supports. Within the conclusion, the perspectives on future developments in metal NP catalysts encapsulated within mesoporous zeolites are lastly discussed.

Keywords: metal nanoparticle, zeolite, mesoporous, nanoconfinement, mass transfer

# INTRODUCTION

Nanocatalysis by metal nanoparticles (NPs) have always attracted sufficient attention from both scientific researches and industrial applications (Astruc et al., 2005; Narayanan and El-Sayed, 2005; Astruc, 2008; Campelo et al., 2009; Cuenya, 2010). Owing to the high surface-to-volume ratio and surface energy, the surface atoms on metal NPs are supposed to be catalytically active for heterogeneous catalytic processes, such as hydrogenation, C-C coupling reaction, carbonylation, oxidation, methanol reforming, fuel cell and so on (Hartwig, 1998; Dyker, 1999; Dupont et al., 2002; Moreno-Manas and Pleixats, 2003; Na et al., 2004; Gilroy et al., 2016). However, precisely because of the high active surface atoms (unstable thermodynamic state), the aggregation and deactivation of the metal NPs by secondary nucleation and recrystallization (or Ostwald ripening) suffer the most serious drawbacks when employed in practical catalytic reactions and hence set great limits on their widely industrial applications (Schmid et al., 1996; Doyle et al., 2003; Challa et al., 2011; Schauermann et al., 2012; Hansen et al., 2013). As a result, the stabilization of metal NPs becomes a challenging research area and can be mainly realized via the combination with soft organic species and hard (or solid) supports (Astruc et al., 2005; Cao et al., 2010). The soft species, also known as capping reagents, are usually classified as polymers (including dendrimers; Corbierre et al., 2001; Crooks et al., 2001), organic or inorganic ligands (Pan et al., 2001), surfactants or micelles (Liz-Marzán and Lado-Touriño, 1996; Kitchens et al., 2003), microemulsions (Capek, 2004), ionic liquids (Dupont et al., 2002; Dupont and Scholten, 2010) and so on. By means of direct chemical binding (generally through heteroatoms) weakly to the surface of metal NPs, these soft protectors will availably block the NP aggregation but also result in the loss of active sites in the meantime due to the subsistent coverage or interfacial interaction (Jin et al., 2017). Additionally, how to resolve the essential question about the recycling potential of these capping metal NPs is always a big challenge. Another significant strategy focuses on the encapsulation of small metal NPs on some hard or solid supports, which will not only prevent NP migration and coalescence but also greatly increase the recycling accessibility (Astruc et al., 2005; Gallon et al., 2007; Campelo et al., 2009; White et al., 2009). Furthermore, the electronic properties of metal NPs on their host support environments can be greatly modified and hence will give rise to various special physical and chemical functions owing to the strong interactions between metal NPs and supports (Ju-Nam and Lead, 2008; Liu et al., 2015, 2017b). Numerous types of nanomaterials can play the role of solid supports, mainly including metal or non-metallic oxides (e.g., TiO2, ZrO2, CeO2, ZnO, Al2O3, Fe3O4, SiO2); (Astruc et al., 2005; Campelo et al., 2009), carbon (e.g., carbon sphere, nanotube, graphene, porous carbon); (Joo et al., 2001; Sun and Li, 2004), functionalized polymers (Kralik and Biffis, 2001), and porous framework (MOFs, COFs, zeolites, silica); (Dhakshinamoorthy and Garcia, 2012; Farrusseng and Tuel, 2016; Zhu and Xu, 2016). The oxide supports can classified as inert (e.g., SiO2) and reactive types based on their chemical reactivity. For example, CeO2 is a representative kind of reactive oxide support as an enhanced anchoring effect of Ce-O-M born of strong metal-support interaction can efficiently suppress the phase transformation and improve the thermal stability of metal NPs under higher temperature (Alessandro, 2002). When superparamagnetic Fe3O4 is utilized as the support, the metal NP/Fe3O4 nanomaterials can be easily recovered using a permanent magnet and therefore reused in many runs (Jacinto et al., 2008). Carbon nanostructures, owing to their intrinsic properties, such as high surface area, high electrical conductivity, and unique physical properties and so on, are extremely attractive supports and have been widely employed in the heterogeneous catalysis process (Wildgoose et al., 2006; Kamat, 2009; Liu et al., 2016, 2018). Furthermore, the carbonaceous surface can be easily modified to stabilize the metal NPs. In recent years, the combination between porous frameworks and metal NP technologies has attracted sufficient research and obtained fruitful results (Joo et al., 2001; He et al., 2003; Yang et al., 2003; Mandal et al., 2004; Rioux et al., 2005; White et al., 2009; Aijaz et al., 2012; Lu et al., 2012; Xu et al., 2012). A porous framework generally composed of interconnected network system of pores with distinct diameters, which exist in many natural substances and can be easily created in most of solid chemical nanomaterials. The pores inside the solid structures (usually called interior pore system, while exterior surface can only be accessed in non-porous structures) can nanoconfine the size and shape of metal NPs and inhibit their aggregation and further growth. The throughout pore network renders the masstransfer more achievable during the catalysis (Satterfield, 1970; Liu et al., 2010, 2017a). Furthermore, the different components of frameworks, dimensions of porous nanochannels, pore sizes and corresponding combinations (hierarchical porosities) endow porous supports with more functional features (Davis, 2002a; Rowsell and Yaghi, 2004). Mesoporous silica with numerous pore sizes, morphologies, and mesochannels have been successfully synthesized by different strategies and also utilized as the support media of metal NPs (Kresge et al., 1992; Zhao et al., 1998; Ciesla and Schüth, 1999; Wan and Zhao, 2007). Additionally, porous carbon materials and MOFs have also been widely prepared to act as the supports for some particular catalytic applications. Each porous support possesses its intrinsic advantages (e.g., carbon receiving greater thermally stabilization) and disadvantages (for example, complex synthesis processes for mesoporous carbon, low framework stabilization of conventional mesoporous silica and MOFs); (Zhao et al., 2006; Arico et al., 2011; Slater and Cooper, 2015).

Compared to abovementioned support media, porous zeolites, due to their inherent structural features, are promised to be an important family of solid supports for the encapsulation of metal NPs. Zeolites, as a kind of highly crystalline inorganic materials having orderly distributed micropores with the diameter <1 nm, have many applications, especially in the field of petrochemical industry (Corma, 1997; Cundy and Cox, 2005; Li and Yu, 2014; Primo and Garcia, 2014; Zaarour et al., 2014). The zeolite framework commonly consists of TO4 tetrahedra (T denotes Si, Al, and P, etc.) which will build lots of senior structures. Up to 2018, there are totally 239 zeolite framework types (http://www.iza-structure.org/databases/) collected by the International Zeolite Association (IZA) (Guo et al., 2017). Based on different pore windows from TO4, conventional zeolites can be categorized into small, medium, large, and extra-large pores. The catalytic activities of zeolites largely depend on the structural and compositional features, including pore sizes, channel types and framework compositions. Compared to other catalyst supports, zeolites possess an important and unique feature, known as shape selectivity (e.g., reactants, products and transition-states). When utilized as the solid supports, the regular cavities and nanochannels in zeolites can act as the spaces and sites for the encapsulation of extraneous NPs. When metal NPs form and grow inside the micropores, they would be wellconfined at the level of ultrasmall sizes and be not easy to escape from these interconnected cavities. For example, widely studied and applied zeolite ZSM-5 (MFI topology) possesses 10-numbered rings and admits the largest pore accessibility of ca. 0.64 nm (Kokotailo et al., 1978), which will greatly confine the crystalline growth of embedded metal NPs. Rigid framework of zeolites offers the important advantages for their applications in harsh catalytic environments; their unique nanochannels also provide the opportunity for shape-selective catalysis (Corma, 1997; Weitkamp, 2000; Davis, 2002a; Moller et al., 2012). So far, many reports on the preparation of zeoliteencapsulated metal NPs show the enhanced catalytic activities. For example, Wang et al. (2016) reported the preparation of ultrasmall Pd NPs encapsulated in situ within nanosized silicalite-1 by one-step hydrothermal synthesis. The resultant catalyst showed a highly efficient H2 generation activity toward the decomposition of formic acid. Liu et al. (2017a) prepared single Pt atoms and Pt NPs with exceptionally high thermal stability within purely siliceous MCM-22 and the catalyst presented good catalytic activity in the dehydrogenation of propane to propylene. Zhang et al. (2017) found that Pd@Beta catalyst exhibited superior selectivity for the hydrogenation of the nitro group when fixing Pd NPs inside Beta zeolite crystals.

It is easy to envision that the mass-transfer processes will be greatly restricted as the micropores have been occupied by metal NPs especially involving organic molecules with larger dimensions (Kärger and Ruthven, 1992; Chen et al., 1994; Tao et al., 2006). For purpose of overcoming this shortcoming, massive efforts have been made to construct zeolitic frameworks with larger aperture (Li and Yu, 2014) or create a new set of pore system (known as mesopores) inside the conventional bulk zeolites (Egeblad et al., 2008; Perez-Ramirez et al., 2008; Lopez-Orozco et al., 2011; Möller and Bein, 2013; Parlett et al., 2013; Perego and Millini, 2013; Serrano et al., 2013). Comparatively speaking, the latter is simpler to implement. As shown in **Figure 1A**, purely microporous zeolites (e.g., ZSM-5), ordered mesoporous materials (e.g., MCM-41), and mesoporous zeolites (e.g., meso-ZSM-5) show obviously different Brunauer– Emmett–Teller curves and pore size distributions based N2 isotherms (Perez-Ramirez et al., 2008). It clearly illustrate that mesoporous zeolites possess both the micropores and mesopores (hierarchical pore systems) inside one crystal, which would be greatly helpful for the mass thransfer of organic molecules. In mesoporous zeolites-encapsulated metal NPs catalysts, metal surfaces (providing the catalytic active sites), the micropores or mesopores (immobilizing and stabilizing the metal NPs) and interconnected mesopores (accelerating the transfer of reactants and products) together built an excellent catalysis system. As schemed in **Figure 1B**, the metal NPs can be immobilized inside both micropores and mesopores. The reactants go through the catalysts along the interconnected mesopores and accelerate the organic reactions. The interior metal NPs also participate in the catalysis, which is not accessible in sole microporous systems.

Considering all of above aspects, this review will highlight the recent progresses on the construction of mesoporous zeolitesencapsulated metal NPs and the investigation of their enhanced activities on various catalytic applications. Conventional zeolitesencapsulated counterparts are not discussed in detail here as some excellent review papers have presented before (Farrusseng and Tuel, 2016; Wang et al., 2018b). In current review, we firstly concern design principles and construction of mesoporous zeolites. Second, we place special focus on how to immobilize or introduce metal NPs into the mesoporous zeolites. Some reported cases will be clearly presented to illustrate the synthetic process of this type of composite catalysts. Third, several typical organic reactions based on mesoporous zeolitesencapsulated metal NPs are recommended to confirm their superior catalytic performances. This review is prospected to provide some important insights on the construction of

stabilized metal NPs within mesoporous zeolites and how to tune the catalytic performances based on the distinct supports media.

# CONSTRUCTION OF MESOPOROUS ZEOLITES

Considering the advantages of hierarchically mesoporous zeolites in various catalytic reactions, many attempts have been made to build the meospores in conventional zeolites. Furthermore, the preparation of mesoporous zeolites is the precondition for the encapsulation of metal NPs.  The main methods can be generally classified as four aspects: (i) hardtemplating, (ii) soft-templating, (iii) post-synthetic treatment, and (iv) stacking or assembly of zeolite crystals. As shown in **Figure 2**, the strategies and the detailed routes are clearly presented schematically.

Most of removable and inert solids with the nanoscale sizes can be employed to act the hard templates for the synthesis of mesoporous zeolites by a hard-templating method, for example, carbon nanostructures (NPs, nanowires, nanotubes, aerogel, ordered porous frameworks), CaCO3 NPs, polystyrene arrays, biological wood and starch (Schwieger et al., 2016). After the removal of these solid species by calcination or acid dissolution, abundant mesopores can be released while the order degree of these mesopores mainly depends on the nanostructures of hard templates. As an instance, three-dimensional ordered mesoporous carbon with large mesopores was developed to construct ordered porous zeolites with distinct framework topologies (e.g., MFI, LTA, FAU, BEA); (Chen et al., 2011). Nevertheless, although the synthesis of mesoporous zeolites using hard-templating methods is greatly high-efficiency, this strategy is often restricted due to the multistep procedures, high costs and health hazards involving during the synthesis.

Soft-templating strategy, referring to the molecular selfassembly of amphiphilic surfactants, is considered as the most possible way to build ordered mesoporous zeolites. The ordered mesopores are greatly favorable for the mass-transfer in the catalytic processes. Choi et al. (2009) and Na et al. (2011) have successfully designed a series of amphiphilic molecules containing quaternary ammoniums, also known as bifunctional surfactants, which were further employed to direct the growth of zeolitic frameworks and ordered mesoscale structures by one-step hydrothermal synthesis. As a consequence, lamellar MFI zeolite nanosheets with a single-unit-cell thickness along b-crystallographic axis and a hexagonally-ordered mesoporous aluminosilicate, consisted of truly crystalline zeolitic walls, have successfully obtained. By means of the introduction of aromatic groups into the hydrophobic chain, Xu et al. (2014a,b) prepared single crystalline mesostructured zeolite nanosheets and the ones with 90◦ rotational boundaries due to the π-π stacking between benzene rings. Additionally, Liu et al. (2012) and Zhu et al. (2014) developed the hydrophilic cationic polymer (e.g., polydiallyldimethylammonium chloride, polystyrene-co-4-polyvinylpyridine) to produce single-crystalline mesoporous zeolite Beta and MFI, respectively. The soft-templating method provides the possibility for the better control of both zeolitic framework topologies and mesoporous structures. However, the current synthesis is limited to zeolite MFI or Beta, and therefore how to design the effective organic surfactants for synthesizing other kinds of zeolitic topologies or mesoporous nanostructures is still a big challenge.

The selective removal or extraction of zeolite framework atoms (e.g., Si, Al, or Ti) via post-treatment by acid, alkali, steaming, or other similar methods, can produce some mesopores inside the bulk zeolites. Through the precise control of the desilication or dealumnization from the zeolite frameworks, one could obtain mesoporous zeolites with different Si/Al ratios and desired catalytic sites. This method can effectively result in abundant intracrystalline mesopores but simultaneously decrease the crystallinity of zeolites partially (Groen et al., 2005). The intercrystalline mesopores are a type of mesopores from the stacking of nanocrystals and can be easily found in most of nanomaterials. Through the synthesis of nanosized zeolite crystals, abundant mesopores although in disordered forms but may be destroyed again after the dispersion in the solvents. After the treatment like swelling, intercalation, delamination and pillaring, some layered zeolites can be turned into ultrathin zeolite layers, which further assemble into mesoporous nanostructures. A typical case is the delamination of as-synthesized MCM-22 into two-dimensional ITQ-2, which possess high surface areas and improved catalytic performances (Corma et al., 1998).

Another strategy to construct mesopores in zeolites focuses on the particular crystalline growth model of some particular zeolites. For example, through the intergrowth of MEL/MFI, selfpillared pentasil zeolite with house-of-cards-like morphology was successfully prepared. The hierarchical porosity was created inside the three-dimensional nanosheet assembly by the repetitive branching of the MFI-type zeolitic nanosheets (Zhang et al., 2012). The synthesis process based on this strategy is relatively simple (no need of special templates or tough condition) but only happens in a handful of zeolitic topologies.

Based on the discussion of the strategies to construct mesopores inside zeolites, the different synthetic routes obviously possess their intrinsic advantages and disadvantages. One should consider these processes comprehensively before the choice of a certain route and further combine the corresponded application fields. For instance, when utilized as a porous support, all of routes are suitable for the post-synthetic encapsulation of metal NPs (discussed in next part) but not for the one-step formation of mesoporous zeolite-encapsulated metal NPs composites.

#### ENCAPSULATION STRATEGIES FOR METAL NPS

The encapsulation of metal NPs inside the porous zeolites can be mainly classified into two types of strategies: post-synthetic process and one-step in situ confinement. The post-synthetic

strategy refers to the introduction of metal NPs after the complete construction of zeolitic framework. Contrastively, the one-step confinement method needs the co-crystallization of zeolites and metallic precursors, and follows an in-situ reduction to obtain metal NPs.

#### Post-synthetic Encapsulation Strategy

The post-synthetic strategy is widely utilized to introduce metal NPs into zeolite crystals as it does not limit the framework sorts of zeolites. The process based on this strategy can be reached by means of the soak of zeolite supports in metal colloids (or NPs) or soluble metal precursors. The channels and pores inside zeolites provide the region for the loading of metal NPs. The diffusion of metal NPs or precursors is usually accelerated in a set of relatively large micropores, especially in mesopores. Actually, the diffusion of metal NPs is quite difficult to enter into the interior structures of zeolites where most of them only load at the surface (Wang et al., 2018a). As a representative example, by means of a particular structural transformation of two-dimensional (2D) zeolite into threedimensional (3D), Liu et al. (2017c) had successfully imbedded single Pt atoms and Pt NPs inside siliceous MCM-22 (Pt@MCM-22). As shown in **Figure 3**, during the swelling process of the lamellar precursor MCM-22 where MWW layers were expanded by the surfactant hexadecyltrimethylammonium, subnanometric Pt species were incorporated into internal channels between individual MWW layers. After the removal of organic species by calcination, lamellar MCM-22 would turn into 3D MCM-22 and simultaneously Pt species were well-confined in external cups (hemi-cages) or encapsulated in the supercages. Subnanometric Pt species were tightly entrapped inside zeolite MCM-22. Due to structural features, Pt@MCM-22 possessed exceptionally high thermal stability even after the treatment in air up to 540◦C and showed desired size-selective catalytic activity for the hydrogenation of alkenes.

Most of cases for the preparation of zeolite-encapsulated metal NP catalysts focus on the strategy using soluble metal precursors, also called wetness impregnation method. In this process, metallic species can easily diffuse into the crystal inner and load in the micropores or mesopores after the removal of solvent (e.g., water). Following a conventional reduction step, the metal precursors can be reduced into elementary metal NPs. A synthesis methodology using thiol-based organometallic complexes was presented to encapsulate metal NPs within LTA zeolites, which were successful in encapsulating 1–2 nm monometallic (e.g., Pt, Pd, Ir, Rh, and Ag) or alloyed NPs (e.g., AuPd, AuPt, and PdPt); (Choi et al., 2010; Otto et al., 2016). In pure microporous zeolites, limited by the small diameter of micropores, the size of resultant metals was usually confined in the region of below 2 nm. As for mesoporous zeolite supports, it is more easy to immobilize metal NPs owing to the presence of the larger mesopores. So far, many kinds of metal NPs (Pt, Pd, Ru, Ag, and the corresponded alloys) has been immobilized in mesoporous zeolites with different framework topologies based on this wetness impregnation method (Wang et al., 2015, 2018a; Mendes et al., 2017; Chen et al., 2018; Zhang et al., 2018). Due to the intrinsically structural advantages of mesoporous zeolites and confinement effects for the immobilized metal NPs, all of these reported composite catalysts had exhibited superior catalytic activity in various catalytic reactions (Liu and Corma, 2018).

# In situ Confinement Strategy

Metal NPs or precursors can also be introduced into the zeolite crystal inners via the one-pot hydrothermal synthetic process. In this method, as-synthesized metal NPs or soluble metallic precursors were firstly mixed with the synthetic gel for zeolites (e.g., structural-directing agents, silica resources, water, sodium hydroxide, etc.), and then turned into high temperature for the crystallization of zeolites. The as-synthesized composite products were further calcined to remove the organic species and reduced under reduction agents (e.g., H2, NaBH4) to form metal NPs. Laursen et al. reported the synthesis of zeolite-encapsulated Au NPs for size-selective satalysis (Laursen et al., 2010; Højholt et al., 2011). Ultrasmall Au NPs were uniformly immobilized in silica

FIGURE 3 | Schematics of the processes for the preparation of Pt@MCM-22. Reprinted with permission from Liu et al. (2017c). Copyright 2016 Macmillan Publishers Ltd: (Nature Materials).

with permission from Wang et al. (2016). Copyright 2016 American Chemical Society.

matrix that would be turned into a zeolite phase (silicalite-1). Au NPs embedded inside slicalite-1 mostly possessed a diameter of 1–2 nm; no obvious change in the original size can be seen even under calcination under 500◦C. The crystalline zeolite frameworks played the role to confine the growth of Au NPs. However, in this process, a phase separation phenomenon often

happened due to the harsh crystalline conditions for zeolites, resulting in the formation of larger metal aggregates. Therefore, the usage of metal NPs as the precursors is not widely applied for the construction of zeolite-encapsulated metal catalysts.

Comparatively speaking, metallic precursors are easy to be encapsulated inside the resultant zeolites. Wu et al. (2014) reported the encapsulation of noble metal NPs (Pt, Pd, Rh, Ir, Re, and Ag) in LTA voids (**Figure 4A**). If only metal cations were employed, premature precipitation of these metal hydroxides would be mainly formed. When some ligands (e.g., NH3, ethylenediamine) were used to stabilize metal cations and protect cationic moieties against precipitation, avoiding premature precipitation resulted in the zeolitic framework units to grow around the ligand-stabilized metal precursors. Finally, a series of metal NPs were successfully encapsulated within the small micropores of LTA zeolite. These composite materials exhibited excellent shape selectivity in both catalytic oxidative dehydrogenation of alkanols and hydrogenation of alkenes. Another important case by using this strategy was reported by Wang et al. (2016). The authors utilized organometallic complex ([Pd(NH2CH2CH2NH2)2] 2+) as the Pd precursor to produce well-dispersed and ultrasmall Pd NPs in nanosized silicalite-1 (MFI). Pd NPs were successfully encapsulated within the intersectional channels of MFI with a diameter of ∼1.8 nm (**Figure 4B**), which could be further tuned via the different synthetic conditions (using NaOH or KOH as the base source). They had also reported the preparation of subnanometric bimetallic NPs of Pd-M(OH)2 (M=Ni, Co) using the similar synthetic methods (Sun et al., 2017). These MFI-encapsulated metal NPs presented superior catalytic performance toward the dehydrogenation of formic acid.

From abovementioned cases, how to balance the crystallization growth of both zeolites and metal NPs is quite important for obtaining the well-dispersed and size-uniform metal NPs inside zeolite crystals. The type of metallic precursors and other synthetic conditions (e.g., alkalinity, silica sources) should be carefully considered. The corresponded catalytic activity may generally depend on how strong the confinement effect is between zeolite frameworks and metal NPs.

#### CATALYTIC APPLICATIONS

When metal NPs are encapsulated inside conventional zeolites, they are mainly confined in micropores and randomly distributed throughout the zeolite crystals. This type of catalysts proves to be appropriate for organic catalysis involving only small molecules. For example, when Pd cluster-containing zeolite catalysts (Pd/S-1, as shown in **Figure 4B**) was utilized to catalyse the dehydrogenation of formic acid and sodium formate (FA-SF) to produce H2 (**Figure 5**), Pd/S-1 exhibited complete decomposition of FA in a short time compared to the commercial Pd/C catalyst, also better than the Pd/S-1 catalyst obtained by the impregnation-prepared catalysts (Wang et al., 2016). However, although exhibiting excellent catalytic activity in the dehydrogenation process, this catalyst may miss their ability in organic reactions for the larger molecules, due to the limitation of small microporous channels or accessible void spaces. Additionally, even if the catalytic reactions involving organic molecules with small diameters (less than the micropores) happens in this type of composite catalysts, relatively slow diffusion in the interior zeolites would greatly influence the catalytic processes.

On the other hand, pure zeolite-encapsulated metal NP catalysts will be powerless to catalyse the reactions with big molecular volume (Ju et al., 2015; Dai et al., 2016; Kosinov et al., 2018). The difficulty for mass transfer and bulky molecule catalysis can be significantly improved via producing mesopores inside conventional zeolites as abovementioned. When immobilized on mesoporous zeolites, metal NPs can be not only distributed inside micropores, but also dispersed on the external surface of the mesopores. Those provide more catalytically active sites and easier ways to across the catalysts.

Han et al. (2018) and Kim et al. (2014) successfully prepared mesoporous zeolite nanosponge-encapsulated Co NPs as the efficient catalysts for Fischer-Tropsch (FT) synthesis. In the presence of amphiphilic surfactants containing multiquaternary ammoniums, zeolite nanosponge with ultrathin walls, which could be irregularly interconnected into three-dimensional mesoporous networks, had been synthesized. By means of the incipient wetness impregnation process using [Co(NO3)2·6H2O] as the metallic precursor, Co NPs were immobilized into the mesopores. Due to the narrow distribution of mesopore diameters, Co NPs possessed the uniform NP diameters of 4 nm (**Figures 6A,B**). Compared to bulk MFI zeolite or γ-Al2O3 supports, the cobalt-supporting zeolite nanosponges had higher catalytic performance in FT synthesis (**Figure 6C**). Owing to the strong confinement effect of mesopores, the catalyst exhibited high resistance to sintering, high conversion of CO, and long catalytic lifetime. Meantime, the thin MFI walls contributed to a high selectivity to branched hydrocarbons in the gasoline range (C5-C11). The authors further increased the loading amount of Co within zeolite Beta or MFI nanosponge, resulting in the construction of Co nanowires or networks along the mesoporous nanochannels (**Figures 6D,E**); (Han et al., 2018). Although the mesopores were blocked by Co nanostructures, the accessibility through microporous windows on the mesopore walls could also ensure a high catalytic activity toward FT

synthesis (**Figure 6F**). Furthermore, the zeolite nanospongeencapsulated Co nanostructures also exhibited high catalytic performance in benzene hydrogenation and furfural-to-γvalerolactone conversion.

Based on in situ confinement strategy, mesoporous silicalite-1 nanocrystals-encapsulated Pd NPs (Pd@mnc-S1) could be wellconstructed by Cui et al. (2016) After mixing silica, Pd2+, small amount of polyvinyl pyrrolidone (PVP) and water, silicalite-1 nanocrystals with built-in mesopores were synthesized following the Kirkendall growth process (**Figure 7A**). The presence of PVP could prevent the aggregation of palladium oxide or hydroxide NPs under the high pH value. Metallic Pd NPs were obtained via successive calcination of the as-synthesized products in oxygen and hydrogen atmospheres. Pd NPs with a diameter in the range of 2∼5 nm exhibited a high thermal and chemical stability, indicating the high encapsulation of Pd NPs inside the framework of silicalite-1 nanocrystals. The intrinsic micropores and Pd NPs in Pd@mnc-S1 resulted in good shape selectivity for a series of model reactions, such as hydrogenation, oxidation and C-C coupling reactions (**Figure 7B**). For example, catalytic hydrogenation conversion of nitrobenzene into aniline could be efficiently completed using Pd@mnc-S1 as the catalyst. However, only a negligible amount of 1-nitronaphthalene would be converted into naphthalen-1-amine over Pd@mnc-S1. As well-known, 1-nitronaphthalene possess the molecule size about 7.3 × 6.6 Å, bigger than the micropores of zeolite MFI (5.3 × 5.6 Å), resulting in the inability of Pd@mnc-S1 to catalyse the bulk organic molecules. The same results of shape selectivity had been found in the oxidation and C-C reactions. Additionally, the presence of mesopores inside zeolite supports greatly enhanced the mass-transfer efficiency and thus gave rise to the high catalytic activity. The authors further found that the number of acid sites in mesoporous H-ZSM-5 frameworks could play as "solid ligands" to activate the embedded Pd NPs for organic synthesis (e.g., Suzuki coupling reactions); (Ke et al., 2017).

In order to catalyse the hydrodeoxygenation of the phenolic compounds in bio-oil to alkanes, Wang et al. (2015) immobilized Ru NPs into b-axis-aligned mesoporous ZSM-5 (Ru/HZSM-5- OM) as an efficient catalyst. Owing to the open mesopores and rich exposed acid sites, Ru/HZSM-5-OM exhibited better catalytic activity in the hydrodeoxygenation of bulky 2, 6 dimethoxyphenol (as a model reaction) than bulk or mesoporous

HZSM-5-encapsulated Ru NPs. This is because the bulky reactant molecule could easily transfer through the open mesopores and complete the reaction on the external/mesopore acid sites in HZSM-5-OM. A high conversion rate and cyclohexane selectivity could be reached when using Ru/HZSM-5-OM as the catalyst. Ru/HZSM-5-OM also presented excellent recyclability, keeping the high conversion (95.8%) and cyclohexane selectivity (69.0%) after recycling ten times (**Figure 8**). Due to the fine confinement effect from HZSM-5-OM, Ru NPs could be well-stabilized in zeolitic framework as no Ru species in the liquid had been detected under the reaction conditions or after the recycling test.

Another interesting strategy to build porous zeoliteencapsulated metal NP catalysts is promised to utilize hollow zeolites as the inorganic supports (Dong et al., 2003; Ren et al., 2007; Li et al., 2014b,c; Pagis et al., 2016; Prieto et al., 2016). The structural features endow them into hierarchically macro-microporous or macro-meso-microporous systems, both of which are structurally beneficial for mass transfer of the reactants and products. The active sites of metal NPs on relatively thin shell are easily accessible. For example, Li et al. (2014b) produced Pt NPs inside hollow silicalite-1 crystals as the composite catalysts for size-selective hydrogenation of aromatic compounds (**Figure 9A**). The hydrogenation product

of methylcyclohexane from toluene possessed a kinetic diameter less than MFI micropores and thus easily diffused through the zeolitic shell. Contrastively, the product of trimethylcyclohexane from mesitylene could not go across the zeolite shell to reach the reaction sites on Pd NPs, resulting in the poor catalytic activity of

Pt@Sil-1 on the hydrogenation of bulky mesitylene (**Figure 9B**). Although the catalyst in this case did not perform the activity toward bulky molecules (only size selectivity), it reminds us that, if some mesopores are created on the zeolite shell (hollow mesoporous zeolites), corresponded catalytic performance will be greatly improved in both shape selectivity and bulky molecular catalysis. Owing to the presence of hierarchically macro-meso-micropores in one zeolite support, the composite catalysts after the encapsulation of metal NPs will be most efficient for various organic reactions.

As for the most cases mentioned above, the mesopores in the final catalysts are not particularly ordered, which may lead to the weak mass transfer efficiency compared to the ordered ones. How to construct a set of ordered mesopores inside the zeolites supports for the encapsulation of metal NPs is always a big challenging subject. The catalytic performances (e.g., conversion, shape selectivity, stability) in these expected catalysts should be greatly improved. The catalysts are theoretically employed in many fields of organic or small inorganic molecule reactions.

#### SUMMARY AND PERSPECTIVES

Mesoporous zeolite is a type of ideal inorganic supports for the encapsulation of metal NPs to produce composite catalysts

#### REFERENCES

Aijaz, A., Karkamkar, A., Choi, Y. J., Tsumori, N., Rönnebro, E., Autrey, T., et al. (2012). Immobilizing highly catalytically active Pt nanoparticles inside the for various organic reactions. In this review, the advantages of mesoporous zeolites as the supports are presented in detail with the comparison to other nanomaterials like metal or nonmetallic oxides, carbon, polymers, and porous frameworks. The main strategies for the construction of mesoporous zeolites and the encapsulation methods for metal NPs have been summarized as these two aspects are extremely important to combine mesoporous zeolites and metal NPs into composite catalysts. Several representative examples exhibited well-catalytic performances in reactant conversions, shape selectivities, and stabilities in many organic catalytic reactions. There are still a lot of challenges in the preparation of mesoporous zeoliteencapsulated metal NPs, such as the degree of mesoporous order, the dispersity and size distribution of encapsulated metal NPs, and the large-scale synthesis, and so on. Additionally, the structure-performance relationship between catalytic reactions and catalyst structures (different mesopores, micropores, and metals) should be carefully revealed in the future.

As for the construction of ordered mesoporous zeoliteencapsulated metal NP catalysts, some purposed strategies are presented. The appropriate process should be the onepot synthesis via the mixing of zeolite gel and optimized metallic precursors (Wu et al., 2014) under the presence of some special structural-directing surfactants (Na et al., 2011) or ordered hard templates (Chen et al., 2011). The most metal precursors can be retained after the removal of templates and reduced in situ into NPs inside the micropores. In this way, the ordered mesopores can run throughout zeolite crystals and meanwhile metal NPs are tightly immobilized inside microporous frameworks, resulting in the higher mass transfer efficiency and enhanced catalytic activity, especially for the bulky molecular catalysis. We hope this review can provide useful insights on the synthesis of zeolites supported metal NP composite catalysts and the understanding of the relationship between the catalytic performance and nanostructure of composite catalysts.

#### AUTHOR CONTRIBUTIONS

All authors listed have made a substantial, direct and intellectual contribution to the work, and approved it for publication.

#### ACKNOWLEDGMENTS

This work is supported by National Natural Science Foundation of China (No. 21501095) and Jiangsu Specially Appointed Professor Plan. The authors also thank the supports from Priority Academic Program Development of Jiangsu Higher Education Institutions, National and Local Joint Engineering Research Center of Biomedical Functional Materials.

pores of metal–organic framework: a double solvents approach. J. Am. Chem. Soc. 134, 13926–13929. doi: 10.1021/ja3043905

Alessandro, T. (2002). Catalysis by Ceria and Related Materials. Singapore: World Scientific.

Kärger, J., and Ruthven, D. (1992). Diffusion in Zeolites. New York, NY: John Wiley.

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

The handling editor declared a past supervisory role with one of the authors BL.

Copyright © 2018 Xu, Lv and Liu. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Bandgap Engineering of Indium Phosphide-Based Core/Shell Heterostructures Through Shell Composition and Thickness

Reyhaneh Toufanian1 , Andrei Piryatinski 2 , Andrew H. Mahler 3 , Radhika Iyer 4 , Jennifer A. Hollingsworth4 and Allison M. Dennis 1,3 \*

*1 Division of Materials Science and Engineering, Boston University, Boston, MA, United States, 2 Theoretical Division, Los Alamos National Laboratory, Los Alamos, NM, United States, 3 Department of Biomedical Engineering, Boston University, Boston, MA, United States, 4 Los Alamos National Laboratory, Materials Physics and Applications Division, Center for Integrated Nanotechnologies, Los Alamos, NM, United States*

Edited by:

*Ou Chen, Brown University, United States*

#### Reviewed by:

*Aiwei Tang, Beijing Jiaotong University, China Shuai Chang, Beijing Institute of Technology, China*

> \*Correspondence: *Allison M. Dennis<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Nanoscience, a section of the journal Frontiers in Chemistry*

Received: *01 September 2018* Accepted: *31 October 2018* Published: *20 November 2018*

#### Citation:

*Toufanian R, Piryatinski A, Mahler AH, Iyer R, Hollingsworth JA and Dennis AM (2018) Bandgap Engineering of Indium Phosphide-Based Core/Shell Heterostructures Through Shell Composition and Thickness. Front. Chem. 6:567. doi: 10.3389/fchem.2018.00567* The large bulk bandgap (1.35 eV) and Bohr radius (∼10 nm) of InP semiconductor nanocrystals provides bandgap tunability over a wide spectral range, providing superior color tuning compared to that of CdSe quantum dots. In this paper, the dependence of the bandgap, photoluminescence emission, and exciton radiative lifetime of core/shell quantum dot heterostructures has been investigated using colloidal InP core nanocrystals with multiple diameters (1.5, 2.5, and 3.7 nm). The shell thickness and composition dependence of the bandgap for type-I and type-II heterostructures was observed by coating the InP core with ZnS, ZnSe, CdS, or CdSe through one to ten iterations of a successive ion layer adsorption and reaction (SILAR)-based shell deposition. The empirical results are compared to bandgap energy predictions made with effective mass modeling. Photoluminescence emission colors have been successfully tuned throughout the visible and into the near infrared (NIR) wavelength ranges for type-I and type-II heterostructures, respectively. Based on sizing data from transmission electron microscopy (TEM), it is observed that at the same particle diameter, average radiative lifetimes can differ as much as 20-fold across different shell compositions due to the relative positions of valence and conduction bands. In this direct comparison of InP/ZnS, InP/ZnSe, InP/CdS, and InP/CdSe core/shell heterostructures, we clearly delineate the impact of core size, shell composition, and shell thickness on the resulting optical properties. Specifically, Zn-based shells yield type-I structures that are color tuned through core size, while the Cd-based shells yield type-II particles that emit in the NIR regardless of the starting core size if several layers of CdS(e) have been successfully deposited. Particles with thicker CdS(e) shells exhibit longer photoluminescence lifetimes, while little shell-thickness dependence is observed for the Zn-based shells. Taken together, these InP-based heterostructures demonstrate the extent to which we are able to precisely tailor the material properties of core/shell particles using core/shell dimensions and composition as variables.

Keywords: quantum dots, bandgap tunability, optoelectronic properties, photoluminescence, type-I quantum dot, type-II quantum dot, successive ion layer adsorption and reaction (SILAR)

# INTRODUCTION

Semiconductor quantum dots (QDs) are used to study the fundamental photophysics of quantum confined systems, as well as, for numerous applications in photovoltaics, solid state lighting, biomedical imaging, and biosensing (Dennis et al., 2012b; Chuang et al., 2014; Lan et al., 2014; Ten Cate et al., 2015; Kagan et al., 2016; Hong et al., 2017; Chandran et al., 2018; Kong L. et al., 2018; McHugh et al., 2018). Quantum confinement is easily observed in simple, single semiconductor nanocrystals (NCs): smaller cores are more quantum confined, resulting in an increase in the energy gap between the conduction and valence bands as observed through higher energy (bluer) photon emission following photoexcitation. Core/shell heterostructures exhibit more complex size and compositional dependencies as the bandgap of the particle depends on both the position and relative offsets of the conduction and valence bands of the separate components (Murray et al., 1993; Alivisatos, 1996; Peng and Peng, 2001; Reiss et al., 2009). This added complexity enables significant tuning of the optoelectronic properties of the NCs well beyond what would be feasible with the individual components in isolation. Thus, the core/shell heterostructure can be designed to exhibit photoluminescent (PL) energies and lifetimes that are not achievable with discrete NCs of either the core or shell material alone. We explore the extent of this heterostructure tunability using indium phosphide cores and four different shell materials, zinc sulfide, zinc selenide, cadmium sulfide, and cadmium selenide.

Indium phosphide QDs have garnered significant interest as a cadmium-free option that is tunable throughout the visible wavelength regime (Micic et al., 1994; Guzelian et al., 1996; Battaglia and Peng, 2002; Xie et al., 2007; Li and Reiss, 2008; Xu et al., 2008; Tamang et al., 2016). As-synthesized InP QDs, however, are non-emissive due to non-radiative relaxation of excitation generated charge carriers facilitated by trapping at surface defects (Dennis et al., 2012a; Omogo et al., 2013; Balan et al., 2017; Reid et al., 2018). Fabrication of emissive cores can be performed through two main synthetic routes: etching the surface of the nanoparticle using a fluoride, or overcoating with a larger bandgap semiconductor to create a core/shell heterostructure (Danek et al., 1996; Dabbousi et al., 1997; Talapin et al., 2002; Adam et al., 2005; Chen et al., 2013; Mordvinova et al., 2014, 2017). The latter approach renders QDs emissive through passivation of surface trap states at the core/shell interface, thus inhibiting non-radiative recombination pathways while protecting the core nanocrystal against photooxidation (Battaglia et al., 2003). The bandgap of the shell material, as well as the position of its conduction and valence bands relative to those of the core NC, dictates the bandgap structure of the resulting heterostructure and, thereby, the resulting optoelectronic properties (Li and Wang, 2004; Wang et al., 2018).

In traditional type-I QDs, a low bandgap core material is coated with a second semiconductor with a wider bandgap than the core material with higher energy valence and conduction band offsets (Chandrasekaran et al., 2017; Brodu et al., 2018). Sufficiently large core size allows both the electron and hole to reside within the core, constituting the so-called type-I localization regime. A decrease in the core radius can eventually allow one charge carrier to delocalize into the shell while the other stays localized, especially if the band offset for either the conduction or the valence bands is small. This situation is distinguished as the quasi-type-II localization regime (Kaledin et al., 2018; Kong D. et al., 2018). In the prototypical example of CdSe/CdS, the core/shell conduction band offset is small, such that the core/shell structure exhibits a quasi-type-II localization regime with the electron wavefunction spreading from the core into the shell material, while the wavefunction of the heavy hole is still confined to the core (Kaledin et al., 2018; Kong D. et al., 2018). In cases where either the conduction or valance band of the shell lies within the bandgap of the core, a type-II localization regime is observed, whereby the electron and hole carriers are spatially separated into the core and shell (Dennis et al., 2012a; Acharya et al., 2015). As each carrier relaxes to its lowest energy state prior to radiative recombination, the resulting bandgap of the heterostructure is smaller than the bulk bandgap of either the core or shell material (**Figure 1**). Generally, the size and composition of a core/shell heterostructure dictates different localization regimes and subsequently determines photophysical properties such as the emission energy/wavelength/color and the photoluminescent (PL) lifetime (Reid et al., 2018).

InP QDs are most frequently shelled with ZnS to provide a protective, passivating layer yielding PL (Tessier et al., 2015). The high lattice mismatch between InP and ZnS (7.8%), however, can result in the introduction of interfacial stresses at the core/shell barrier, leading to the formation of defects and non-radiative

recombination pathways. A shell of ZnSe offers a significantly reduced lattice mismatch and still induces PL, but is itself prone to photo-oxidation, making the final material prone to degradation over time. The current preferred solution is to use a multilayered or alloyed shell of ZnSe then ZnS to satisfy the requirements of low lattice mismatch at the core/shell interface and high stability at the shell/media interface (Ippen et al., 2012; Lim et al., 2013). Shelling with other semiconductor materials changes the band alignment and more dramatically impacts the particle photophysics (Htoon et al., 2010; Vela et al., 2010). For example, we produced type-II InP/CdS QDs (Dennis et al., 2012a) that have been shown by us and others to exhibit shellthickness-dependent suppressed-blinking behavior in the NIR (Dennis et al., 2012a; Smith et al., 2017).

In this work, we address the role of core size, shell composition, and shell thickness on tuning the PL emission colors and radiative lifetimes of type-I and type-II systems with an InP core. A range of visible and NIR emitters have been produced by synthesizing core/shell QDs using three different sizes of core InP nanocrystals (small, medium, and large), four different shell compositions (ZnS, ZnSe, CdS, and CdSe), and varying the number of shell deposition iterations (1 through 10). A schematic of possible band alignments for each type of heterostructure is shown in **Figure 1**. In addition, we show the relative bulk conduction and valance band positions for each of the semiconductors used in this study. We demonstrate that in our cadmium-free type-I systems, the emission color is tuned throughout the visible wavelength range through variation of the core size alone, where the shell material acts as an optically inactive passivating layer, protecting the core NC. In our chargeseparated type-II systems, the core size is less influential, while the shell thickness strongly impacts the emission peak wavelength and radiative lifetimes of the resulting NIR emitters.

#### MATERIALS AND METHODS

#### Materials

Oleylamine (OAm, 80–90%), 1-octadecene (ODE, 90%), diphenylphosphine (DPP, 98%), sulfur (S, 99.999%), tris(trimethylsilyl) phosphine (TMS3P, 98%), cadmium oxide (CdO, 99.95%) and hexanes (99%) were purchased from Fisher Scientific. Indium (III) acetate [In(Ac)3, 99.99%], oleic acid (OA, 90%), n-trioctylphosphine (TOP, 97%), selenium pellets (Se, 99.99%), zinc acetate [Zn(Ac)2, 99.0%], dioctylamine (98%), and ethanol (99.8%) were purchased from Sigma-Aldrich. ODE was heated to 120◦C under vacuum for several hours prior to use and stored in an argon glovebox. All other chemicals were used directly without further purification. Quartz glass cuvettes were purchased from Starna Cells Inc. All air-sensitive materials were stored and handled in a glovebox under argon.

#### Precursors

The metal and chalcogenide precursors were prepared with minor modifications to previously reported protocols. A 0.2 M stock solution of indium oleate with a molar ratio of 1:3 was prepared by mixing 3.2 mmol indium (III) acetate (0.93 g) and 10.61 mmol (3.35 mL) oleic acid and heating the reaction flask to 120◦C for 2 h with multiple argon backfill and evacuation cycles to ensure complete removal of acetic acid. After a clear and colorless solution was obtained, 12.58 mL degassed ODE was added to the solution. Once the temperature of the solution reached 120◦C, it was cooled to 90◦C and vacuumed for another hour. A 0.2 M stock solution of cadmium oleate (Cd:OA 1:4) was prepared by mixing 7 mmol CdO (0.9 g) and 31.15 mmol (9.83 mL) oleic acid. The rust colored solution was degassed at 110◦C for 1 h, backfilled with argon, and heated to 280◦C until a clear and colorless solution was obtained. This solution was cooled to 100◦C and vacuumed for an hour to remove any water produced during the cadmium oleate formation reaction. The flask was backfilled with argon and stored at 80◦C for future use. A 0.2 M stock solution of zinc oleate (Zn:OA 1:4) was prepared by mixing 20 mmol zinc acetate (4.39 g) with 88.88 mmol (28.05 mL) oleic acid and degassing under vacuum at room temperature until no vigorous bubbling was observed. The temperature was slowly raised to 120◦C and maintained for 2 h, until the solution was clear and no bubbling was observed. At this point, 71.95 mL degassed ODE was added to the solution. The flask was stored at 60◦C for further use. Solutions of sulfur in both TOP and ODE were used as sulfur precursors. For the preparation of a 0.2 M stock solution of sulfur in ODE, 3.11 mmol (0.1 g) elemental sulfur and 15.59 mL ODE was mixed and heated to 80◦C under vacuum for 2 h. The flask was backfilled with argon and heated at 80◦C overnight to ensure complete dissolution of sulfur in ODE, and stored at the same temperature for future use. A similar procedure was followed for the preparation of the 0.2 M stock solutions of sulfur in TOP and selenium in TOP. TOP:S and TOP:Se were stored at room temperature in an argon glovebox for future use.

Diphenylphosphine selenide (DPP:Se) was synthesized as previously described (Evans et al., 2010). In an argon-filled glovebox, 20 mmol (3.48 mL) DPP, 20 mmol (1.58 g) selenium, and 25 mL anhydrous toluene were loaded in a 3-necked round bottom flask attached to a Findenser (Radleys, UK). The flask was heated at 110◦C for 16 h. Following the removal of toluene in a rotary evaporator, white crystalline precipitates of DPP:Se were obtained. The collected precipitate was washed with cooled toluene to remove impurities. A 0.2 M stock solution of DPP:Se was prepared by mixing 8 mmol (2.12 g) of this powder and 39.98 mL ODE and heating to 100◦C until the crystals completely dissolved in ODE. The solution was kept at 80◦C to maintain solubility.

#### Synthesis of InP Core Nanocrystals

InP core nanocrystals of three different sizes–small, medium, and large (S, M, and L, respectively) were prepared using hot injection synthesis. For the synthesis of the small and medium cores, 2 mL 0.2 M In(OA)3, 1.5 mL OAm, and 2 mL ODE were loaded into a 100 mL round bottom flask. The indium oleate solutions were heated under vacuum to 120◦C and backfilled with argon multiple times, while syringes of 0.1 M and 1 M (TMS)3P in ODE were prepared in the glovebox. An initial 0.2 mL of 1 M (TMS)3P was rapidly injected into each flask of indium oleate at 191◦C with the thermocouple temperature set to 178◦C for the growth phase of the reaction. Starting within 45 s of the first injection, 2 mL of 0.1 M (TMS)3P was added to the flask via dropwise injection. Small cores were reliably produced by quenching this reaction after 3.5 min by removing the heating mantle and adding 12 mL of degassed, room temperature ODE. To generate medium cores, more indium and phosphorus was repeatedly added to the InP cores to provide precursors for growth with a total of 6 mL of warm 0.2 M In(OA)3 and 12 mL of 0.1 M (TMS)3P added dropwise in 4 separate aliquots each over the span of 20 min. After 32.5 min of total reaction time, the core growth was quenched by removing the heating mantle and adding 24 mL of degassed, room temperature ODE. The large InP cores were produced similarly to the small core protocol described above, except that the OAm was replaced by 0.5 mL dioctylamine and the reaction proceeded at 178◦C for 14 h.

For the synthesis of InP/ZnS and InP/ZnSe, core InP QDs were used in their original reaction solution. For the synthesis of InP/CdS and InP/CdSe, core nanocrystals were cleaned by a precipitation and resuspension procedure. To clean the reaction product, aliquots were diluted in hexane and centrifuged to remove any precipitates in an argon glovebox. In a clean centrifuge tube, ethanol was added to the supernatant to precipitate the QDs, and the sample was centrifuged again. At this step, the clear supernatant was discarded and the colored precipitate was resuspended in hexane. A final centrifugation step removed any residual precipitates. The transparent, but colored, supernatant was used for the synthesis of InP/CdS and InP/CdSe. The same procedure was performed for cleaning the core/shell reaction products for further characterization.

#### Synthesis of Core/Shell Nanocrystals

Four distinct shell compositions comprising CdSe, CdS, ZnSe, and ZnS were deposited on the InP core nanocrystals using a modified SILAR technique, adapted for colloidal synthesis (Li et al., 2003). In this approach, aliquots of cationic and anionic precursors were added to the core solution between long anneals at growth temperatures (150–240◦C). The required amounts of anion and cation precursor were calculated based on the number of surface atoms needed to add a monolayer to a certain size of a core/shell nanocrystal. For each reaction, 1 mL OAm and enough ODE to make up a total reaction volume of 10 mL were mixed in a 100 mL round bottom flask on a Schlenk line. The flask was vacuumed at room temperature for 30 min before the solution temperature was raised to 120◦C for multiple vacuum/argon cycles. To protect the InP cores from Ostwald ripening caused by sudden exposure to heat, the flask was cooled to 60◦C for the addition of 200 nmol of InP cores.

The flasks were kept under dynamic vacuum for an additional 15–20 min before the reaction temperature was raised. In addition, the reaction flask for InP/ZnSe was spiked with 0.21 mL (1.207 mmol) of DPP at this time. For all shell compositions, the cation precursor for the first deposition was injected at 150◦C; the first anion injection proceeded 10 min later at the same temperature. Immediately after the anion injection, the temperature was raised to 240◦C and the solution was annealed as described in **Table 1**. For InP/CdS and InP/CdSe heterostructures, all subsequent shell precursors were injected dropwise at 240◦C (Dennis et al., 2012a). However, for InP/ZnS TABLE 1 | Summary of shell reaction parameters.

and InP/ZnSe heterostructures, a temperature cycling procedure was followed as previously described (Xie and Peng, 2009). Specifically, after the anion annealing at 240◦C, the temperature of the flask was lowered to 150◦C for the subsequent cation and anion injections before the temperature was again increased to 240◦C for a longer anneal. To monitor the progress of the reaction, aliquots of the reaction solutions were taken at regular intervals. The reaction was terminated by cooling the solution to room temperature.

#### QD Characterization

Absorption spectra were obtained using a NanoDrop 2000c spectrophotometer, using the pedestal. Samples were cleaned by precipitation and resuspended in hexane or TCE (to eliminate hexane reabsorption in NIR measurements) for PL measurements and measured in quartz cuvettes. Photoluminescence spectroscopy below 900 nm was measured with a Horiba NanoLog spectrofluorometer (Horiba Jobin Yvon) outfitted with a 450 W Xenon arc lamp, double excitation and double emission monochromators, and silicon CCD detector. NIR emission peaks were measured with a similar instrument outfitted with a liquid nitrogen-cooled InGaAs CCD array. Spectra were taken with 400 nm excitation and a 450 nm, 550 nm, or 800 nm longpass emission filter, depending on the wavelength region being examined. The collected spectra were corrected for lamp power and detector sensitivity using instrument-specific correction files.

Absolute quantum yield values were obtained using the spectrofluorometer described above in conjunction with a Quanta-phi six-inch integrating sphere (HORIBA Jobin Yvon). Cuvettes were loaded with hexanes in order to generate a blank reference following excitation at 400 nm. QD samples were diluted directly into the blanked cuvettes and measured with identical conditions. The absolute quantum yield (i.e., the photons emitted divided by the photons absorbed) was determined using the Horiba software.

Transmission electron microscopy (TEM) images were acquired using a JEOL 2100 LaB6 high-resolution microscope operating at 200 kV. Each sample was precipitated and resuspended in hexane twice and drop-cast on a copper TEM grid, which was washed successively by hexane, ethanol, and DI water. The grid was then gently heated to evaporate any residual water. TEM images were sized manually using ImageJ. TEM images have been size matched for presentation and the reported sizing information is based on 108–346 QDs, depending on the sample. The shell thickness was calculated by taking the difference between the core/shell radius determined via TEM and the core radius determined using the 1S absorption peak position (Xie et al., 2009).

Photoluminescence lifetime measurements were performed using an Edinburgh Instruments DeltaFlex time-correlated single photon counting (TCSPC) fluorescence lifetime instrument. QD samples were suspended in hexanes and loaded into quartz cuvettes. Samples were excited with a 405 nm EPL picosecond pulsed diode laser using a 1 microsecond pulse period. For each sample, the measurement continued until 10,000 photons were collected in the highest-intensity bin to ensure accurate fits. Photons were collected at the emission peak with a bandwidth window of 15–50 nm, depending on the FWHM of the sample. After collection, data was analyzed with F980 software (Edinburgh Instruments) by fitting the tail to the triexponential equation R (t) = A1e −t/τ1 + A2e −t/τ2 + A3e −t/τ3 , where An is the percent weight of each exponential. The average photoluminescence lifetime, τ , was calculated using the following equation (Lakowicz, 2006; Jones and Scholes, 2010):

$$\overline{\tau} = \sum\_{n} \frac{A\_n \tau\_n^2}{\sum\_m A\_m \tau\_m}.$$

#### Effective Mass Model of Core/Shell Heterostructures

The electronic structures of the heterostructures were calculated using a two-band effective mass model applied to each core/shell composition as has been previously described (Efros et al., 2000; Piryatinski et al., 2007; Dennis et al., 2012a). The confinement potential has been parametrized using values for the bulk bandgaps, conduction band energies, electron and hole effective masses, and material dielectric constants listed in **Table S1**. The 1S exciton energy is evaluated as a sum of the electron and hole 1S state energies due to the confinement potential (kinetic energies) plus their Coulomb binding energy (perturbative corrections).

#### RESULTS AND DISCUSSION

In order to examine the impact of the composition and dimensions of InP-based core/shell heterostructures on their optical properties, we used an effective mass model to generate wavefunctions for the electron and hole carriers as a function of the core and shell compositions and dimensions and, from the overlap integral, predict the bandgap energy of the heterostructure (**Figure 2**). The modeling data predict that InP/ZnS heterostructures with core diameters larger than 2 nm are exclusively type-I emitters with both the electron and hole localized in the core. The vertical bands of equal energy show that the final emission color is primarily dictated by the starting core size and uninfluenced by ZnS shell thickness. The InP/ZnSe data similarly exhibit vertical bands in the bandgap energy plot after a shell ∼1 nm thick is deposited, but a quasi-type-II structure is predicted for core sizes up to 4 nm in diameter—a range that is more relevant for this study than the very small core sizes predicted to exhibit quasi-type-II behavior for the InP/ZnS system. In this context, we would expect the core size to primarily dictate the color of both InP/ZnS and InP/ZnSe heterostructures. The InP/ZnSe particles, however, are predicted to exhibit more red-shifting with the deposition of thin shells than InP/ZnS. After deposition of several atomic monolayers, neither material is expected to show photophysical changes as a function of shell thickness.

In contrast, modeling of the InP/CdS(e) systems predicts a strong influence of both core size and shell thickness on the resulting bandgap. In these systems, larger cores naturally exhibit redder emission, but the bandgap energy also shifts quickly and significantly with CdS(e) shelling. The hole wavefunction is shown to localize in the core in all cases where the core is >1 nm in diameter. The electron wavefunction on the other hand indicates localization of the electron in the shell as soon as the shell is >1 nm (1.5 nm) for InP/CdS(e). Because of these clear localization patterns with such small core sizes and relatively thin shells, nearly all of the InP/CdS(e) parameter space is classified as falling into a type-II localization regime. One of the hallmarks of a type-II system is that it can exhibit radiative recombination at photon energies lower than is possible for either the core or shell material in isolation. Although not all of the bandgaps predicted for the InP/CdS(e) core/shell structures in the type-II regime are lower energy than the bulk bandgap of InP at 1.35 eV, these composition and dimension combinations are still type-II, as they exhibit lower energy bandgaps than would be possible for either InP or CdS(e) of the same size (i.e., with confinement).

It is well established that semiconductor nanocrystals smaller than their exciton Bohr radius are quantum confined and exhibit size-dependent optoelectronic properties (Murray et al., 1993; Battaglia and Peng, 2002). This was seen for the InP core particles generated for this study as well. InP cores were nucleated by injecting the highly reactive phosphorus precursor, tris(trimethyl silyl)phosphine, into an indium oleate solution heated to 191◦C. Oleylamine facilitated the fast, uniform growth of the smallest cores; additional indium and phosphorus precursors were added at the growth temperature (178◦C) to yield the medium cores. The large cores were synthesized overnight at 178◦C in the presence of a secondary amine (di-n-octylamine) without further precursor addition. The resulting cores exhibit exciton 1S peaks (small and medium) or shoulders (large) at 445, 543, and 622 nm, which correlate to particle diameters of 1.5, 2.5, and 3.7 nm, respectively (**Figure S1**) (Xie et al., 2009).

To observe the effect of shell composition and thickness on the optical properties of InP, epitaxial shells of four different semiconductor materials (ZnS, ZnSe, CdS, and CdSe) were grown on each of three sizes of InP cores with shell growth confirmed through TEM imaging (**Figure 3**). Lattice fringes are evident in the images, demonstrating the crystallinity of the nanoparticles. The core sizes determined from the 1S peak position and the average particle diameters obtained from TEM sizing data (**Figure S2**; **Table S2**) were used to determine the average shell

thickness for each sample, assuming spherical core and core/shell particles. The number of shell monolayers (MLs) deposited was determined by dividing the shell thickness in nm by the thickness of one cation/anion layer (i.e., half the lattice constant of the shell material; **Table S1**).

It is apparent from an examination of the sizing data (**Figure S2**; **Table S2**) that the ZnS shell deposition proceeded significantly less efficiently than the deposition of the other shell materials. While the ZnS monolayer is the thinnest of the four materials with a lattice constant of 5.41 Å (**Table S1**), this difference in unit cell size does not account for all of the difference in the final size. The ZnS and ZnSe shells are only 4.1 and 6.5 MLs thick, respectively, after 10 iterative rounds of SILAR deposition. For the cadmium-free systems, the shell growth procedure is not as effective as that of the cadmium-containing systems for which the chemistry is very well developed. This is mainly attributed to the less reactive nature of the zinc precursor compared to the cadmium precursors. The impact of this lower reactivity can be observed in the PL spectra as well, particularly for some InP/1ZnS(e) samples (**Figure 4**; **Figure S3**), which exhibit low energy tails indicative of trap emission. This suggests incomplete formation of the ZnS(e) shell in the early stages of the reaction. This tail disappears and the full width at half maximum (FWHM) of the PL peak decreases with subsequent shelling (**Figure S4**), providing evidence for complete shell formation with successive precursor addition. We suspect that a combination of high lattice strain, due to the large lattice mismatch between InP and ZnS (7.8%), and the lower reactivity of the zinc precursor combines to produce the least efficient shell deposition in the case of ZnS. Several reports describe increasing the reactivity of zinc by means of the addition of phosphorous containing compounds, with the aim of the formation of more reactive

FIGURE 3 | Size-matched TEM images of InP/ZnS, InP/ZnSe, InP/CdS, and InP/CdSe core-shell heterostructures after 1, 4, 7, and 10 successive shell depositions. The scale bar for all images is 10 nm. The average ± standard deviation diameter of the core/shell particles is included in the inset text in the corresponding image. The size statistics are based on *n* = 108–346 particles, depending on the sample.

zinc phosphine complexes, particularly for the synthesis of thick-shelled InP/ZnSe heterostructures (Joo et al., 2009; Evans et al., 2010; Yu et al., 2011).

To improve the efficiency of the ZnS(e) shell depositions, several Se and S precursors were tested for effective shell growth with zinc oleate. Both trioctylphosphine sulfide (TOP:S) and sulfur dissolved in the non-coordinating solvent octadecene (S/ODE) were tested for ZnS deposition. The addition of sulfur caused significant etching of the core nanocrystals at elevated temperatures, especially in the first stages of the synthesis, as seen through the blue-shifting of the 1S exciton peaks (**Figure S5**). Using TOP:S, however, the 1S peak remained stable, indicating that the core nanocrystal remained intact. The selenium precursors tested include trioctylphosphine selenide (TOP:Se) and diphenylphosphine selenide (DPP:Se). When DPP:Se was used, TEM images indicated QDs of a larger average diameter and a more uniform morphology compared to TOP:Se (**Figure S6**). This is likely due to the formation of the more reactive zinc phosphine monomers, resulting in a higher reaction yield and a more effective shell growth with increasing shell thickness (Yu et al., 2011). We found that spiking the reaction flask with 0.21 mL (1.207 mmol) of DPP in addition to using

DPP:Se as the anionic precursor improved the overall quality of the InP/ZnSe QDs (**Figure S6**).

For CdS(e) shell formation, successful shell growth was observed via TEM imaging using S/ODE and TOP:Se as anionic precursors, respectively, with cadmium oleate as the cationic precursor. Little optimization of the SILAR protocol was needed for CdS(e) shell growth, as many of the thick-shell examples in the literature focus on cadmium systems (Zhou et al., 2016, 2017; Niu et al., 2017). The primary struggle in adapting these protocols to the InP-based system hinged on deposition of the first shell layer without excessive etching or dissolution of the InP core. This is successfully achieved by lowering the initial shell growth temperature to 150◦C, as we previously reported (Dennis et al., 2012a).

Optical characterization was performed on samples comprising three core sizes and four shell thicknesses for each of the four core/shell heterostructure compositions. This set of forty-eight samples spans a wide range of particle sizes, emission wavelengths, and photoluminescent lifetimes, demonstrating the interplay between the core/shell structure and the photophysical function. By examining the PL peak position, we immediately observe the impact of the different band alignments (**Figures 4**, **5**, **Figure S3**). InP/ZnS and InP/ZnSe heterostructures have shell materials with significantly larger bandgaps than their core material (**Table S1**) and thus exhibit type-I behavior. Both electrons and holes are confined to the core nanocrystal, providing the exciton with protection against oxidation and photobleaching. We can see that the PL peak positions of type-I heterostructures do not red-shift in spite of successful shelling demonstrated by an increase in the average diameter of the samples. The thicker shell deposition is, however, accompanied by a slight decrease in the FWHM of these samples, which demonstrates improved passivation of the core and the absence of electron leakage into the shell. The trap emission evident as a red tail present in some PL spectra indicates incomplete shelling after the deposition of the first ZnS(e) monolayer, but disappears following subsequent shell additions.

Although the modeling results predict stronger confinement for the InP/ZnS system compared to InP/ZnSe, we did not find this to be evident in the PL spectra. During initial experiments, InP/ZnS synthesized using S/ODE exhibited measurably bluer emission than InP/ZnSe synthesized with TOP:Se, but visually observable color shifts during the addition of the sulfur was indicative of core etching (**Figure S5**). Once the sulfur precursor was changed to TOP:S to mitigate etching as discussed above, the difference in the peak PL emission wavelength between the InP/ZnS and InP/ZnSe generated with the same starting cores narrowed. Our modeling predicts that smaller cores will exhibit a more quasi-type-II behavior in the InP/ZnSe system than larger cores, which are predicted to be type-I (**Figure 2**). We would expect this to present as a larger red-shift for InP/ZnSe with smaller core sizes, which we observe in fact for both InP/ZnS and InP/ZnSe systems (**Figure 5**). The red-shift for the smallcore InP/ZnS was significantly larger than expected compared to other reactions. This reaction may have been affected by Ostwald ripening, which is more likely with smaller cores than larger ones, prior to successful deposition of the ZnS shell.

Based on the position of the exciton 1S peak of the cores and taking into consideration the average Stokes shift of 60 nm associated with InP, both InP/ZnS and InP/ZnSe largely exhibit type-I behavior (Mic´ic et al., 1996; Fu and Zunger, 1997 ´ ). Both the EMA model and our empirical data show that for the type-I systems, the starting core size is the determining parameter in tuning the emission color. The vertical contour plot lines for InP/ZnS(e) in **Figure 2** indicate that quantum confinement and the heterostructure bandgap do not change with additional ZnS(e) shell thickness after the initial deposition of a limited shell. Our graphs of InP/ZnS(e) peak position with respect to SILAR iterations in **Figure 5** confirm this behavior empirically as well. Our results confirm what others have observed—namely that InP/ZnS(e) emission wavelengths are tunable throughout the visible wavelength range by altering core size.

In contrast with the type-I heterostructures, deposition of CdS and CdSe on the InP core is accompanied by a significant red-shifting of the PL peak into the NIR wavelength ranges, which plateaus after the 7th round of SILAR deposition. Using CdSe or CdS as the shell material, the PL was successfully tuned throughout the red and into the near-infrared wavelength ranges. Larger core sizes resulted in emission at even longer wavelengths. In type-II systems, the significant influence of shell thickness on emission tuning is due to the separation of charge

carriers: the confinement of holes in the core and electrons in the shell. In both cadmium shell systems, carrier recombination can take place across the core/shell interface at a lower energy than the bandgap of either of the constituent semiconductor materials.

It is important to note that in addition to the expected type-II NIR emission from the InP/CdS(e) structures, thicker shelled (>3 ML) InP/CdSe—but not InP/CdS—also exhibited shell-thickness dependent visible emission (**Figure S7**). In an extensive study that is beyond the scope of this paper (Buck et al., in preparation), we demonstrate that both of these photon energies (visible and NIR) are emitted by a single particle, appearing to originate from the type-II recombination described here, as well as, recombination events where both electron and hole carriers are confined in the CdSe shell. It is further notable that this behavior has not been observed in the InP/CdS system. It is likely that the core etching observed upon addition of S/ODE during the InP/ZnS synthesis also occurs when S/ODE is added for CdS deposition. This remodeling of the InP surface could impact surface states, preventing the visible emission originating in the shell. While we have spectroscopic evidence of the InP/ZnS etching (**Figure S5**, discussed above), we can only surmise that the S/ODE behaves similarly during the InP/CdS synthesis because the red-shifting of the optical features is pronounced even at the earliest depositions.

The contrast between the type-I and type-II systems is pronounced in the radiative lifetimes of the heterostructures, as well as, in the PL energy. **Figure 6** illustrates the PL lifetime for the thickest shell of each heterostructure, as well as, the change in lifetime with shell thickness. The lifetime measurements were taken for the emission range comprising one third of the total emission peak width in order to focus on bandgap emission rather than the redder trap state emission visible in the thin-shelled InP/ZnS(e) samples. The InP/ZnS(e) heterostructures exhibit negligible change in PL lifetime with additional shell deposition, as the carriers are both confined in the InP cores; the radiative recombination occurs relatively quickly and the average lifetime does not exceed 100 ns. The InP/CdS(e) structures exhibit dramatically longer lifetimes than the InP/ZnS(e) structures, consistent with type-II behavior. In addition, InP/CdS(e) QDs exhibit a shell thickness-dependent lengthening of the lifetime as the excited electron spreads through a larger shell volume with PL lifetimes approaching 700 ns. Even at particle diameters between 3 and 6 nm, where the particle size across different heterostructures are similar, the radiative lifetimes can be up to an order of magnitude higher for the InP/CdS(e) systems than for InP/ZnS(e).

To see shell-thickness dependent trends in PL intensity, we plotted the absorbance-normalized PL spectra of the medium core samples for four shell thicknesses of each of the four

shell compositions (**Figure S8**). For InP/ZnS and InP/ZnSe, the brightest samples are obtained after the deposition of 1.96 and 3.81 shell MLs, respectively. We observe a clear increase in brightness as the InP surface is completely passivated, but thicker shells result in decreased PL intensity, presumably due to strain effects and defect formation. It's notable that this decline is observable at thinner ZnS shells than ZnSe shells (4.1 vs. 5.9 MLs, respectively); this difference correlates with the higher lattice mismatch for InP/ZnS vs. InP/ZnSe (7.8 vs. 4.9%). For the InP/CdS(e) samples, the first round of shell deposition passivates the InP core, enabling relatively bright emission from the previously non-emissive core. With subsequent shell deposition, however, the effects of spatial charge separation are observed as a significant decrease in the PL intensity. This decrease in PL intensity trends with the increase in PL lifetime as a result of the type-II nature of the heterostructures.

While the trends within each shell composition group are very consistent, we were also interested in the differences in brightness amongst compositions. Here, it is important to note that the absorbance-normalized PL measurements were taken many months after the particles were synthesized. We qualitatively observed that the InP/ZnSe samples significantly dimmed with time through casually monitoring brightness by eye under UV illumination, an option that is not feasible for the NIR-emitting InP/CdSe samples. The thinner-shelled InP/ZnS samples also lost some of their PL intensity with age: months earlier, both the 1.4 and 2.0 ML InP/ZnS exhibited absolute QYs of 43–45%, whereas at the time of the absorbance-normalized PL measurements, the 2.0 ML sample still exhibited a 43% QY, but the relative emission of the 1.4 ML sample had decreased (**Figure S8**). It is challenging to compare the InP/ZnS(e) samples to the InP/CdS(e) samples because of the different detector used for the NIR PL measurements—i.e., the y-axis scaling between the InP/ZnS(e) and InP/CdS(e) plots in **Figure S8** is quite different. We'll note, however, that we were never able to take absolute QY measurements of even the thin-shelled InP/CdS(e) systems on our instrument because the particles were so dim. The high absorptivity of the InP/CdS(e) samples is helpful for brightness, but the QY measurements require very dilute/low absorbing samples, making it more challenging to get a measurement that we were confident in from the very dim samples. In this context, despite the different y-axes of the InP/ZnS(e) and InP/CdS(e) plots, we are quite confident in saying that the type-II heterostructures were considerably dimmer than the type-I QDs.

# CONCLUSIONS

We shine light on the optical tunability of InP-based core/shell heterostructures through compositional and dimensional modifications, resulting in a head-to-head comparison of type-I and type-II InP QDs. The emission energy of the InP/ZnS(e) type-I core/shell QDs was tuned by altering the core size, while the emission of the InP/CdS(e) type-II counterparts was primarily shell thickness-dependent. We demonstrate that charge separation in type-II systems caused by the valence and conduction band alignments gives rise to strongly shell thicknessdependent increase in average PL lifetimes compared to either thin-shelled InP/CdS(e) or InP/ZnS(e) of any shell thickness. Our systematic study of these InP core/shell heterostructures provides insight into QD design for a variety of biological and optoelectronic applications.

# AUTHOR CONTRIBUTIONS

AD and JH conceived of and designed the project. RT performed the experimental work and data analysis with assistance from AM for sizing of the TEM images. RI performed PL spectroscopy measurements in the NIR wavelength range. AP performed the effective mass modeling of the heterostructures. RT and AD wrote the manuscript.

# FUNDING

This publication was supported by the National Center for Advancing Translational Sciences, National Institutes of Health, through BU-CTSI Grant Number 1KL2TR001411. Its contents are solely the responsibility of the authors and do not necessarily represent the official views of the NIH. Funding sources include a Laboratory Directed Research and Development program of Los Alamos National Laboratory (DR20170001). AD was supported in part through a KL2 Fellowship, BU-CTSI Grant Number 1KL2TR001411. RT was funded in part through a Boston University Dean's Catalyst Award. AM was supported in part through Boston University Undergraduate Research Opportunities Program (UROP). AP, RI, and JH were supported by the Laboratory Directed Research and Development program of Los Alamos National Laboratory under DR20170001. This work was performed in part at the Center for Nanoscale Systems (CNS), a member of the National Nanotechnology Coordinated

#### REFERENCES

Infrastructure Network (NNCI), which is supported by the National Science Foundation under NSF award no. 1541959. CNS is part of Harvard University. The Dennis Lab is part of the Photonics Center, BUNano, the Precision Diagnostics Center and the Neurophotonics Center at Boston University.

#### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2018.00567/full#supplementary-material

the process resulting in high photoluminescence efficiency. J. Phys. Chem. B 106, 12659–12663. doi: 10.1021/jp026380n

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2018 Toufanian, Piryatinski, Mahler, Iyer, Hollingsworth and Dennis. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Colloidal Synthesis of Bulk-Bandgap Lead Selenide Nanocrystals

Thulitha M. Abeywickrama, Asra Hassan and Preston T. Snee\*

*Department of Chemistry, University of Illinois at Chicago, Chicago, IL, United States*

Lead selenide quantum dots (QDs) are low-bandgap IV-VI semiconducting nanomaterials that have been studied for a variety of applications. Their preparation using colloidal methods can create small spherical to larger cubic nanocrystals, with an upper limit of ∼17 nm reported to date. Here we describe methods for preparing cubic PbSe nanocrystals over a 20–40 nm size range using a twostep procedure. Specifically, ∼10 nm PbSe QDs are generated using the rapid injection method, the products from which are overcoated with additional lead and selenium precursors. The use of two lead reagents were studied; lead oleate resulted in a maximum of 20 nm cubes, while more reactive lead hexyldecanoate resulted in much larger nanomaterials with bulk bandgaps. However, PbSe samples prepared with lead hexyldecanoate also contained agglomerates. Special care must be taken when characterizing larger strained nanomaterials with X-ray powder diffraction, for which the Scherrer equation is inadequate. A more rigorous approach using the Williamson–Hall method provides characterizations that are consistent with electron microscopy analysis.

#### Edited by:

*Ou Chen, Brown University, United States*

#### Reviewed by:

*Huaibin Shen, Henan University, China Liangfeng Sun, Bowling Green State University, United States*

> \*Correspondence: *Preston T. Snee<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Nanoscience, a section of the journal Frontiers in Chemistry*

Received: *31 August 2018* Accepted: *30 October 2018* Published: *22 November 2018*

#### Citation:

*Abeywickrama TM, Hassan A and Snee PT (2018) Colloidal Synthesis of Bulk-Bandgap Lead Selenide Nanocrystals. Front. Chem. 6:562. doi: 10.3389/fchem.2018.00562* Keywords: lead selenide, quantum dots, semiconductor, nanomaterials, conductivity

# INTRODUCTION

Quantum dots (QDs) are nanoscale particles that have size-tunable optoelectronic properties that may differ substantially from their bulk counterparts. The materials can be prepared using low-cost colloidal methods, which has attracted attention for their use in electronic devices (Talapin and Murray, 2005; Talapin et al., 2010), photovoltaics (Tang et al., 2011; Etgar et al., 2013), sensors (Snee et al., 2006), and in bioimaging (Alivisatos, 2004). Quantum dots are often classified based on the elemental grouping of their composition, which in turn is associated with specific applications. For example, II–VI's such as CdSe are visible emitters often employed in biological sensing applications and have been incorporated into commercially available displays (i.e., television sets). The III–V's such as InP and InAs are of interest for their near-IR emission, and the IV–VI's such as PbS and PbSe are low bandgap materials that have utility in electronics, photodetectors, and solar cells. Lead selenide has been extensively studied due to the high quality of the nanomaterials that can be prepared (Pietryga et al., 2004; Capek et al., 2015) and the ease of tunability of the electronic structure due to the low bulk bandgap (0.27 eV) and large excitonic Bohr radius (46 nm) (Yanover et al., 2012). There are many procedures to synthesize PbSe QDs, although the most robust utilize the rapid injection method since it can produce QDs with narrow, near-atomically precise size distributions (Capek et al., 2015; Shapiro et al., 2016).

Previously reported methods to prepare PbSe QDs often employ high boiling point organic solvents (diphenyl ether or 1-octadecene) containing stabilizers such as oleic acid, into which lead salt and phosphine selenide precursors are injected. For many years the reaction yields were reported to be very low, which was found to result from the near ubiquitous use of trioctylphospine selenide (TOPSe) as the calchogenide precursor (Evans et al., 2010). Specifically, Evans et al.

**72**

discovered that purified TOPSe is non-reactive with lead and that previous reports on preparing PbSe QDs using this reagent relied on the fortuitous presence of secondary phosphine impurities (Evans et al., 2010). As such, it is common to incorporate diphenylphosphine with TOPSe in a preparation. At present, there are reports of the colloidal synthesis of small magicsized clusters to larger PbSe cubes on the order of 15–17 nm, which imparts bandgaps from the higher energy near infrared to shortwave infrared spectral regions (Shapiro et al., 2016).

There is interest in the synthesis of larger nanomaterials using colloidal techniques. For example, Zamkov and co-workers prepared very large (∼25 nm) nanocrystals by ligand-induced digestive ripening of as-prepared CdSe and CsPbBr3 QDs among others (Razgoniaeva et al., 2018). Nanomaterials of this size have minimal interfacial areas, which enhances the conductivity of films made from these systems (Razgoniaeva et al., 2017). We sought to prepare very large monodisperse PbSe nanocrystals for these reasons and to challenge the present ∼15–17 nm size limit (Pietryga et al., 2004; Shapiro et al., 2016). To this end, we explored the use of lead acid precursors of varying reactivity, which resulted in the development of two protocols to produce cubic PbSe QDs over a 20–40 nm in size range. This was achieved by a two-step process whereby ∼10 nm core QDs were overcoated to obtain much larger cubic QDs. The main factors that dictate the sizes and quality of the products are the growth temperature and the identity of the lead precursor. Special care must be taken when characterizing larger strained crystalline nanomaterials using powder X-ray diffraction techniques, for which the Scherrer equation is inadequate. Last, it was found that the larger sized PbSe samples are significantly more conductive than their smaller counterparts.

#### EXPERIMENTAL

#### Materials

Acetone (99.5%), diphenyl phosphine (98%), hexane (mixture of isomers, 98.5%), 2-hexyldecanoic acid (96%), lead(II) oxide (PbO, 99%), methanol (99.8%), 1-octadecene (technical grade, 90%), oleic acid (90%), oleylamine (>98% primary amine), 2-propanol (99.5%), tetrabutylammonium iodide (98%), triethylamine (99%), and trifluoroacetic anhydride (99%) were purchased from Sigma-Aldrich. Acetonitrile (HPLC Grade) and trifluoroacetic acid were purchased from Fisher Scientific. Trioctylphosphine (97%) was purchased from Strem. Trioctylphosphine selenide (TOPSe) were prepared by mixing Se metal into trioctylphoshpine and stirring overnight to prepare a 1M solution. Diphenylphosphine selenide was prepared according to the procedure of Evans et al. (2010). Oleic acid was recrystallized according to the procedure of Arudi et al. (1983).

#### Preparation of Pb(C18H33O2)2

Lead oleate was synthesized using lead oxide, trifluoroacetic anhydride, and oleic acid as described by Hendricks et al. (2015).

#### Preparation of Pb(C16H31O2)2

Lead hexyldecanoate was synthesized using a modification of the method reported by Hendricks et al. (2015). Lead(II) oxide (5.000 g, 22.4 mmol) and 10 mL acetonitrile were added to a 100 mL round bottom flask maintained in an ice bath. After stirring for 10 min, trifluoroacetic acid (0.34 mL, 4.48 mmol) and trifluoroacetic anhydride (3.20 mL, 22.4 mmol) were added. The solution turned clear after 15 min stirring and was then warmed to room temperature. In another 250 mL round bottom flask, 2-hexyldecanoic acid (12.2 mL, 45.0 mmol) and triethylamine (7.13 mL, 50.5 mmol) were mixed in 90 mL of 2-propanol. The lead trifluoroacetate solution was then added to the 2 hexyldecanoic acid solution. The resulting colorless mixture was heated to reflux for 2 h and then cooled to room temperature. The solution was concentrated under vacuum and the precursor was separated as a semisolid by filtration after cooling to the point where solid materials separated. To remove excess triethylamine, the crude product was heated with an attached distillation apparatus with a temperature probe until a condensate appeared at 89◦C, which is the boiling point of triethylamine. The product was left at this condition for ∼1 h to assure the removal of the amine and was then characterized with NMR (**Figure S1**) and combustion analysis (**Table S1**). Deviations from the expected results are due to residual triethylamine; these data are provided in the supporting information.

#### Preparation of "Core" PbSe QDs

PbSe QDs were synthesized using the hot injection method via two protocols that are derivatives of that reported by Shapiro et al. (2016).

#### Method (1) Stoichiometric Ligand Condition

Into a 50 mL three neck round bottom flask containing 7.5 g 1-octadecene (ODE) solvent, either 0.483 g lead oleate (0.63 mmol) or 0.450 g lead hexyldecanoate (0.63 mmol) was added. The solution was heated under vacuum at 100◦C until the pressure stabilized. The selenium precursor in the injection solution was prepared inside a glovebox by mixing 0.4 g of 1.0 M trioctylphosphine selenide (TOPSe, 0.48 mmol) with 0.15 g of diphenylphosphine (DPP, 0.81 mmol) with 1.2 g of trioctylphosphine (TOP), which was then loaded into a 3 mL syringe. The lead precursor/ODE solution was heated to 260◦C under nitrogen, at which point the heating mantle was removed. The selenium solution was rapidly injected when the solution cooled to 240◦C. This solution was maintained at 230◦C for another 10 min before cooling to room temperature.

#### Method (2) Excess Ligand Condition

Another set of PbSe QDs were prepared following the same procedure above with excess ligand present in the lead precursor/ODE solution. Specifically, 0.551 g of oleic acid (1.95 mmol) was added to the 50 mL three neck round bottom flask when using lead oleate as a precursor. Alternatively, 0.500 g of 2-hexyldecanoic acid (1.95 mmol) was added when using lead hexyldecanoate. The preparation of magic-sized PbSe QDs was attempted using lead hexyldecanoate as above, with the exception that the solution was maintained at room temperature for a 24 h growth period.

An XPS survey spectrum for the processed lead oleate "core" QDs was measured (**Figure S2**). The elemental composition

from the XPS analysis is presented in **Table S2**. These data reveal the presence of lead and selenium as well as some oxide. The Pb:Se ratio is 2:1, which is a common observation from PbSe QDs and likely results from the presence of excess Pb precursors that were not removed upon precipitation/washing of the sample before analysis (Jawaid et al., 2011).

# Overcoating of PbSe QDs

Due to the low quality of "core" PbSe nanomaterials prepared with lead hexyldecanoate, all subsequent manipulations were performed with PbSe QDs from Method 2 with the lead oleate precursor. This method produced the largest, most homogeneous QDs. When overcoating a sample, a portion (typically ½) of a batch of the as-prepared PbSe QDs were transferred into a 20 mL vial. Addition of hexane and increasing amounts of 2-propanol resulted in precipitation, and the QDs were separated by centrifugation. The precipitate was re-dissolved in a hexane/2-propanol mixture and centrifuged again; this step was repeated twice more.

#### Method (1) Overcoating With Pb(C18H33O2)2

Approximately half of a batch of as-prepared "core" QDs were processed as above. The precipitated nanocrystals were added to 20 g ODE containing 0.121 g oleic acid (0.43 mmol) in a 100 mL four neck round bottom flask. The mixture was degassed under the vacuum at 100◦C until the pressure stabilized. Then the solution was heated to either 130, 160, or 190◦C under a nitrogen atmosphere. Meanwhile, two precursor solutions were prepared for injection. One for lead was prepared by dissolving 0.290 g of lead oleate (0.38 mmol) in 3 mL TOP, and the selenium solution was prepared by dissolving 0.100 g diphenylphosphine selenide (DPPSe, 0.38 mmol) in 3 mL TOP. These manipulations were performed in a glovebox. Once the QD/oleic acid/ODE bath reached the target temperature, the lead and selenium precursor solutions were added using a syringe injector at a 3 mL/h rate while maintaining a constant temperature. After complete addition of the precursors, the mixture was maintained at the overcoating temperature for another 5 m before cooling to ambient. Products were stored as-prepared in a glove box until later analysis.

Two additional protocols were investigated to prepare larger materials. Both methods sought to increase the precursor to QD ratio in the overcoating process. One was realized by doubling the precursors' masses; specifically, 0.580 g of lead oleate was (0.75 mmol) was added to 3 mL TOP and 0.200 g of DPPSe was likewise added to 3 mL TOP, which were subsequently injected into a QD/oleic acid/ODE solution at 160◦C. Another method used the original masses of lead and selenium precursors, which were applied to ∼¼ of a batch of "core" PbSe QDs.

#### Method (2) Overcoating With Pb(C16H31O2)2

This procedure is similar to Method 1 except where noted. Approximately half of a batch of "core" QDs prepared via Method 2 using lead oleate were processed using hexane/2-propanol as discussed previously. The solvent consisted of a solution of 0.110 g 2-hexyldecanoic acid (0.43 mmol) in 20 g ODE. The injection solutions consisted of 0.540 g lead hexyldecanoate (0.75 mmol) in 3 mL TOP and 0.200 g DPPSe (0.75 mmol) in 3 mL TOP. The overcoatings were performed at either 130 or 160◦C. Products were stored as-prepared in a glove box until later analysis.

#### Characterization

Samples were characterized by removing an appropriatelysized portion from storage in the glovebox, which was centrifuged to separate the PbSe precipitate. The solid was re-dissolved in hexane and centrifuged again; this step was repeated once more. The solid material was re-dispersed in hexane or tetrachloroethylene and was used for further characterizations. Transmission Electron Microscopy (TEM) analyses were performed using JEOL JEM-3010 microscope operating at 300 KeV. TEM grids were prepared by dropcasting a drop of the processed nanoparticle solution onto a 300 mesh carbon-coated Cu grid from Ted Pella. X-ray diffraction measurements were performed using a D8 Advance ECO Bruker XRD diffractometer with Cu Kα (λ = 1.54056 Å) radiation. The instrument was used in both a low and high-resolution modes. The angle-dependent instrument response function of the XRD diffractometer was determined using NIST SRM 640e, a strainfree Si standard composed of ∼4.1 micron sized particles. X-ray photoelectron spectroscopy (XPS) analysis was performed on a Kratos Axis 165 with Al Kα source operate at 12 kV and 10 mA. Fourier-transform infrared spectroscopy analysis was performed using an Thermo Nicolet 6700 FTIR and a Bruker Tensor 27 FTIR. Samples for FTIR analyses were coated onto a KBr pellet. 1H NMR and 13C NMR spectra were recorded using a Bruker Avance DRX 400 NMR spectrometer, and elemental analysis was performed by Midwest Microlab.

Conductivity measurements were performed on "core" ∼12 nm PbSe QDs prepared with method 2 and larger ∼22 nm PbSe overcoated samples. The processing is based on the method of Ning et al. (2014), whereby 1 mL of growth solution was washed with anhydrous hexane and centrifuged 3×. The samples were redispersed into hexane and incubated with 0.6 mL of a 0.6 M solution of tetrabutylammonium iodide in oleylamine for 15 min in a glove box. The samples were washed with excess dry methanol and then hexane, and were redispersed in 2 mL hexane with sonication before spin coating onto a lithographically prepared Si substrate consisting of two Ti electrodes with a 3µm gap. The samples were dried under vacuum, and then taken out into ambient air and light for measuring current vs. applied voltage using a Kiethley 2450 Sourcemeter, with the voltage swept from −3.00 to +3.00 V in steps of 0.1 V. Five devices

were measured and were all found to be qualitatively similar, with differences attributed to lithographic imperfections and the presence of residual photoresist in the device's electrode gap.

# RESULTS AND DISCUSSION

Large 20–40 nm cubic-shaped PbSe colloidal nanocrystals were synthesized from "core" PbSe QDs that were overcoated with additional lead and selenium precursors. The first consideration toward the development of protocols was given to precursors. For the chalcogenide, there has been considerable effort to study the chemistry of selenium since the report by Evans et al. whom demonstrated the necessity of a secondary phosphine species in the production of lead selenide QDs (Evans et al., 2010). As a result, many groups have employed TOPSe with additional secondary diphenylphosphine to prepare PbSe QDs. We have used the same system for selenium, and as such we sought to study more effective lead-carboxylate precursors to realize larger nanomaterials. In this regard, recently Peng and co-workers reported the use of "entropic" carboxylate fatty acid ligands for quantum dot core and core/shell synthesis (Yang et al., 2016). In their vocabulary, an entropic ligand includes 2-hexyldecanoic acid, which is a non-viscous liquid at room temperature despite its very high molecular weight. The Peng group reported that this ligand enhances the solubility of large QDs, and that the extra solubility aids in the overcoating of CdSe QDs into larger CdSe/CdS core/shell sizes (Zhou et al., 2017; Lai et al., 2018). As a result, we examined the use of lead oleate as well as "entropic" lead hexyldecanoate as precursors for PbSe QDs synthesis and subsequent overcoating. Several other parameters such as temperature and stoichiometries were also investigated.

#### Synthesis of "Core" PbSe QDs