Introduction
Emerging energy and environmental crisis necessitate the development of advanced nanomaterials to provide fertile platforms for field applications of chemical conversion or energy storage, including CO2 reduction,[1–6] hydrogen evolution reaction,[7–12] oxygen evolution reaction,[13–17] N2 fixation,[18–23] chemicals synthesis,[24–28] pollutants degradation,[29,30] solar cells,[31,32] batteries,[33,34] supercapacitors,[35,36] and so on. To the scientific and engineering communities, there has been a surge of interest to integrate unlimited types of nanomaterials into purposeful, organized, and single-entity nanostructures with elaborate functionalities, collective properties, and enhanced stability.[37–40] Over the past decades, the wide selection and combination of isotropic or anisotropic nanomaterials into well-defined heterostructures have become a scientific frontier relevant to the sustainable development of society.[41]
Intricate in the structure and arrangement of nanomaterials, the recent paradigm of heterostructures have paid particular attentions to the design of hierarchically ordered heterostructures, in which the self- and/or directed assembly of nanounits is spatially orderly. Compared with the randomly aggregated heterostructures, the ordered heterostructures can have optimized interfaces and electronic structure with emergent synergistic properties/effects, by shortening the pathway of charge transfer and mass diffusion, promoting the atomic usage, strengthening the interaction of nanounits, and so on.[38] The properties of ordered heterostructures, in principle, will closely correlate to the intrinsic crystal phase, size, shape, lattice orientation of each nanounit in heterostructures,[42] and also highly depend on the spatial distribution and interaction forces between fine nanoparticles (NPs).[43–45] Concomitantly, it has imposed stringent demands for precise synthesis of purposeful ordered heterostructure and unraveling the underlying form-to-function relationship.
Recently, extensive evidence along with some review articles have emphasized the development of synthesis strategies of heterostructures and highlighted their tremendous potential applications, with some of them discussing the type (type I,[46,47] type II,[48–50] Z-scheme,[51,52] and S-scheme[53,54]) and electron-transfer pathways of heterostructures,[55] the dimensions (e.g., 0D–2D,[56,57] 1D–2D,[58,59] and 3D–2D[60,61]) of NPs in heterostructures, and the field applications of heterostructures for water splitting,[39,62,63] photocatalytic,[51,64–67] battery,[68–70] and other electrochemical energy systems,[41] etc. There is a rising number of reports revealing that the ordered architecture can endow the heterostructure with optimal atomic utilization as well as strong capabilities in harvesting of solar light and promoting the separation and transfer of photogenerated charge carriers.[71–73] Despite these advances, investigation of ordered heterostructure is still in its infancy. While crystallization and self-assembly of biomaterials into ordered structures abound in nature,[74,75] the controlled alignment and spatial distribution of multidimensional NPs remain challenging, since they usually require a judicious selection of the types of nanounits with designed lattice topology and precise control of their crystallization and surface modification, since nanounits of different nanomaterials tend to have different synthesis condition, growth mechanism, and kinetics.[76] Therefore, ordered assembly of anisometric NPs (e.g., rods, wires, or nanosheets) is likely to be more difficult, holding rising interest to unlock the precise synthesis of ordered heterostructure and unravel the driving force involving.
Herein, we briefly examined and discussed the recent progress in preparing ordered heterostructures with tailoring geometrical and electronic properties, with a focus on the emerging synthetic methodologies and drawing an image about how building blocks of multidimensional and structural difference can be assembled orderly to strengthen the particle–particle interaction for respective sustainable applications. Finally, challenges and future perspectives that lie ahead for the development of vital strategies to engineer the ordered heterostructures were discussed.
Surfactant-Mediating Assembly of Ordered Heterostructures
Whereas the small NPs of distinct phases, sizes, shapes, and dimensions could be precisely engineered in the wet synthesis chemistry, a fine control over the assembly and spatial organization of nanounits to prepare ordered heterostructure remains highly challenging.[43] To tackle the bottleneck, frequently the surfactants are introduced for tailoring the growth and assembly of NPs taking advantage of interaction forces of hydrophobic–hydrophobic interactions, repulsion, or dipolar interactions. For instance, Zhang et al. showed sodium dodecyl sulfate (SDS) can direct the self-assembly of Au NPs and CdSe NPs in ≈100 nm spherical superparticles, in which both the Au and CdSe NPs have small crystalline size, about 2.8–9.0 and 3.3 nm, respectively (Figure 1a–c).[77] Wang et al. revealed the amphiphilic surfactants (e.g., Tween 20 [neutral], cetrimonium bromide [CTAB, cationic], and SDS [anionic]) can not only guide the ordered arrangement of Au, Ag, Fe3O4, and CeO2 NPs on metal–organic framework (MOF) particles, but also can control their spatial distribution (Figure 1d).[43] Through the selection of the kind of surfactants, the spatial arrangement of Au NPs on MOF substrate can further be engineered precisely, with the Au NPs to be monodispersed on MOF particles or distributed to the central of MOF particles.
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While ordered heterostructures can also be prepared by growth of new phase on pre-synthesized substrates,[62,78] the site-selective nucleation and arrangement of new phase on substrate are indispensable, and normally under the assisting of surfactants, such as hexadecyl trimethyl ammonium (CTAB), polyvinylpyrrolidone (PVP), sodium alginate (SA), ethylenediamine (EDA), sodium 4-dodecylbenzenesulfonate (SDBS),[79] SDS,[80] and so on.[81,82] For instance, Li et al. showed that the PVP can direct the site-selective growth and deposition of Co–ZIF-67 (ZIF: zeolitic imidazolate framework) onto α-Fe2O3 nanoarray, as the PVP chemisorbed on α-Fe2O3 could bring additional coordination interaction between pyrrolidone ring (C=O) of PVP and Co ion and thus exert a strong affinity to Co ions (i.e., the precursor of Co–ZIF-67).[83] Surfactants can further prevent the random aggregation of newly formed NPs, and strengthen the particle–particle interaction for their self-assembly.[84] In a water-in-oil synthesis system to prepare the MnO2/graphene heterostructure, the CTAB endows positive charges to the surface of MnO2, and the electrostatic attraction drives the assembly of MnO2 on negatively charged graphene.[85] In some cases, the surfactants could work as the structure-directing agents and control the morphology (e.g., sphere, wire, rod, sheet, cube, etc.), size, and exposure facet of nanounits.[86] For instance, Zhai et al. showed that CTAB can guide the nanocrystals growth on their favorable faces, leading to the generate of MoS2/NiS heterostructure built of elongated wires or cuboids building blocks.[87]
Sometimes, the surfactants can impair the chemical bond of the substrate, and convert the single-component substrate into ordered heterostructure. For example, Xu et al. reported that the SA and EDA can impair the V–C bonds of V2C nanosheet, and following in situ oxidization of newly exposed metal V sites will create a V2C–VO2 nanoribbon heterostructure rich of heterointerface (Figure 2a–h).[88] Despite advances in wet chemistry, the solution-based synthesis of ordered heterostructures with nanoscale building blocks remains challenging, and it still calls for robust and general approaches to realize a precise control of the size, shape, orientation, dimension, and spatial alignment of nanounits in heterostructure.[43]
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Sacrificial Templating for Assembly of Ordered Heterostructures
Sacrificial templating method can convert solid-state nanomaterials into heterostructure that may inherit the chemical composition, size, shape, and pore structure from templates,[89] thereby creating intriguing heterostructures that can hardly be prepared by wet chemistry.[90] To date, existing templates include metal oxides,[91,92] metal hydroxides,[67,93] metal sulfide,[94,95] MOFs,[83,96–98] carbon materials,[99] polyoxometallate,[100,101] etc., as discussed in the following.
By Calcination
Direct calcination of templates at high temperature can create ordered heterostructures, through the dehydration or decomposition of templates with the removal of H2O molecular or organic species from their lattice structure. Porous MOFs comprising tailorable metal sites linked with organic ligands are emerging class of sacrificial templates for advanced materials design.[102–104] The common way is by decomposition along with oxidation of the metal cluster and organic linkers in the framework, in the process of which the single-atom metal or metal particles (oxides, nitrides, pnictides, chalcogenides, etc.) can be generated and in situ distributed on support (e.g., carbon matrixes) derived from carbonization of MOFs.[105] Through a careful control of the calcination condition (e.g., template or time, and atmosphere), the obtained heterostructures could have open channels and inherit the high porosity from MOFs template.[106] For instance, calcination of Ce/Co–tetrakis(4-carboxyphenyl)porphyrin frameworks (TCPP) MOFs under air flow creates CeO2/Co3O4 heterostructure with nanoflowers-like architecture inheriting from MOF.[107] Owing to the ordered architecture, the CeO2/Co3O4 have abundant active sites for activation of peroxymonosulfate (PMS), creating a large number of sulfate radicals (SO4• −) for the degradation of antibiotics in environment (Figure 3a).[108]
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Layered double hydroxides (LDHs) are distinguished 2D layered materials, being widely deployed in many applications for alleviating the energy and environment crisis.[109–111] As layers of LDHs carry two (or more) metal elements, direct calcination of LDHs can create dehydrated layered double oxides (LDOs) with well-preserved metal component and 2D feature of LDHs.[112] Through annealing of ZnCo–LDH at 400 °C under air atmosphere for 6 h, Xu et al. prepared an oxide heterostructure, in which ≈7 nm thicknesses Co3O4 nanosheets with highly active (112) facets are well dispersed on the (001)-faceted ZnO.[113] Notice that a combination of calcination treatment and ions exchange strategy allows to tune the chemical component of heterostructured LDOs. Our group showed after the dehydration of 2D MgAl LDH to give 2D MgAl LDO, additional CuO can be fixed in 2D MgAl LDO nanosheets, only through ion exchange between Mg in LDO and Cu ions in solution.[114] Hence, by stirring of MgAl LDO in diluted Cu solution at room temperature and following annealing treatment, high-density CuO NPs with a defined and parallel lattice orientation appear and they are well dispersed in matrix of MgO and Al2O3 (Figure 3b–d). The well-defined CuO/MgO/Al2O3 catalyst can activate the persulfate efficiently to generate abundance of SO4• − radicals for scavenging the phenol, with a high rate constant k = 0.335 min−1 that outperforms many other Cu-based catalysts with k = 0.0033–0.147 min−1.
Calcination treatment of materials in absence of O2 normally called a pyrolysis process, and could create heterostructures that have redox-active metal or metal oxides anchored on porous carbon carbonaceous materials.[83,115–120] For instance, Khan et al. showed the pyrolysis of Ni–MOF-74 creates Ni NPs highly dispersed on the carbon support.[121] The derived Ni@C-600 demonstrates superior catalytic performance toward the methanation of CO2 and H2 into CH4, with production rate of 488 mmol g−1 h−1 obtained under UV–visible–IR irradiation. Notice that after ten consecutive cycling tests or more than 12 h of reaction under continuous flow configuration, no particle aggregation or significant loss of activity happens. Zhong et al. showed regulation of pyrolysis parameters (e.g., the temperature, reaction time, and gas atmosphere) can control the structure of MOF-derived heterostructure.[117] In addition to MOFs as template, in situ pyrolysis of polymetallic oxides (e.g., nickel molybdate,[122] iron vanadate,[123] nickel silicate,[124] and nickel cerate[125]) paves another straightforward means to fabricate ordered oxide-based heterostructure with intimate interface. As showed by Liu et al., the gradient pyrolysis of NiMoO4·xH2O nanowire at 550 °C in Ar/H2 (95/5 vol%) atmosphere will produce a MoO2–Ni heterostructure, with the MoO2–Ni nanowires covering on the nickel foam (Figure 4a–d) and the Ni NPs homogenously distributing on MoO2 nanowires.[126] In this process, the pyrolysis temperature is a key to control the pyrolysis process and thus the structure of heterostructure. While it gives Mo–Ni alloys at rising temperature >600 °C, MoO2–Ni heterostructure can be obtained at a relatively low temperature (<600 °C). Analogously, Hao et al. found the in situ pyrolysis of (NH4)6Mo7O24/GO at 400, 500, and 600 °C will give MoO3/rGO-400, MoO3@MoO2/rGO-500 and MoO2/rGO-600 sheets, respectively (Figure 4e,f, GO: graphene oxide, rGO: reduced graphene oxide).[127] With ordered heterostructure of MoO3@MoO2/rGO-500 as example, the lamellar hybrid of MoO3@MoO2 will have width of ≈200 nm and thin thickness of ≈5 nm, and such oxide hybrid can be homogeneously dispersed on thin rGO nanosheets. The thin MoO3@MoO2/rGO heterostructure has intimate interface, ensuring an effective charge transfer for high performance in catalysis and other related applications (e.g., batteries and other energy-storage devices).
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By Reduction Reaction
High-template-reduction treatment can convert templates into heterostructures carrying highly dispersed small metallic NPs.[112,128,129] By H2 reduction of ZnFeAl–LDH at 500 °C, Zhao et al. created a Fe-based heterostructure (denoted as Fe-500), which has ≈5 nm Fe0 and FeOx well dispersed on ZnO and amorphous Al2O3 derived from decomposition of LDHs (Figure 5a,b).[130] Owing to the high dispersion and intimate contact between the metallic Fe0 and FeOx on oxide support, the Fe-500 catalyst is efficient in driving the CO hydrogenation, with the high initial selectivity toward hydrocarbons (89%, with light olefins product reaching 42%). Likewise, Chen et al. found the reduction and annealing of NiAl–LDHs can give heterostructured Ni/H-Al2O3, with about 8.7 ± 2.6 nm Ni NPs well decorated on the Al2O3 nanoflake.[131] Unlike the H2-reduction treatment performed under high temperature, Zhang et al. reported an in situ reduction of ternary CuZnAl–LDH in solution at room temperature using ascorbic acid as reducing agent, to generate uniform dispersion of Cu2O NPs with an average size of 2.22 nm on LDH nanosheets.[132] In this process, regulating the concentration of ascorbic acid and the reaction time can control the density of ultrafine Cu2O (Figure 5c,d). Notice that the catalytic reduction of N2-NH3 conversion for the synthesis of nitrogen-based fertilizers and other commodity chemicals is of great significance, while it is suffered from the extremely high bond energy (941 kJ mol−1) of N≡N.[133–136] The dispersed and ultrafine Cu2O on LDH can act as the robust catalysts to reach efficient visible-light-driven photocatalytic reduction of N2 to NH3, with a generation rate of NH3 (30.31 μmol g−1 h−1) better than many other reported benchmark photocatalysts.
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By Sulfuration and Phosphorization
Introducing heteroatoms (e.g., S and P) to heterostructure by sulfuration and phosphorization treatment can create ordered heterostructure with improved hydrophilicity and electronic conductivity.[137] Over the past decades, metal sulfides (e.g., In2S3, ZnIn2S4, CdIn2S4, etc.) with appropriate bandgap have demonstrated unique electronic and optical properties for photocatalytic redox reactions of H2 evolution,[138,139] CO2 photoreduction,[140,141] and organic photosynthesis.[142,143] In this aspect, sulphuration is a common way to create multifunctional metal sulfides with controlled metal component, size, and shape like the templates.[144] Wang et al. showed that sulfidation of 1D prism-shaped MIL-68 (MIL: Materials Institute Lavoisier) can create 1D In2S3, while following hydrothermal cation exchange will give hierarchical nanotube comprising ordered In2S3–CdIn2S4 (Figure 6a).[145] Likewise, the successive decomposing, thermal treating and vulcanizing of core–shell polyacrylonitrile@MOF-74 nanofibers can create tubular CoO/Co–Cu–S heterostructure—an assembly of CoO and CoS2 nanoneedles (Figure 6b).[146] Owing to synergistic effect of different metallic ions and the high surface area of the hollow structure, the CoO/Co–Cu–S is a good candidate to prepare all-solid-state hybrid supercapacitor with a high specific capacity of 320 mAh g−1 at a current density of 2.0 A g−1 (Figure 6c,d).
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Phosphorization of templates can create heterostructures by assembly of phosphides with tunable crystalline phases, electronic, geometrical structure, and unique performance.[147] Quan et al. reported a two-step procedure to prepare NiFeP@NC/Ni2P loop-sheet heterostructure (Figure 6e).[148] The Ni–Fe Prussian blue analogues (PBA) nanocubes are in situ and directionally grown on the peripheral edge of Ni(OH)2 nanosheets to form Ni–Fe PBA/Ni(OH)2 loop-sheet heterostructure, in which the [Fe(CN)6]3+ coordinated with Ni2+ ions released from H+-etching of Ni(OH)2, and phosphorization treatment of intermediate gave a NiFeP@NC/Ni2P loop-sheet heterostructure. The NiFeP@NC/Ni2P is by assembly of oriented and interconnected Ni2P NPs with a high density of grain boundaries, and is efficient in driving water oxidation with an overpotential of 223 mV at a current density of 20 mA cm−2 and a small Tafel slope of 46.1 mV dec−1, owing to the rich accessible active sites and improved mass/charge transfer of ordered architecture (Figure 6f). Analogously, Li et al. prepared heterostructured Co/CoP embedded within ordered macroporous–mesoporous–microporous carbon via the carbonization and phosphorization of single-crystalline, ordered macro-microporous ZIF-67.[149] The Co/CoP interfaces favors the exposure of active sites and also optimize hydrogen/water absorption free energy via electronic coupling, thus leading to outstanding catalytic activity with overpotentials of only 120 and 260 mV at 10 mA cm−2 for the H2- and O2-evolution reactions in 1.0 m KOH, respectively.
By Chemical Etching
Chemical etching has long been recognized as appealing way to tune the chemical composition, phase, surface, and pore structure of templates.[150] For instance, chemical etching can convert MOF template into heterostructure comprising dispersed 2D LDHs with adjustable layer thickness and tunable component. Zhang et al. reported in presence of urea and nickel nitrate in mixed solution of ethanol and distilled water, etching of spindle-like MIL-88 A comes along with the coprecipitation of LHDs to give a double-shelled Ni–Fe LDH nanocage.[151] The shell number (one to two or to several) of Ni–Fe LDH could be tailored by adjusting the volume ratio of ethanol and distilled water. Of optimal structure, the double-shelled Ni–Fe LDH nanocage delivers high-electrocatalytic oxygen evolution reaction activity in an alkaline electrolyte and shows a current density of 20 mA cm−2 at a low overpotential of 246 mV. To investigate the etching mechanism and growth kinetics of ZnCo–LDH on ZIF-8 by adding Co2+, Wang et al. employed the liquid-phase TEM imaging technology (Figure 7a).[152] They found a prerequisite to obtain a well-defined hollow structure is to reach a balance between the etching rate of MOF and the growth rate of LDH. At an etching rate faster than the growth rate of LDH, the MOF template will collapse and no ordered hollow structure can be obtained.
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When etching of MOF can be controlled to create an MOF@LDH intermediate, a posttreatment of MOF@LDHs by calcination, phosphorization, or sulfurization can expand the family of accessible ordered heterostructured materials.[150] Chen et al. created a hollow H-LDH/Co9S8 nanocage through in situ etching of ZIF-67 with Ni2+, as the Ni2+ can extract the Co2+ ions to create cage ZIF-67 and simultaneously form a NiCo–LDH shell on the cage (Figure 7b).[153] Further thioacetamide-assisting sulfidation reaction will create hollow H-LDH/Co9S8 of intriguing cage architecture. By such well-defined heterostructure, the large cavity of H-LDH/Co9S8 cage could help to inhibit the diffusion of lithium polysulfides and accommodate sulfur in structure, while the intimated interface promotes the electron conductivity and Li+-ion diffusivity (Figure 7c,d). For exploring robust and ultralong life span electrodes, the H-LDH/Co9S8 delivers a high discharge capacity (1339.1 mAh g−1) at the current density of 0.1 C with an ultrastable life span over 1500 cycles.
Chemical etching can create complex nanoarchitecture. Our group have reported that when being etched in K3[Co(CN)6] solution and following annealed, the prism-like MIL-68(In) crystals can be transformed into branched In2O3 microtubes, whose shell is composed of aligned and elongated In2O3 nanowires with length of over 200 nm.[154] The elongated nanowires are by assembly of tiny In2O3 nanorods of a diameter of ≈5 nm and with the same crystallographic orientation, which forms microtube mesocrystal that can diffract the electron analogous to that of the single crystal (Figure 7e–i). Under identical conditions, the direct annealing of pristine MIL-68(In) without a pre-etching treatment only produces microtube composed of randomly aggregated In2O3 NPs. The derived ordered-In2O3–ZnIn2S4 photocatalyst has better catalytic performance over H2 evolution than the disordered In2O3–ZnIn2S4, as the ordered architecture of branched In2O3 can significantly boost the short-range electron transfer in interface of In2O3–ZnIn2S4 heterojunction.
Once a site-selective etching can be achieved, an in situ topological transformation of template may occur, paving additional pathway to create intriguing heterostructure.[155–157] Inspired by the larger coordination constant of [Zn(NH3)x]2+ than [In(NH3)x]3+, our group showed that the NH3·H2O can selectively extract the Zn element from ultrathin half-unit-cell ZnIn2S4 (referred to as HZIS) (Figure 8a). The topological atom extraction of Zn from ultrathin HZIS creates an in-plane 2D In2O3/ZnIn2S4 heterostructure, in which the newly generated In2O3 is confined in 2D plane with atomic coherence to the lattice of residual 2D HZIS.[158] density functional theory (DFT) calculations show the charge-transfer ability of In2O3/HZIS is highly dependent on the nanoscale distance between In2O3 and ZnIn2S4 in heterostructure, and the “zero distance contact” of obtained in-plane In2O3/HZIS 2D heterostructure has the optimal charge-transfer ability when compared with the out-plane heterostructures of larger distance (Figure 8b–f). The 2D in-plane In2O3/ZnIn2S4 heterostructure further induces remarkable charge redistribution at the heterojunction interface and creates local electric field confined within ultrathin layer, driving the visible-light-driven CO2–CO conversion efficiently, which exhibited a CO production rate up to 5624 μmol h−1 g−1, 5.3 times that of pristine HZIS (1063 μmol h−1 g−1).
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Solid-to-Solid Phase Transition to Generate Ordered Heterostructure
In passing decades, most of related reports were based on the crystallization and growth of heterostructure in liquid media. In recent years, principle of solid-to-solid (S–S) phase transition has been developed to prepare advanced heterostructures.[159] For instance, S–S phase transformation can help to customize the phase of transition metal dichalcogenide (TMDs) that have held great promise for electrolysis, supercapacitors, batteries, etc.[94] Given there is a very small difference between the free energy of semiconducting 2 H phase and metallic 1 T′ phase MoTe2, the S–S transformation of MoTe2 normally occur. In light of this, Xu et al. created a coplanar 2D 1 T′–MoTe2/2 H–MoTe2 heterostructure with seamless interface and small contact resistance (Figure 9a–d).[160] They found in preparation of 1 T′–MoTe2 film by tellurizing the Mo film at high temperature, the 1 T′ phase MoTe2 was first generated, while thereafter, a spontaneous atomical rearrangement of Mo and Te atoms in plane of 2D 1 T′ MoTe2 happened. Partial S–S transition of 1 T′ to 2 H MoTe2 then took place and generated a coplanar 1 T′/2 H MoTe2 heterostructure. Chemical agents may provoke the S–S phase transition of intermediate. During the preparation of perovskite, Wang et al. found as the ethylamine iodide (EAI) can interact with the NH2 group of δ-CH(NH2)2PbI3 through hydrogen bond, the S–S transformation of yellow phase δ-CH(NH2)2PbI3 into the highly oriented black phase α-CH(NH2)2PbI3 takes place.[161] The α-CH(NH2)2PbI3 manifests great potential applications in optoelectronics due to its strong light absorption capability and high carrier mobility.
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Compared with 3D bulk-sized materials, the 2D materials of a thin thickness can be more readily to be engaged in the S–S phase transition, due to their high surface energy, large surface-to-volume ratio and the atomically thin layers.[162] Within the 2D materials, the rearrangement of atom and molecular of metastable 2D intermediate can form intriguing 2D heterostructures.[159] He et al. employed the electron injection technique to engineer the phase transition of 2 H MoS2, and fabricated an ordered heterostructure built of conductive TiO that is chemically bonded to 1 T–MoS2 (TiO–1 T–MoS2) nanoflowers.[163] It shows that the electron injection can trigger a reorganization of the Mo 4d orbitals in MoS2 and provoke the phase transition of MoS2 from 2 H to 1 T in 2D plane. Thanks to the higher electronic conductivity and lower Na+ diffusion barrier of 1 T–MoS2 than 2 H–MoS2, the TiO–1 T–MoS2 nanoflower (NF) delivers a high performance on constructing the sodium-ion battery, with a high rate capability (650/288 mAh g−1 at 50/20 000 mA g−1, respectively) and good cyclability (501 mAh g−1 at 1000 mA g−1 after 700 cycles) (Figure 9e–f).
S–S transition of 2D MOF also opens access to prepare atypical hybrid MOF. To this end, researchers proposed a growth, etching, structural transformation (GET) principle to customize the growth and phase transformation of 2D MOF intermediate, which involves of three distinctive processes, including MOF-on-MOF growth (G), etching (E), and structural transformation (T).[164] With each GET progress regulated by controlling the reaction parameters, ordered 2D MOFs heterostructure of distinct architecture could be obtained. Taking advantage of S–S transition of leaf-shaped ZIF-L via GET, Lee et al. produced the ring of ZIF-8@ZIF-67(Co, Zn), ring of ZIF-8@ZIF-8, and plate of ZIF-8@ZIF-67 heterostructure (Figure 10a).[164] With the formation of core–shell ZIF ring as example, the growth of ZIF-67 shell starts from the outside of the MOF particles at the early growth stage, accompanied with a change of the morphology of particles from leaf-like particles to elongated hexagon. As H+ is generated by the hydrolysis of Co2+, etching occurred at the boundary between the newly grown ZIF-67 and the MOF template, creating a large void in the center of MOF template (Figure 10b–j). At prolonging reaction time, the transformation of MOF template to ZIF-8 is accomplished. When the three GET process is regulated, it produces the distinct assembled plates of ZIF-8@ZIF-67. Worthy of note, further pyrolysis of the ordered rings and plates MOFs can give a large family of derivatives comprising metal or metal oxides well distributed on carbon.[165,166]
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S–S phase transition can also convert the bulk-sized intermediate into ultrafine particles of high dispersion and create ordered architecture. Recently, our group showed in preparation of LDHs-related heterostructure, the surfactant-free hydrolysis of Mn(COO)2·2H2O and NiCl2 in MeOH at 180 °C (Ni/Mn=0.3) at the early growth time of 1 h first generates an amorphous intermediate that is composed of Ni and Mn elements.[167] The S–S phase transition of amorphous intermediate occurs at 18 h and spontaneously gives an intriguing 2D/2D heterostructure—an ordered assembly of ultrathin nanosheets of oxygen-deficient Mn3O4 (OVs–Mn3O4) anchoring on ultrathin nanosheets of Ni-rich Ni–Mn LDHs (Figure 11a,b). The well-defined 2D/2D heterostructure manifests both high CO2 reduction reaction (CRR) efficiency and CO selectivity (with VCO reaching up to 16.10 mmol h−1, CO selectivity of 95.0%), owing to the ensemble effect oxygen vacancies and catalytically active nickel sites, and the fast charge transfer in layer-on-layer 2D/2D heterostructure (Figure 11c).
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S–S principle can create ordered heterostructure by assembly of building blocks that have distinct lattice match and parallel crystallographic orientation.[168] Our group have reported a spontaneous S–S phase transition of bulk-sized CuAl–LDHs crystal into ultrafine CuO clusters that are of atomical dispersion on 2D AlOOH with defined orientation (Figure 11d).[169] In this process, a precipitation of CuCl2·2H2O, AlCl3·6H2O, and urea in deionized water at 120 °C first gives the microscale Cu–Al LDH crystals at growth time <6 h (Figure 11e). These bulk Cu-rich LDH crystals are metastable, and further self-decompose into <2 nm CuO clusters that are oriented and atomically dispersed on ≈4 nm thickness AlOOH nanosheets. It reveals the parallel orientation between the (060) plane of AlOOH NSs and the (001) plane of CuO clusters, owing to the close d-spacing between the (002) of AlOOH (dAlOOH(002) = 0.143 nm) and (220) plane of CuO (dCuO(220) = 0.150 nm) (Figure 11f–h). The lattice-match effect also tightly immobilizes the ultrafine CuO clusters on ultrathin AlOOH. With catalytic hydrogenation of 4-nitrophenol as model reaction, the oriented CuO–AlOOH reaches a high reaction rate constant of 130.0 s−1 g−1, significantly outperforming the commercial Pd/C catalysts and reported state-of-the-art noble-metal catalysts (1.89–117.2 s−1 g−1) as shown in Figure 11i.
Lattice-Match-Driven Epitaxial Growth of Oriented Heterostructure
Epitaxial growth can guide the growth and oriented attachment of new component on preformed substrate with a defined lattice direction. The components in system of epitaxial growth normally have a small lattice mismatch,[170–173] so as to minimize the interfacial energy of heterostructure.[45] Over the past decades, solution-based epitaxial growth has been reported to create a number of nano-heterostructures, composed of metal–metal, metal chalcogenide–metal chalcogenide, metal chalcogenide–metal oxide, metal–metal chalcogenide, metal–metal oxide, and carbon–metal chalcogenide hybrids. Zhao et al. revealed the controllable epitaxial growth of secondary MOF on seed MOF to create ordered hybrid MOF with tunable phase and chemical component.[174] Only by adding the starting materials of Zr-BTB (i.e., ZrOCl2·8H2O and dichloroacetic acid dissolved in N,N-dimethylformamide) in a solution containing PCN-134 (PCN: porous coordination network) nanoplates, the nanosheets of Zr-BTB can selectively grow on the six edge planes of PCN-134 nanoplates, spontaneously having their (010) planes parallel to each other. With further regulation of the phase, size and shape of seed MOF (or the secondary MOF) in developed strategy, distinct types of MOF heterostructures composed of lattice-oriented 1D/0D PCN-222/PCN-608, 1D/1D PCN-222/NU-1000, and 1D/2D PCN-222/PCN-134 with tunable band structure and enhanced charge-transfer capability can be obtained.
Many reports evident the adaptive structure variation in the epitaxial growth of hybrid MOFs, in which the secondary MOF may regulate their growth behavior and orientation to have their structure, morphology, and orientation to be fit to those of the seed MOF. Wang et al. investigated the growth behavior of Fe–BDC (BDC is 1,4-benzenedicarboxylic acid) and Fe–NDC (NDC is naphthalene-1,4-dicarboxylic acid) on a diversity of seed Zr-MOFs.[175] They found that the lattice change of seed Zr-MOF can affect the structure and morphology of newly formed Fe–BDC and Fe–NDC. It reveals the morphology of heterostructures can be affected by the degree of lattice difference between the template MOF and the epitaxial MOF. Five topologically identical Zr–MOFs (MOF-801, UIO-66, DUT-52, UIO-67, and BUT-30) with the same face cubic-centered topology but different ligand lengths were set as templates for epitaxial growth of Fe–BDC (Figure 12). Whereas the epitaxial growth of Fe–BDC on UiO-66 generated irregularly shaped particles, the epitaxial growth of Fe-BDC on MOF-801 or DUT-52 results in an octahedron and eight uniform protrusions perpendicular to the faces of the octahedron. Although using MOF-801, UIO-66, and DUT-52 as templates all create octahedral heterostructures, the heterostructures by epitaxial growth of Fe–BDC have different morphological feature. Lee et al. also reported the self-adjustment of lattice of Fe–MIL-88C in the tip-to-middle anisotropic growth of Fe–MIL-88B-on-Fe–MIL-88C.[176] Although Fe–MIL-88C and Fe–MIL-88B with large mismatched cell lattices, the growth of Fe–MIL-88C mostly began at both the ends of the Fe–MIL-88B template and then effectively prolonged in the c direction, and gradually covering the template. These tip-to-middle growth behavior showcases the self-adjustment of cell lattice of MOF and confirms the spontaneous manipulation of the structure and composition in growth of nanomaterials.
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Epitaxial growth of building blocks with an lattice-oriented alignment can also take place between the two components that have distinct structures and topologies.[161] Feng et al. showcased the epitaxial growth of BiOI on ZnO nanowires.[177] By pre-fixing of the precursor of BiI3 on ZnO, an epitaxial growth of 2D BiOI nanoflakes on 1D ZnO nanowires occurs, driven by the similar interplanar spacing of the (002) face of ZnO and (111) face of BiOI. The ordered ZnO/BiOI heterostructure has lattice match at the interface, allowing a fast charge-carrier transfer and providing enormous active sites to reach the high photocatalytic activities over the photodegradation of bisphenol-A (BPA) in water. Das et al. also showed the lattice-match effect can induce the site-selective nucleation and facet-directed epitaxial growth of Pb4S3Br2 on nanostructured CsPbBr3.[178] The structure of final CsPbBr3–Pb4S3Br2 heterostructure highly depends on the shapes and exposure facets of CsPbBr3, as the Pb4S3Br2 tends to grow on the {110} and {200} facets of rhombicuboctahedrons CsPbBr3, while prefers to grow on the {110} facets of hexapods CsPbBr3, and on the {002} facets of dodecahedron of orthorhombic CsPbBr3, respectively (Figure 13a).
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Lattice-match epitaxial growth can drive oriented dispersion of small metallic particles with defined exposure facet on a support. Lv et al. showed that ≈4 nm Pt NPs that are of a high density and evenly dispersed on single-crystal LiTiO2 (Sc-LiTiO2) can be prepared only by the thermal reduction of H2PtCl4 by NaBH4 in aqueous solution containing Sc-LiTiO2.[179] There is identical in lattice spacing (both 0.14 nm) of cubic LiTiO2 and cubic Pt, and the Pt NPs thus have the defined {111} facets well parallel to that of the {111} facet of Sc-LiTiO2, triggering oriented growth and high dispersion of tiny Pt NPs on Sc-LiTiO2. DFT calculation suggests that there will generate strong interaction at the interface of aligned LiTiO2 (111) and Pt (111), with its electrochemical active surface area evaluated to be 80.9 m2 g−1. The Pt NPs/Sc-LiTiO2 manifests a high-electrocatalytic ORR activity and durability in acid solution in reference to that of commercial Pt/C. Under such principle, it is feasible to change the substrate to tune the exposure facets of metallic particles for optimized performance and diverse applications.
Through the epitaxial growth model, organization of nano-sized building block into superlattice nanowires is also feasible. In recent years, the superlattice nanowires have held great promise for optoelectronics and solar-to-fuel conversion,[180,181] as they carry engineered band structures and geometric parameters enabling the customization of the charge-transfer pathway of heterostructure and realization of full utilization of sunlight.[137] Prior researches suggest the epitaxial growth allows the oriented and site selective growth of new phase on host nanowire to create heterostructured superlattice nanowires. As such, Li et al. reported in preparation of CdS–ZnS quantum dots in nanowire that the catalyzed growth (Ag2S as the host catalyst) of nanowires can integrate CdS and ZnS in a superlattice nanowire (Figure 13b,c).[182] The growth of nanowire can be switched from the component of ZnS to CdS, by injecting the Cd precursors that initiates the nucleation of CdS quantum dots (QDs) at the catalyst–nanowire interface. With the depletion of Cd precursor, the epitaxy growth of ZnS restarts. In this procedure, the slow catalyzed growth is a key to realize the precise control of axial composition and creation of the sharp interface of CdS–ZnS superlattice nanowires (Figure 13d,e). On this basis, they further proposed a low-temperature cation-exchange strategy to transform the aforementioned colloidal CdS–ZnS superlattice nanowires,[137] into a big family of axial superlattice nanowires, including MxSy–ZnS (M=Cu, Ag, Co, Ni, Pb, Hg) and CuMS2–ZnS (M=In, Sn) with diverse applications to mitigate the energy and environmental burden. Worthy of note, Lu et al. showcased that utilizing the difference of precursor reactivity can control the sequential growth of ZnS on CdS nanowires to prepare heterostructured CdS–ZnS nanowires by one-step metal–organic chemical vapor deposition (MOCVD) process.[183,184]
Lattice-match epitaxial growth of two components with distinct structures, sometimes, can occur in one-pot synthesis strategy spontaneously. Recently, our group proposed that the inorganic TiO2 nanowires and MOF particles can align with each other to form a lattice-oriented TiO2/MOF mesocrystal, which can diffract the electrons like a single crystal owing to their oriented alignment (Figure 14a,d).[185] Specifically, the hydrothermal growth of Ti(SO4)2, K4Fe(CN)6·3H2O and PVP in distilled water under 80 °C for 24 h spontaneously gave a self-assembly of ultrathin TiO2 nanowires and Prussian blue (PB) particles. The TiO2 and PB shows oriented alignment to each other as directed by a lattice matching of the spacing (0.297 nm) of TiO2 and the dPB(222) spacing (0.292 nm) of PB (Figure 14e). The ordered architecture accelerates the electron transfer from TiO2 to PB, leading to a fast color switching of PB under ultraviolet (UV), visible light, and sunlight. Accordingly, the ordered OA–TiO2/PB (OA: oriented attachment) can be considered as an ideal candidate for the development of efficient photoreversible color-switching system (PCSS) and advanced rewritable paper technologies.
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Vapor Deposition and Van der Waals Assembly
The epitaxial growth in liquid media could be suffered from the uncontrollable and random aggregation of nanoscale building blocks. To date, while the vertical and lateral stacking 2D heterostructure have demonstrated great potential in applications of electronics, photonics, and catalysis, there lacks of robust principles to control the vertical or lateral stacking of 2D materials. To tackle the bottleneck, physical vapor deposition (PVD) strategy could provide a high-level control over the stoichiometry and assembly behavior of 2D nanomaterials for preparing vertical or lateral heterostructure, for example, CsPbBr3/SrTiO3,[186] AlGaN/GaN,[187] WO3–WS2,[188] SnSe/SnSe2,[189] etc.[190] Kim et al. utilized the pulsed laser to induce the PVD growth of WO3/LaAlO3/Nb:SrTiO3 heterostructure, in which three unit cells of atomically thin LaAlO3 deposit on the single-crystal Nb-doped SrTiO3 substrate (Figure 15a–d).[191] Yan et al. also applied molecular beam epitaxy technique to grow an AlGaN/GaN quantum-well heterostructure on top of ultrathin crystalline NbN superconductor.[187] As an alternative to PVD, the chemical vapor deposition (CVD) can introduce a diversity of chemical reactions to the material synthesis. In CVD growth, the reactant gases will be pumped into the reactor, and thereafter, the heterogeneous reactions at the gas–solid interface create the new phases to be deposit onto substrate,[192] offering high flexibility for producing 2D heterostructures with tunable chemical component at controlled manner.[45,193]
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Epitaxial CVD growth of ordered 2D heterostructures normally has a lattice match between two growing 2D materials, or between 2D material and the substrate.[194] For instance, Ruzmetov et al.[195] reported the CVD epitaxial growth of single-layer 2D MoS2 on 3D GaN substrates, as GaN and MoS2 have the same hexagonal crystal system with only a small lattice mismatch <1% (i.e., GaN = 3.19 Å, MoS2 = 3.16 Å). The GaN has m-plane (1100) orientation, and the sides of MoS2 triangles will be aligned on m-plane of GaN with an alignment of lattices of GaN and MoS2. Fan et al. also reported the epitaxial growth of 2D/1D CsPbBr3/Bi2O2Se heterostructure for high-performance applications in electronics and optoelectronics (Figure 15e–h).[196] Given the lattice mismatch of Bi2O2Se and CsPbBr3 (−1.8% for 3 × ≈ dCsPbBr3(001) and −7.9% for 2 × ≈ dCsPbBr3(1-10)) shown in Figure 15i,j, an epitaxial growth of CsPbBr3 nanowires on Bi2O2Se nanoplates is feasible following the epitaxial relationships of [001]||[200] and [1–10]||[020].
The CVD growth is a powerful tool to prepare intriguing 2D TMDs heterostructure, such as WS2–WSe2, WS2–MoSe2, WS2–WSe2–MoS2, WS2–MoSe2–WSe2, and WS2–WSe2–WS2–WSe2–WS2, and so on.[197] Since the CVD growth does not require high-vacuum working environments,[192] Zhao et al. proposed a water-assisted CVD growth of 2D WS2–MoS2 heterostructure.[198] In this process, thermal dehydration of solid CaSO4.2H2O creates a water vapor, and the amount of water vapor controls the supersaturation of vapor reactions (Figure 16a). It is suggests that a low supersaturation prefers the screw dislocation growth mode of MX2, while a high supersaturation may trigger a layer-by-layer growth mode.[198–201] Hence, by switching the MX2 source in CVD growth and controlling the supersaturation, the lateral epitaxial growth of MoS2 on the edge of WS2 is accomplished (Figure 16b–e). Nguyen et al. also reported a lateral 1 T’ RexMo1–xS2-2 H MoS2 heterostructure by two-step CVD growth process.[202] In CVD furnace, a substrate coated with liquid precursor of NaReO4 and NaMoO4 is heated to an elevated temperature of 680 °C. The formation and growth of 1 T′ RexMo1–xS2 occur by pumping ammonium sulfide atmosphere. Under further increase of temperature to 800 °C, the lateral 1 T′ RexMo1–xS2-2 H MoS2 heterostructure appears, where the outer 2 H MoS2 and the inner 1 T′ RexMo1−xS2 were separated, the 1 T′ RexMo1−xS2 remains a monolayer at ≈0.74 nm over the whole region (Figure 16f–i). The coherence with the junction relationship of (2-20)ReS2//-110)MoS was revealed by annular dark-field scanning transmission electron microscopy (ADF STEM) at the atomically sharp interface between 1 T′ RexMo1−xS2 and 2 H MoS2. Compared with many TMDs catalysts, the lateral 1 T′ RexMo1–xS2-2 H MoS2 exhibits higher hydrogen evolution reaction (HER) catalytic performance with an overpotential of ≈84 mV at current density of 10 mA cm−2 and Tafel slope of 58 mV dec−1.
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Therefore, to produce a lateral- or vertical-stacking 2D heterostructure, it requires a selective nucleation and growth of 2D material on the edge or top of seed 2D crystal, respectively, which can be controlled through tuning the flow rate and the concentration of vapor precursors, as well as the reaction temperature and growth time of system.[45,203,204] Gong et al. revealed that the temperature-dependent CVD growth can create either vertically stacking or in-plane stacking WS2/MoS2 heterostructure.[205] The lateral growth of WS2 on edges of MoS2 occurs around 650 °C, while the in-plane WS2/MoS2 with intimate contact and atomically sharp interface appears at 850 °C. In addition to TMDs, CVD growth can integrate atomic layer of graphene and h-BN to create in-plane 2D carbon-based heterostructure.[206] Geng et al. proposed an one-pot confined CVD growth of large-area, in-plane graphene/h-BN heterostructure (Figure 17a–d).[206] In this process, the growth of h-BN is templated by the pre-grown hexagon graphene, and the in-plane graphene/h-BN has the inner hexagonal graphene crystal decorated by outer h-BN ribbons at uniform continuity. Also notice both graphene and h-BN are monolayer. Likewise, Gigliotti et al. reported the highly ordered h-BN/graphene deposited on silicon carbide by lateral epitaxial deposition,[207] and the h-BN layers are atomic thin and orderly stacking at interface (Figure 17e–g). Owing to the epitaxy of the h-BN film on graphene, the ordered grain boundaries can span in region over microscale.
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Worthy of note, van der Waals (vdW) integration can create a large family of vdW heterostructures (vdWHs) beyond the limit set by lattice matching.[208] Duan et al. proposed a general route to create 2D vdWHs arrays between semiconducting TMDs (s-TMDs) and metallic counterparts (m-TMDs) in spite of their lattice differences.[209] To this end, they first patterned periodic arrays of nucleation sites on monolayer or bilayer s-TMDs (e.g., WSe2, WS2, MoS2), on which the selective nucleation and growth of additional m-TMDs can form a periodic array of m-TMD/s-TMD in vdWH (Figure 18a). Elemental mapping reveals the spatially elemental distribution of W, Se, V, and Si in these vdWH with atomically sharp boundary along the interface (Figure 18b,c).
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Many other vdWHs have also been reported, such as Bi3O4Cl/g-C3N4,[210] MOF-on-MOF,[211] and GeS–SnS.[212] For instance, only by mixing of g-C3N4 and Bi3O4Cl in ethanol solution under ultrasonic treatment, the vdWH of Bi3O4Cl/g-C3N4 is formed and carries an built-in electronic field that enables a fast separation of electron–hole pairs for reaching high photocatalytic CO2 reduction activity (Figure 19).[210] In a stamping strategy proposed by Yao et al., the vdW force can create the highly oriented MOF-on-MOF thin film, by integration of second MOF layer onto the seed MOF layer, for example, the deposition of Cu–TCPP layer onto the a Cu–HHTP (HHTP=hexahydrotriphenylene) layer (Figure 20).[211] The repeating stamping allows a control of the thickness of Cu–TCPP layer on Cu–HHTP. The preparation of lateral 2D heterostructure is also feasible through vdW assembly. Sutter et al. established the high-quality vdWHs composed of central SnS seed crystal laterally connected on the edge of GeS.[212] The direction of these lateral heterostructure interfaces comply with the {110} face of the SnS seeds, while vertical (001) vdW interface was formed combine the top layer of GeS with the SnS flakes center. TEM exhibits a fine central region surrounded by a wide edge band with different contrast and further identified by SAED. The top layer of GeS and the seeds of SnS have the same crystal structures with anisotropic lattice mismatch. The mismatch is small along the [010] direction (denoted as b) with (bSnS − bGeS)/bSnS = −0.27%, but big mismatch is along [100] (denoted as a) with (aSnS − aGeS)/aSnS = 8.9%. The bilayer TMDs heterostructures can deliver a myriad of applications such as for optoelectronics and valleytronics. In light of such principle, the rising community of 0D, 1D, and 2D nanounits can provide a rich library of building blocks for creating assembled vdWHs.[213,214]
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Perspective and Challenge
Nowadays, heterostructured nanomaterials constitute one of the most intriguing categories of functional materials to alleviate the energy and environmental crisis. To access the aesthetic structures and optimal performance, researchers have fueled constant interests to the construction of purposeful ordered heterostructures of intriguing interfacial synergy and escalated properties. Notwithstanding the advances over past years, significant challenges and opportunities still lie ahead, putting them out of reach for practical applications.
First, the precise control of assembly of constituent nanounits in heterostructures remains a huge challenge.[76,215] Given that the formation, growth, and assembly of nanounits in heterostructures under different conditions all proceed through a dynamic processes, the liquid-phase transmission electron microscopy (LP-TEM) and cryogenic TEM (cryo-TEM) can be powerful and in situ tools to examine the time–course nucleation, crystallization, growth, and assembly of nanomaterials.[216] In this process, unraveling the growth mechanism and the driving force involving can provide key insights into the correlation between the selection of nanounits and their ordered self- or directed assembly for the generation of ordered heterostructure. In addition to the structural control, a proposal of green, facile, and scalable strategy for the synthesis of ordered heterostructures comprising earth-abundant elements is also highly demanded, for the sake of large-scale applications, given that many developed strategies still suffer from drawbacks of tedious procedure, high cost, low product specificity, etc.
Second, engineering the heterointerface is vital to reach optimal physicochemical properties of heterostructures, while the properties of ordered assembly built of nanoscale or atomic entities can exhibit properties distinct from the macroscopic counterparts. Engineering the interfaces of heterostructure at nanoscale and atomic level is then imperative, yet remains very challenging. For many reported heterostructures, constituent nanounits are normally physically attached with each other, creating interfaces with large boundaries and random orientation.[217] In this aspect, lattice-match effect could drive the oriented assembly of nanounits, offering a facile and powerful tool to the creation of well-defined heterointerface with defined atomic configuration and regulated defect structure.[169,185] As emerging heterogeneous materials, the single-atom catalysts (SACs) have witnessed their great superiority for field applications,[165,166,218,219] which heavily relies on a controllable regulation of the metal active sites and surrounding microenvironment of single atoms in SACs heterostructures. As such, Zhou et al.[220] prepared a heterostructure comprising NiO/Ni deposited on Ag NWs via electrodeposition strategy, and thereafter the single-atom Pt can be anchored selectively on interface of NiO/Ni. They found that the Pt atoms at the NiO/Ni interfaces can induce a higher occupation of Pt 5 d band at the Fermi level, accounting for the enhanced performance of catalysts toward the HER. To date, the robust and general principles to engineer the interface of heterostructure at atomic level and strengthen the involving interactions awaits to be exploited.
Third, as the microstructure (phase, size, and shape), spatial location, and distribution as well as the interaction of nanounits can all impact the properties of ordered heterostructures, it becomes elusive and sometimes controversial regarding the description and prediction of “form-to-function” relationship of ordered heterostructures. It is then timely required to combine the advanced characterization technologies, for example, aberration-corrected STEM, AC-STEM, or in situ Fourier transform infrared spectroscopy, in situ Raman, to examine both on the structures and properties of heterostructures under in situ environment, and the DFT calculations to describe and predict the geometrical structure and electronic properties of heterostructures at atomic level.
In this review, we briefly summarized the recent process on the rational design of ordered heterostructures, which could open door to advance the material design and underpin the fundamental knowledge on the growth and assembly mechanism of nanomaterials to afford complex yet efficient architectures. Ordered structures assembled from functional units have long been existent in nature. Should artificially ordered assembly of nanounits with a wider selection of building components into purposeful heterostructures be readily achieved, greater versatility can become possible in constructing a myriad of functional materials. We anticipate that this review can fill the currently missing part of ordered assembly and derived heterostructures so that a complete picture from material design to field applications can be realized for relieving the energy and environmental crisis.
Acknowledgements
T.S.Z. and B.X.Y. contributed equally to this work. This work is financially supported by the National Key Research and Development Program/Key Scientific Issues of Transformative Technology (Grant no. 2020YFA0710303), the National Natural Science Foundation of China (Grant nos. U1905215 and 52072076), the Fujian Science Foundation Grant (Grant no. 2022J01554), and the Key Project of Science and Technology Innovation of Fujian Provincial Department of Education (Grant no. 2022G02002).
Conflict of Interest
The authors declare no conflict of interest.
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Copyright John Wiley & Sons, Inc. 2023
Abstract
Precise synthesis of high‐quality, sophisticated heterostructures by ordered assembly of small nanomaterials is a key step to gain advanced materials that have elaborate functionalities, collective properties, and enhanced stabilities for mitigating the energy and environment crisis. Intricating in the structure, size, and shape of nanomaterials, ordered assembly of isotropic or anisotropic nanoscale building blocks to create specified heterostructures remains challenging, owing to the extraordinary challenges in design of lattice topology of distinct nanounits and in control of their crystallization, growth, and assembly mechanism/kinetics. Herein, the emerging methodologies to prepare a diversity of ordered heterostructures with strengthened particle–particle interaction are examined and synergistic properties are enhanced. It is aimed to unlock the principles to regulate the geometrical and electronic properties of these intriguing kinds of heterostructures for respective sustainable energy and environmental applications. Current challenges and opportunities in customization of ordered heterostructure at the nanoscale and atomic level are also discussed.
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Neither ProQuest nor its licensors make any representations or warranties with respect to the translations. The translations are automatically generated "AS IS" and "AS AVAILABLE" and are not retained in our systems. PROQUEST AND ITS LICENSORS SPECIFICALLY DISCLAIM ANY AND ALL EXPRESS OR IMPLIED WARRANTIES, INCLUDING WITHOUT LIMITATION, ANY WARRANTIES FOR AVAILABILITY, ACCURACY, TIMELINESS, COMPLETENESS, NON-INFRINGMENT, MERCHANTABILITY OR FITNESS FOR A PARTICULAR PURPOSE. Your use of the translations is subject to all use restrictions contained in your Electronic Products License Agreement and by using the translation functionality you agree to forgo any and all claims against ProQuest or its licensors for your use of the translation functionality and any output derived there from. Hide full disclaimer
Details
1 Key Laboratory of Advanced Materials Technologies, Fuzhou University, Fuzhou, Fujian, China