3D Printing of glasses with tunable UV–VIS–IR photoluminescence via low-temperature nanoscale engineering

Domain: physics
Published: 2026-01-19
Editorial outcome: minor_revision

Abstract
Recent breakthroughs in three-dimensional (3D) printing of glass have enabled precise shaping of glass structures from macro to micro scale while retaining their excellent physical and chemical properties. However, current functional developments in 3D printed glass primarily focus on leveraging its intrinsic optical transparency, with limited exploration of photonics application. 3D-printed glass functionalized by quantum dots (QDs) will hold great promise for creating desired photonic functional macro- and microstructures. However, voxel-precision customization of photonic functions in 3D printed glass by the monodispersed and size-controlled QDs is fundamentally conflict with the high-temperature sintering process required in 3D printed glass. Here, we present a simple, cost-effective method for creating a library of QDs functionalized 3D-printed glasses with tunable UV-VIS-IR photoluminescence (PL) by printing nanoporous glass (NG) followed by low-temperature nanoscale engineering of QDs growth. The integration of uniform and size-controlled QDs imparts 3D printed glass with fine tuning in composition, architecture, and photonic functionalities, while nanoscale engineering of the micro-environment enables precise tailoring of various QDs with enhanced optical performance and stability. This universal approach establishes a benchmark for 3D-printed functional glasses with advanced photonic capabilities.

Introduction
Glass stands as one of the most important materials known to mankind, with exceptional optical transparency, thermal/chemical stability, mechanical properties, and composition tunability, and has become a cornerstone of advanced photonic systems. In the realm of glass manufacturing, 3D printing, also known as additive manufacturing, holds the potential to combine the automation of modern industrial methods with a level of geometric complexity that exceeds the limitations of traditional artisanal techniques1–5. Recent advances, including stereolithography (SLA)3, digital light processing (DLP)6, two-photon polymerization (TPP)7, and computed axial lithography8 have demonstrated success in fabricating silica-based optical components through photopolymerization of resin–glass composites. These technologies enable the integration of glass’s outstanding physical and optical properties into the creation of arbitrary macro- and microstructures2,3. Further photonic functionalization is highly desirable.

Photonic functional glasses, such as glass containing metal and metal oxide nanocrystals9, rare-earth ions10, or QDs11, have played crucial and central roles in the birth and development of photonics, which deals with the generation, transmission, modulation, amplification, conversion, and detection of light12–15. Among them, semiconductor QDs exhibit quantum-confined emission tunable from ultraviolet to visible to infrared (UV–VIS–IR), coupled with the highest photoluminescence quantum yields (PLQY) up to 100%, narrowband emission, and a broad range of tunable chemical and physical properties due to their large surface-to-volume ratio16–20. 3D printing of QD composite offers a fabrication route to engineer devices with advanced functional properties. The challenge is that current 3D printing technologies for QDs typically utilize photocurable resins, which suffer from low mechanical strength and degraded properties21. In contrast to resins, glass possesses exceptional optical transparency and physical properties. However, the high-temperature treatment required to remove organic compounds in glass 3D printing technology will cause severe aggregation due to Ostwald ripening and surface ligand decomposition of semiconductor QDs9,22,23, which will prohibit the integration of QDs with 3D-printed glass. Moreover, uniform monodispersed distribution and precise size control of QDs in 3D-printed glass are hindered by uncontrollable growth kinetics and weak confinement effect in the solid medium. In addition, given the rich chemistry of QDs, a universal methodology is highly desired for the variety of QD systems in 3D-printed glass.

Here, we overcome these challenges by developing a universal nanoscale post-functionalization for 3D-printed glass, which offers a low-temperature, precise micro-environment chemistry strategy enabling the integration of QDs accessible to modern 3D printing glass techniques. By 3D printing NG with desired metal ion precursors, a library of 3D printed glasses functionalized by QDs were constructed relying on the spatial and chemical dual nanoconfined growth of diverse QDs (Fig. 1). The method is to exploit the controllable diffusion of metal ions in the confined environment, which tailors the delivery of precursors to the reaction and constrains the nanocrystal size in the quantum confinement regimes. This post-functionalization strategy circumvents the thermal limitations that lead to QD degradation. In addition to its convenience and energy-saving benefits, the nano-reaction space provided by this low-temperature approach also effectively prevents the aggregation of QDs, as it avoids the Ostwald ripening process that typically occurs during the melt-quenching route. Considering the inherent flexibility in incorporating QDs in 3D-printed glass via precise micro-environment chemistry, this will inspire the development of approaches to 3D printing of functional glass with the precise control over complex geometry, broad chemical compositions, and tunable photonics properties. These QDs functionalized 3D-printed glasses exhibit broad tunable PL ranging from UV (300 nm) to near-IR (2 μm) with luminescence lifetime from 16 to 699 ns and demonstrate voxel-level control over the QD bandgap. Benefiting from the multi-functionalization of the 3D-printed glass, various promising applications, such as photocatalysis, information encryption, and decryption, could be realized.

Results
3D printed glasses functionalized by QDs
Figure 1a systematically details the manufacturing paradigm through four evolutionary stages: sol–gel ink formulation, vat photopolymerization 3D printing of green body, 3D-printed metal ions-doped NG (M-NG) formation, and final QD post-functionalizing. The engineered sol–gel ink comprises a photopolymerizable silica–alumina glass system optimized for vat photopolymerization printing, with functional additives enabling dual-phase control of nanoscale confinement environment and QD crystallization. Key components include: Tetraethyl orthosilicate (TEOS) as silica precursor, 3-(Trimethoxysilyl)propyl acrylate (APTMS) as covalent network reinforcement (Table S1). Metal salts [Pb(NO3)2, Cd(NO3)2, Zn(NO3)2, AgNO3, and In(NO3)3] were added as precursors for Pb based [CsPbX3 (X=Cl, Br, I), PbS, PbSe)], Cd based (CdS, CdSe, CdTe), ZnS, Ag2S, and AgInS2 QDs, respectively. Sudan Orange G was added as a photo-absorber to enhance vertical resolution and maintain dimensional accuracy. Besides, aluminum lactate is selected as a nanopore former in sol–gel-derived M-NG. TEOS, aluminum lactate, and APTMS undergo hydrolysis and condensation reactions, jointly contributing to the formation of a spatial network structure modified with acrylate groups5,6,24.

A moderate drying step transformed the printed wet gel structure into xerogel, followed by low-temperature treatment at 650 °C in air2, which converted the green body into 3D-printed M-NG, resulting in linear shrinkage of ~40%. The heat treatment decomposed organic compounds and removed carbon through oxidation. The thermal gravimetric analysis (TG) curve of the green body shows no significant weight change above 650 °C, indicating the complete conversion to inorganic glass (Supplementary Fig. 1). Differential scanning calorimetry (DSC) analysis of NG reveals a typical glass transition temperature (Tg) of 747 °C (Supplementary Fig. 1). The composition of NG is confirmed by the energy-dispersive X-ray spectroscopy (EDS) spectra, exhibiting an atomic percentage of all elements consistent with the designed SiO2–Al2O3 network and no detectable carbon contamination (Supplementary Fig. 2). Raman spectroscopy further confirms the complete removal of organic components after sintering, showing spectra that closely resemble those of commercial fused silica (Supplementary Fig. 3). The as-prepared 3D-printed NG showcases the nanoporous structure with a pore size centered at 4 nm with smooth surface (Rms ~1.1 nm) and excellent transparency over 90% in visible range (Supplementary Figs. 4–6), which is highly suitable for optical devices. Due to the presence of nanopores, NG has a slightly lower refractive index compared to commercial fused silica glass (Supplementary Fig. 7).

All-inorganic perovskite (CsPbX3) QDs functionalized 3D-printed glasses are used here as a representative system. The doped metal ions (Pb2+) in M-NG primarily exist as glass modifiers, distributed uniformly in the form of single atoms or atomic clusters to facilitate QD crystallization, as shown in high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images (Supplementary Fig. 8). Then, the 3D printed M-NG was saturated in QD precursor solution to facilitate the in-situ growth of QDs. Ultimately, post-functionalization of 3D-printed glass by embedding QDs subsequently enabled the fabrication of arbitrary QD glass structures with complex spatial architectures and rationally designed optoelectronic properties (Fig. 1b and Supplementary Fig. 9). The voxel size is 180.0×180.0×5.6 μm, with QDs uniformly distributed within a single voxel (Supplementary Fig. 10). The optical image (Fig. 1c) shows the printed Loong-embedded CsPbBr2I QDs, revealing intricate spatial details and texture, emitting orange light under UV excitation.

This low-temperature nanoconfined strategy for 3D-printed glass is generally applicable for various semiconducting QDs, including lead halide perovskite, CdS, CdSe, CdTe, PbS, PbSe, ZnS, AgInS, and Ag2S (Fig. 1d and Supplementary Fig. 11). Corresponding transmission electron microscopy (TEM), HAADF-STEM micrograph, and EDS mapping of 3D printed glasses functionalized by QDs confirm their composition (Supplementary Figs. 12–20). These glasses exhibit emission with tunable peak from UV (380 nm) to near-infrared region (1.6 μm), with lifetimes varying from 16 ns for CsPbBr3 QDs to 699 ns for PbS QDs (Supplementary Fig. 21). In addition, this method establishes two-dimensional bandgap engineering capabilities within monolithic glass objectives, permitting independent control of both QD composition and physical dimensions at micron resolution. By adjusting the QD composition, specifically the halide ions (from CsPbBrxCl3−x to CsPbBryI3−y), this approach enables an expected emission peak tunability spanning from 450 to 710 nm (Supplementary Fig. 22a). Furthermore, by regulating the pore size of M-NG (from 2.6 to 7.3 nm) through tuning the Al content, the size of CsPbBr3 QDs can be precisely controlled, enabling fine tuning of the emission peak from 465 to 510 nm (Supplementary Figs. 22b, 23). Within a single 3D-printed glass object, a controllable spatially resolved red-to-blue color gradient across the structure is achieved through spatial sectionalized treatment with different QD precursor solutions, and green-to-blue dual emission confined to designated regions is realized through selective tuning of QD size (Fig. 1e). Overall, our method successfully enables the QDs functionalized 3D-printed glass with high spatial resolution and complex geometries, allowing for the synthesis to a variety of QDs with controllable emission and lifetime. This would provide insights for the functionalization of the 3D-printed glass.

Low-temperature nanoscale engineering of QDs growth kinetics
Achieving monodisperse QDs with uniform sizes within a glass matrix has long been a central challenge for QD-functionalized glass, and rational growth kinetics control has proven to be an effective strategy for achieving monodispersed QDs25,26. To address this, we propose a precisely engineered micro-environmental chemistry strategy. Our approach incorporates two key considerations. First, we utilize metal ion dopants in the M-NG matrix as the QD precursor source. These ions are gradually released into the confined nanochannel and become enriched, forming spatially localized heterogeneous nucleation centers (Fig. 2a). Unlike conventional glasses, where atomic diffusion is largely frozen below the Tg, the M-NG exhibits notable ion mobility even at room temperature (Supplementary Fig. 24). This behavior stems from its exceptionally high specific surface area (~564 m²/g), which exposes a substantial fraction of atoms at the internal pore surfaces27. Based on molecular dynamics (MD) simulations, this surface contribution was estimated to correspond to over 60% of the total atoms (Supplementary Fig. 25). The high specific surface area and the nanoscale thin pore walls of the porous glass matrix result in elevated surface energy, which significantly facilitates the diffusion of Pb ions within the glass. As a result, the spatial distribution of the synthesized QDs could inherit the uniform distribution of Pb2+ introduced during the sol–gel process, resulting in uniform monodispersed QDs within NG. Second, during the reaction with precursor solution, the precursors undergo rapid reaction upon reaching the nucleation threshold25. The QD experiences rapid growth while being spatially confined by the nanochannel of the M-NG matrix. Meanwhile, the precursor depletion drives the system rapidly into the Ostwald ripening regime, in which smaller nuclei that fail to achieve sustainable growth undergo selective dissolution28. Under the dual regulation of precisely controlled precursor release and nanoconfined growth, a natural size-focusing process is achieved within the nanochannel, ultimately ensuring the exceptional size uniformity of the resulting QDs in M-NG.

We used perovskite CsPbBr3 QDs (PQDs)—a representative system with notoriously fast growth dynamics26—as a model to clarify our precisely engineered micro-environmental chemistry strategy. We monitored the formation process of PQDs through a combination of synchrotron in-situ small-angle X-ray scattering (SAXS) and in-situ absorption spectroscopy (Fig. 2b, c). SAXS analysis reveals the time-dependent evolution of PQD concentration, which shows a rapid initial increase followed by a gradual slowdown, eventually reaching saturation at ~1500 s (Fig. 2d top). Ex-situ XPS analysis further reveals the prolonged reaction process (Supplementary Fig. 26). Complementary in-situ absorption spectroscopy (Fig. 2c, d bottom) provides a similarly resolved concentration profile, enabling efficient and facile parametric screening. Based on the evolution of SAXS patterns, the average size of synthesized QDs has no significant change during the reaction and is estimated to be 4 nm (Fig. 2d right and Supplementary Text S1), which corresponds to the average nanochannel diameter of 4 nm. In addition, the TEM image shows the size distribution of PQDs ranging from 2 to 6 nm (Supplementary Fig. 27), which is consistent with the in-situ SAXS results. These results highlight the effective spatial confinement imposed by the nanoscale reaction environment on QD growth. In contrast to the extraordinarily rapid reaction kinetics typically observed in PQD synthesis, our nanoconfined synthetic strategy results in a significantly prolonged and controllable formation process, generating uniform and monodispersed QDs.

To gain further insight into the factors affecting the diffusion-regulated growth process, we systematically investigated the influence of precursor conditions, specifically, the Pb2+-doping level in the glass matrix and the concentration of CsBr in solution, on the reaction kinetics. Obviously, the Pb2+-doping level primarily determines the final absorbance, which is directly related to the concentration of QDs, while an increase in CsBr concentration significantly shortens the time required for the reaction to reach saturation (Fig. 2e, f). These observations imply that CsBr diffusion serves as the key rate-limiting factor governing the reaction kinetics in this nanoconfined system. We employed a one-dimensional transient diffusion model based on Fick’s second law to describe the time-dependent concentration profile of CsBr diffusing into the NG matrix (Fig. 2g)29,30. Detailed procedure and discussion are shown in Supplementary Text S3. Once the CsBr concentration in the NG reaches a critical threshold (Cc), CsPbBr3 rapidly crystallizes and fills the nanopores, forming the QD formation region. As CsBr continues to diffuse inward, the QD formation region propagates from the glass surface toward the center, and the reaction is saturated when this region reaches the core of the glass matrix. We observed good agreement between theoretical predictions and experimental results concerning the dependence of reaction saturation time on CsBr concentration (Fig. 2h), confirming that CsBr diffusion is the primary rate-limiting factor governing QD growth in this heterogeneous system and leading to a diffusion-mediated growth mode.

Physical and chemical dual confinement of QDs
The 3D printed PQD glass of the Oriental Pearl Tower model exhibits bright blue-green emission under UV excitation (Fig. 3a inset), which can be attributed to the high transparency of 3D-printed glass (Supplementary Fig. 29), benefiting from uniform QD distribution inside and the sub-diffraction limit size of the nanopore, as shown in Fig. 3a. The high-resolution transmission electron microscopy (HRTEM) image of PQD glass reveals a lattice spacing of 2.92 Å, assigned to (200) crystal facet of CsPbBr3 31. The X-ray diffraction (XRD) and wide-angle X-ray scattering (WAXS) of PQD glass (Fig. 3b and Supplementary Fig. 30) show diffraction peaks indexed to the cubic CsPbBr3. Importantly, the uniform distribution of PQDs throughout the M-NG is supported by cross-sectional SEM and EDS elemental mapping, as well as PL mapping from both the surface and interior, all of which indicate homogeneous QD formation and emission (Supplementary Figs. 31–34). This is attributed to the spatially confined and gradual growth mechanism in our precise micro-environment chemistry strategy. The transient absorption (TA) also shows a single bleaching peak without any energy transfer process, indicating that few QD aggregations occur (Supplementary Fig. 35). These results validate the efficient and homogeneous nanoconfined growth of PQDs within M-NG.

In-situ confined-grown crystals typically exhibit enhanced interactions with the substrate material, which can improve the overall stability and homogeneity of the material24,32,33. We further investigate the interfacial interaction between confined PQD and M-NG through X-ray absorption spectroscopy (XAS) and high-resolution XPS. X-ray absorption near edge structure (XANES) spectra show that the Pb L3-edge absorption of M-NG and PQD glass is shifted to higher energies compared to that of Pb foil, indicating that lead in these samples exists in an oxidized state (Supplementary Fig. 36). The k2-weighted Fourier transform extended X-ray absorption fine structure (EXAFS) and wavelet transform (WT) analysis (Fig. 3c) reveal a Pb–O peak for M-NG at 1.6 Å (phase uncorrected). The Pb–O bonds could facilitate the bonding of QDs with the glass matrix at the interface. In contrast, the peaks for the PQD in 3D-printed glass are primarily observed at 2.5 Å after precursor reactions, which can be assigned to Pb–Br, with an additional peak at low R region. WT analysis (Fig. 3d) reveal that the additional feature originates from a light element, suggesting that oxygen in the glass is the likely backscatterer. To exclude the influence of unreacted PbO in the QD growth, we conducted inductively coupled plasma (ICP) tests (Table S2). The result reveal that the 3D-printed glass has an almost equal atomic ratio of Cs to Pb, which confirm that most of the Pb element can be converted into PQDs. Furthermore, comparing high-resolution XPS spectra of Pb4f in M-NG and PQDs glass (Supplementary Fig. 37), the Pb4f peaks exhibit a discernible shift to lower binding energy due to the conversion from Pb–O to Pb–Br. These values are lower than those of colloidal CsPbBr3 QD due to the transfer of electron density from Pb atoms to SiO4 tetrahedra within the glass. These results imply the formation of Pb–O bonds at the interface between PQDs and the M-NG matrix (Supplementary Fig. 38). Density functional theory (DFT) calculations reveal that interfacial Pb–O bonding can effectively passivate surface defects of QDs, thereby enhancing their optical performance. Br vacancy leads to the formation of a trap state that is located at ∼0.6 eV below the conduction band with a localized charge density within the lattice (indicated by yellow symbols), acting as an electron-trapping center for nonradiativerecombination (Fig. 3e, f). Once the Br vacancy on the surface is modified by O, it removes the midgap states through contributions from the glass matrix, resulting in a significant reduction of wavefunction localization within the materials (Fig. 3g, h). Additionally, DFT calculations show that the interaction between Pb and O at the CsPbBr3/SiO2 interface leads to a higher adhesion energy, indicating a more stable interface for the QDs34 (Supplementary Fig. 39).

The physical and chemical dual nanoconfined environment of M-NG effectively regulates the emission properties of QDs. The temperature-dependent (77–300 K) PL of the PQDs glass revealed a higher EA value (exciton activation energy from thermal dissociation) of 165 meV, twice that of colloidal QDs (79 meV) (Supplementary Fig. 40 and Supplementary Text S4). The significant EA value and minimal change in PL intensity suggest that the exciton in PQD glass has a higher energy threshold for thermal dissociation due to the enhanced spatial confinement of electrons and holes in M-NG35. Owing to interfacial defect passivation and enhanced exciton binding energy, the PQD glass exhibits a high PLQY up to 82%. Moreover, the physical isolation by the nanopores along with the interfacial chemical bonding effectively enhances the stability of the QDs. As shown in Fig. 3i, PQD glass maintains 80% of its initial PL intensity after 180 days in ambient conditions of 25 °C and 50% relative humidity (RH), highlighting its excellent long-term stability. In contrast, the PL intensity of colloidal QDs decreased to 23% of the initial intensity after 30 days. Even under high humidity conditions (25 °C, 80% RH), the PQD glass retains 74% of its initial PL intensity after 600 h, while the pure QDs only retain 15% after 117 h (Supplementary Fig. 41). Moreover, we assess PQD glass photo-stability under continuous 405 nm, 120 mW laser irradiation in ambient conditions. The PQD glass retains 86% of its PL intensity after 16 h, while the colloidal QD film decreases to 44% (Supplementary Fig. 42).

Demonstration of 3D multifunctional glass architectures
This approach for 3D printing of QD glass can be used to fabricate multifunctional photonic and optical devices, which was first demonstrated by printing QD glass architecture for photocatalytic CO2 reduction (Fig. 4a). The 3D design freedom offered by this additive manufacturing approach enables simultaneous optimization of multiple performance-determining parameters in photocatalytic systems. Specifically, the integration of QDs into a transparent, M-NG matrix provides both efficient photon harvesting and spatially confined photocatalytic reaction environments, while 3D printing allows for deterministic control over hierarchical geometric complexity that surpasses conventional fabrication, including the realization of bioinspired fractal surface features. These structures emulate nature’s efficient light-harvesting designs, such as tree-like or forest architectures36, and are known to enhance photon scattering, increase illuminated surface area, and promote mass transfer efficiency in photocatalytic processes. Finally, we fabricated three types of 3D-printed architectures with different surface fractal micro-features: smooth surface (#1), cylindrical microarray (#2), and channel-like microarray (#3) (Fig. 4b and Supplementary Fig. 43). As shown in the SEM images, the 3D-printed QD glass exhibits well-defined microscale fractal surface features, characteristic of the designed hierarchical architecture. Control experiments using pristine M-NG without PQD loading show negligible photocatalytic activity, confirming that the CO2 reduction functionality originates from the post-functionalized PQDs (Fig. 4c). Importantly, the introduction of surface microstructures significantly enhanced the photocatalytic performance: dome structures with increasing geometric complexity exhibited progressively higher gas production rates compared to the planar QD glass. The CH4 evolution rate increases from 2.9 to 3.2 μmol g−1 for #1, 4.6 μmol g−1 for #2 and 7.4 μmol g−1 for #3, representing an improvement up to 2.6-fold. Similarly, the CO evolution rate rises from 4.5 μmol g−1 to 4.8 μmol for #1, 9.9 μmol g−1 for #2 and 13.4 μmol g−1 for #3, achieving an increase up to 3-fold. Moreover, to showcase the potential of this 3D-printing technique as a multifunctional photonic platform, we printed QD-embedded glass architectures as building blocks for optical information encryption and decryption (Supplementary Fig. 44). By spatially doping different metal ions and triggering selective luminescence through precursor-mediated reactions, we realized chemically responsive optical information encryption and decryption.

Methods
Chemicals
The materials were used as received without further purification: Aluminum lactate (L-Al, 95%) and Diphenyl (2,4,6-trimethyibenzoyl) phosphine oxide (TPO, 98%), 3-(trimethoxysilyl) propyl Acrylate (APTMS, 99%), tetraethyl orthosilicate (TEOS, 99%) were purchased from Adamas-Beta. Sudan orange G (96%), ethanol (99.9%), N,N-dimethylformamide (DMF, 99.5%), Tris–hydrochloric acid buffer (0.05 M, pH=7.4), 1-butanol (99.5%), methyl acetate (99%), octylamine (99%), oleic acid (99%), and HCl (AR, 37 wt%) were purchased from Sigma-Aldrich. CsBr (99.99%), CsCl (99.99%), CsI (99.99%), Na2S (99.9%), Pb(NO3)2 (98.5%), Cd(NO3)2·4H2O (99%), In(NO3)3 (99.9%), AgNO3 (99%), Zn(NO3)2·4H2O (AR), selenium (99.9999%), tellurium (99.999%), sodium sulfite (98%) were purchased from Aladdin. MABr (99.5%) and FABr (99.5%) were purchased from Xi’an Yuri Solar Co., Ltd. Ammonium acetate (99.9%) and PbBr2 (99%) were purchased from Macklin.

Preparation of UV-curable sol ink
In our case (as can be calculated from Table S1), the sol–gel ink was prepared as follows: first, a solution (M1) was prepared by dissolving L-Al, a metal salt, and deionized water. The metal salts included one or more of the following: Pb(NO3)2, Cd(NO3)2·4H2O, AgNO3, Zn(NO3)2, and In(NO3)3. L-Al and metal salts were weighed in proportion and dissolved in deionized water, then stirred with a magnetic stirrer for 30 min. Second, a mixture of TEOS and ethanol (EtOH) (denoted M2) was added to M1 during the stirring process by slowly adding dropwise, and stirring the mixture until TEOS was fully reactive. Then, the M3 composed of 3-(Trimethoxysilyl) propyl Acrylate (APTMS) and Diphenyl (2,4,6-trimethyibenzoyl) phosphine oxide (TPO) was added, then stirred with a magnetic stirrer for 30 min. Finally, NH4Ac solution (denoted M4) was added and sealed in a dark vessel and stirred for ≈45 min until the formed sol could form a gel matrix upon exposure to UV radiation with a wavelength of 405 nm. The weight percentage L-Al:TEOS:H2O:EtOH:APTMS:TPO:NH4Ac at all stages of the process is presented in Table S1.

3D printing of M-NG
The 3D models of a sample are designed by Blender and printed using a commercial DLP 3D printer (Max UV, Asiga). The printing parameters include using a UV-LED light source (405 nm) with a light intensity of 15 mW cm−2. After printing, the structures were placed in a closed vessel with a tiny hole and aged, dried in a 50 °C oven for 7 days. Then, dried samples were sintered at 650 °C (0.5 °C/min to 650 °C and kept for 120 min) in a muffle furnace to obtain 3D-printed M-NG.

In-situ confined growth of QDs
Precursor preparation. For perovskite precursors: CsX (X=I, Br, Cl): Dissolved in anhydrous methanol (5 mg/mL) with magnetic stirring at 70 °C for 2 h in a glove box. MABr/FABr: Dissolved in n-butanol (5 mg/mL) with magnetic stirring at 70 °C for 2 h in a glove box. For chalcogenide precursors: Na2S: 0.1 M solution in Tris–HCl buffer (pH 7.4, 20 mL) with magnetic stirring for 30 min. Na2SeSO3: Synthesized by refluxing Se powder (5 mmol) and Na2SO3 (6 mmol) in deionized water (50 mL) at 80 °C for 3 h. NaHTe: Generated via NaBH4 reduction (2.5 g in 5 mL H2O) of Te powder (0.127 g) in a Schlenk flask under N2 flow (40 °C, 20 min vigorous stirring) and stored in a glove box. Desired M-NG was activated by vacuum at 100 °C for 1 h to remove adsorbed gas and moisture.

The synthesis of QD-functionalized glass. Perovskite QDs: CsPbX3 (X=Cl, Br, I): Pb-M-NG immersed in methanolic CsX solutions (60 °C, 2 h). MAPbBr3/FAPbBr3: Pb-M-NG immersed in methanolic MA/FABr solutions (60 °C, 2 h). Chalcogenide QDs: Cd-based systems: CdS: Cd-M-NG immersed in 0.1 M Na2S (Tris–HCl buffer, pH 7.4) for 1 h. CdSe: Cd-M-NG reacted with 0.08 M Na2SeSO3 for 1 h. CdTe: Cd-M-NG treated with freshly prepared NaHTe (0.05 M in degassed H2O) for 1 h in a glove box. Pb-based systems: PbS: Pb-M-NG in Na2S solution (1 h, static immersion). PbSe: Pb-M-NG in Na2SeSO3 solution for 1 h. Ag-based QDs: Ag2S: Ag-M-NG in Na2S solution (1 h, dark condition). AgInS2: In/Ag-M-NG (1 molar ratio) reacted with Na2S solution for 1 h. Zn-based QDs: ZnS: Zn-M-NG in Na2S solution for 1 h. After QD growth, vacuum drying was performed at 50 °C to remove the solvent.

Preparation of colloidal CsPbBr3 QDs. Colloidal CsPbBr3 QDs were prepared using the hot-injection method37. Cs oleate precursors were prepared by reacting Cs2CO3 with oleic acid in 1-octadecene. The mixture was degassed at 130 °C for 1 h in a three-necked flask, then heated to 150 °C under nitrogen. Then, PbBr2 precursor solution (PbBr2, oleylamine, and oleic acid mixed in 1-octadecene) was dried under vacuum at 120 °C for 1 h in a three-necked flask. The temperature was raised to 160 °C, followed by the rapid injection of Cs oleate precursor. The reaction was quenched with an ice water bath, and CsPbBr3 QDs were collected by centrifugation (9056×g, 0 °C), washed three times with hexane and methyl acetate, and dispersed in toluene (~100 mg/mL). 50 μL of QD solution was drop-cast onto a glass slide and dried at 60 °C until the solvent evaporated, forming a uniform film.

In-situ absorption and PL measurements
The reactions were carried out in a modified quartz cuvette. The absorption spectra were recorded with an Ocean Optics deuterium–tungsten light source (DH2000-BAL-TTL-24V) and a Thorlabs CCS200/M spectrometer. For in-situ PL and PL mapping measurements, the excitation beam was introduced to a confocal system and focused on the samples through an objective (×10). The PL signals were collected by the objective and detected by a spectrograph with 1200 mm−1 grating and a silicon-charged coupled resolution of 0.04 nm.

Photocatalytic CO2 reduction tests
The photocatalytic CO2 reduction performance was explored using Labsolar-6A system (Perfect Light Co., Ltd., China) and a gas chromatograph. A 300 W Xe lamp was used as the photocatalytic light source to simulate sunlight. QD glass samples were added to the reaction chamber. Before the test, the system was dehydrated to remove O2 and other gases, and then CO2 was introduced. The content of Cs element in the glass matrix was determined by ICP testing, thus obtaining the content of CsPbBr3 QDs.

Small and wide-angle X-ray scattering (SAXS/WAXS) measurements
The SAXS experiment was conducted at beamline BL19U2, and the WAXS experiment was carried out at BL17B1 in Shanghai Synchrotron Radiation Facility (SSRF). For in-situ SAXS, X-ray photon energies were 12 keV with a Pilatus 2M detector. Initially, the M-NG was confined within a cylindrical reactor. To eliminate the influence of methanol diffusion within the M-NG pores during the reaction, the M-NG was pre-soaked in methanol to ensure complete infiltration of the nanopores. Following this pre-treatment, an excess amount of precursor solution was introduced by an electrically driven automatic injector to initiate the reaction. For WAXS, X-ray photon energies were 10 keV with a Pilatus 2M detector. The q-range and the scattering intensity were calibrated with LaB6.

Morphology and composition measurement
N2 adsorption–desorption isotherms were measured at 77 K using a QuantaChrome Autosorb iQ-MP gas sorption analyzer. Specific surface area was calculated using the Brunauer–Emmett–Teller (BET) method in the relative pressure range of 0.1–0.3. Pore size distribution was derived from the desorption branch using the non-local density functional theory (NLDFT) model. The refractive index was measured using a GAERTNER L116 ellipsometer. High-resolution TEM and HAADF-STEM imaging were performed using a Thermo Fisher Spectra Ultra scanning transmission electron microscope. Scanning electron microscopy (SEM) images were acquired using a ZEISS Gemini 500 field-emission SEM. Atomic force microscopy (AFM) was performed using a Bruker Dimension Icon system in tapping mode. Elemental analysis was performed using an Agilent 5110 ICP-OES system. XAS measurements at the Pb L-edge were performed at the BL16U1 and BL17B1 in SSRF. Data was collected from 100 eV below the Pb L-edge to ∼800 eV above the edge.

Computational methods
The dense glass matrix was modeled using classical molecular dynamics (MD) simulations with a Morse potential38,39. Long-range Coulombic interactions were computed using Ewald summation with a 10 Å cutoff. A simulation box containing 17,997 atoms and periodic boundary conditions (PBC) in all directions was employed. The system was heated to 3000 K over 0.57 ns, equilibrated at this temperature for 1.0 ns, then cooled to 300 K over 0.57 ns and further equilibrated at 300 K for 0.2 ns under the NPT ensemble (0 bar), using LAMMPS40. This process yielded a dense aluminosilicate glass structure with a cubic box size of 63.3 Å×63.3 Å×63.3 Å. Eight idealized spherical cavities with a diameter of 2.8 nm were introduced at the center of the box by removing the corresponding glass atoms. The porosity and specific surface area were designed based on experimental Brunauer–Emmett–Teller (BET) analysis.

Theoretical methodologies based on DFT calculations were performed using the Vienna ab-initio simulation package (VASP)41–43. The generalized gradient approximation (GGA) with Perdew–Burke–Ernzerhof (PBE) function was used as the exchange-correlation function41. The projected augmented wave (PAW) method was used to describe the interactions between ions and electrons44. Grimme’s D3 correction was used to consider the van der Waals interactions45. Monkhorst–Pack mesh and kinetic energy cutoff were set to be a grid of (3×3×1) k-points and 500 eV, respectively44. For the computation of adhesive energy, the bottom atoms of the SiO2 substrate were constrained to maintain the bulk environment. To evaluate the stability of the CsPbBr3/SiO2 interfaces, we calculate the interfacial binding energy (Eb), which is defined as
Eb = Einterface − ESio2 − Eperovskite (1)
where Einterface, ESio2, and Eperovskite denote the total energies of the heterointerface supercell, the SiO2 substrate, and the perovskite slab.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The study presents an interesting result of 3D printed optical glasses. Through functionalized by QDs, the 3D printed glass with full spectrum luminescence across UV–VIS–IR were realized through doping nanoporous glass with desired metal ion precursors. This work addresses critical challenge in the field of 3D printed glass, namely the conflict between high-temperature sintering process required in 3D printed glass and the combination of QDs without degeneration. The 3D printed functionalized glasses were constructed relying on the spatial and chemical dual nano-confined growth of diverse QDs under low-temperature, which is innovative and effective. Furthermore, this low-temperature combined strategy for 3D printed glass is universal, which is applicable for various semiconductors QDs. These results would open the door of 3D printed glass for complex, spectrally-engineered applications. I believe this work is a valuable contribution to the field of 3D printed glass, and recommend its publication in Nature Communications after minor revisions to address the points below.\n1. The quantitative atomic ratios of silicon, oxygen, and any potential contaminants (such as carbon) should be provided to demonstrate the amorphous silica has a 1:2 stoichiometric ratio and no contaminants. In addition, it is recommended that Raman spectroscopy characterization should be presented to further confirm the glass composition.\n2. The size of the QDs is controlled by tunning the pore size of the nanoporous glass, which is adjusted via the aluminum lactate concentration. However, the manuscript lacks direct characterization of the different pore size. The authors should provide pore size distribution data for the corresponding samples.\n3. In the analysis of the relationship between precursor concentration and reaction time (Figure 2h), the authors draw a conclusion based on a comparison between a theoretical model and experimental data. However, only four experimental points are presented. The authors should provide more data points to demonstrate the statistical reliability of their claims.\n4. Can the concentration of the QDs be quantitatively controlled? The manuscript briefly discusses a relationship between QD concentration and the precursor doping level. More quantitative analysis of the final QD concentration and its controllability should be provided.\n5. To better highlight the method’s versatility, please consider including the organic-inorganic hybrid perovskite QDs in Figure 1d.\n6. In Figure 1a, the capitalization of initial letters in the labels should be made consistent. In Figure 2, the y-axis for “QD concentration” is presented without specific units or values. In Figure 3, the labels appear to be misaligned with their corresponding diffraction peaks. Please carefully check and correct to ensure the accuracy of the figures.", "reviewer_2": "This manuscript reportes the 3D printing of functional glass materials: a low-temperature post-functionalization strategy for nanoporous glass (NG). This strategy addresses the critical contradiction of \"high-temperature sintering-induced quantum dot (QD) degradation\" that arises during the integration of QDs with 3D-printed glass. Leveraging the intrinsic properties of nanoporous glass, the authors propose an approach to realize the in-situ growth of diverse QDs through spatial and chemical dual nano-confinement effects. Furthermore, the manuscript employs a comprehensive set of characterization techniques to fully validate the underlying mechanism and provides compelling demonstrations of potential applications. I recommend this work for publication in Nature communications after the authors address the following issues.\n1. The authors did not provide any refractive index data for nanoporous glass, metal-ion doped NG (M-NG), and 3D-printed QDs-functionalized glass. Since refractive index is critical for the optical behaviors (e.g., light coupling, waveguiding) vital to the proposed optical applications, these data should be given.\n2. When describing the PL stability of PQD (perovskite QD) glass, the manuscript refers to \"ambient condition\" (Fig. 3i) but does not define its specific parameters-especially relative humidity (RH), a key factor which influences the stability of perovskite. What is the RH level of the \"ambient condition\"? In addition, can the PQD glass maintain its PL intensity and structural integrity under high-humidity environments (e.g., 80% RH or 95% RH)?\n3. In the photocatalytic CO2 reduction experiments (Fig. 4), the authors attribute enhanced performance of complex 3D structures (e.g., channel-like microarray) to \"enhanced photon scattering and surface area.\" However, they do not distinguish whether the improved activity stems from the porous nature of the glass matrix (e.g., mass transfer of CO2/O2 through pores) or solely from the increased surface area of the 3D architecture.\n4. The manuscript claims \"voxel-level control over the QD bandgap\" but does not specify the voxel size of the 3D-printed structures (e.g., lateral resolution × layer thickness) or how QD uniformity is verified across individual voxels. What is the voxel dimension achieved via DLP printing?\n5. For the in-situ QD growth process, the authors mentioned that M-NG is \"saturated in QDs precursor solution\" but did not report the influence of precursor concentration on the QD growth and the light emitting performance. Will the increase of precursor concentration lead to excessive QD aggregation, and if so, what is the optimal concentration window to maintain monodisperse QDs?\n6. There is a spelling error in Line 102 of Page 3 in the manuscript, which should be corrected to Zn(NO3)2."}, "author_response": "We thank all reviewers for the positive comments and valuable suggestions. In response to these comments, we have made the following corrections and improvements point by point. The manuscript has been carefully checked to ensure clarity and grammatical accuracy.\nReviewer #1 (Remarks to the Author):\nThe study presents an interesting result of 3D printed optical glasses. Through functionalized by QDs, the 3D printed glass with full spectrum luminescence across UV–VIS–IR were realized through doping nanoporous glass with desired metal ion precursors. This work addresses critical challenge in the field of 3D printed glass, namely the conflict between high-temperature sintering process required in 3D printed glass and the combination of QDs without degeneration. The 3D printed functionalized glasses were constructed relying on the spatial and chemical dual nano-confined growth of diverse QDs under low-temperature, which is innovative and effective. Furthermore, this low-temperature combined strategy for 3D printed glass is universal, which is applicable for various semiconductors QDs. These results would open the door of 3D printed glass for complex, spectrally-engineered applications. I believe this work is a valuable contribution to the field of 3D printed glass, and recommend its publication in Nature Communications after minor revisions to address the points below.\nResponse: We appreciate the reviewer for the thoughtful and positive evaluation of our work.\n1. The quantitative atomic ratios of silicon, oxygen, and any potential contaminants (such as carbon) should be provided to demonstrate the amorphous silica has a 1:2 stoichiometric ratio and no contaminants. In addition, it is recommended that Raman spectroscopy characterization should be presented to further confirm the glass composition.\nResponse: We thank the reviewer for the constructive comments. We have performed SEM-EDS mapping analysis on the printed NG and conducted quantitative analysis of the elements (Si, O, and Al) using EDS ranging from 0 to 10 keV in Fig. R1. The measured atomic ratios of Si, O, and Al are in agreement with the designed SiO2-Al2O3 composition, with no detectable carbon contamination. We have added this figure as in the revised SI (Supplementary Fig. 2).\nAccording to the suggestion, we have performed Raman spectroscopy of the 3D printed NG before and after 650°C sintering (Fig. R2, Supplementary Fig. 3 in revised SI). The spectra of the organic peaks (C-H, C=O), which were significant before sintering, completely disappeared after sintering at 650°C. In the Raman spectrum of NG after sintering, the ω1 and ω3 bands can be observed, which correspond to the bending vibrations of the Si-O-Si bridges in Si(O1/2)4 tetrahedra. Moreover, the D1 and D2 peaks correspond to the symmetric stretching of the Si-O rings are observed [Phys. Rev. B 1983, 28, 3266-3271]. An additional peak at 953 cm-1 between the peaks indicates the presence of non-bridging oxygen (NBO) in the structure, which results from doping with metal elements in the NG. These results demonstrate the glass composition primarily consists of Si-O-Si networks, which is similar with the commercial fused silica. The corresponding description has been added in the revised manuscript (Page 4, Line 120-125) and revised SI.\nFigure R1. (a) SEM image and (b-d) elemental mappings of O, Al, and Si, respectively. (e) EDS spectrum and quantified atomic percentages, consistent with the designed Si-O-Al composition (Al2O3 : SiO2 = 4: 92) and confirming the absence of carbon contamination.\nFigure R2. Raman spectra of printed green body, NG after 650°C sintering and commercial fused silica. The spectra of the printed green body show significant organic peaks (C-H, C=O), which completely disappear after sintering at 650°C. The spectra of NG after 650°C sintering is similar with the characteristic of commercial fused silica, indicating the complete composition transformation of the printed gel into SiO2 without organic residual.\n2. The size of the QDs is controlled by tunning the pore size of the nanoporous glass, which is adjusted via the aluminum lactate concentration. However, the manuscript lacks direct characterization of the different pore size. The authors should provide pore size distribution data for the corresponding samples.\nResponse: We thank the reviewer for this valuable suggestion. According to the reviewer’s comment, N2 adsorption/desorption measurements have been performed on NG with different Al contents. As shown in Fig. R3 (Supplementary Fig. 23), the pore diameter gradually increased from ~2.5 to ~7.1 nm as the Al content rose from 2% to 25%. These results confirm that the pore size of the M-NG can be effectively tuned by adjusting the content of aluminum lactate, which act as nanopore former in sol-gel derived Al2O3-SiO2 NG. These results and discussion have been added in the revised manuscript (Page 5, Line 153-155) and revised SI.\nFigure R3. N2 adsorption/desorption isotherms of the NG with different Al:Si ratios of (a) 2:98, (b) 12:88, (c) 15:85, (d) 20:80, and (e) 25:75. Insert show corresponding pore size distributions. (f) The relationship between pore diameter and Al content, indicating that increasing Al concentration leads to enlarged pores.\n3. In the analysis of the relationship between precursor concentration and reaction time (Figure 2h), the authors draw a conclusion based on a comparison between a theoretical model and experimental data. However, only four experimental points are presented. The authors should provide more data points to demonstrate the statistical reliability of their claims.\nResponse: We thank the reviewer for the constructive comment. We have done corresponding experiment at more CsBr concentrations of 6, 15, and 17 mg/mL as shown in Fig. R4 (Figs. 2f and 2h in revised manuscript). All the experimental data show good agreement with the theoretical prediction.\nFigure R4. In-situ recorded absorbance of QDs at (f) various CsBr concentration showing QD concentration variation during the reaction. (h) Theoretical predictions and experimental observations regarding the dependence of the reaction saturation time on the CsBr concentration.\n4. Can the concentration of the QDs be quantitatively controlled? The manuscript briefly discusses a relationship between QD concentration and the precursor doping level. More quantitative analysis of the final QD concentration and its controllability should be provided.\nResponse: We thank the reviewer for the valuable suggestion. We have calculated the CsPbBr3 QD concentration (in molarity) using the Beer-Lambert law [Science, 2022, 377: 1406; J. Phys. Chem. Lett., 2018, 9, 3093],\nA = εcL\nwhere A is the absorbance at 400 nm, c is the concentration of QDs, L is the optical path length which is the thickness of the QD glass. ε is the extinction coefficient per mole of particles, which can be obtained from the intrinsic absorption coefficient μi,\nε = Ai · π(d/2)3 / ln10\nwhere d is average particle diameter of the QDs, μi = 7.7×104 cm-1 [J. Phys. Chem. Lett., 2018, 9, 3093].\nWe establish a relationship between the Pb content and the obtained QD concentration in Fig. R5. The results show a linear correlation between the Pb content and QD concentration, suggesting that the QD concentration in 3D-printed glass can be precisely and quantitatively controlled by adjusting the Pb doping content.\nFigure R5. The relationship between Pb content and QD concentration. Linear fitting results reveal the concentration of CsPbBr3 QDs can be quantitatively controlled by adjusting the Pb doping content.\n5. To better highlight the method’s versatility, please consider including the organic-inorganic hybrid perovskite QDs in Figure 1d.\nResponse: Thanks for this valuable suggestion. We have added the photoluminescence (PL) spectra of organic-inorganic hybrid perovskite QDs (MAPbBr3 and FAPbBr3) in Fig. 1d in revised manuscript.\n6. In Figure 1a, the capitalization of initial letters in the labels should be made consistent. In Figure 2, the y-axis for “QD concentration” is presented without specific units or values. In Figure 3, the labels appear to be misaligned with their corresponding diffraction peaks. Please carefully check and correct to ensure the accuracy of the figures.\nResponse: We thank the reviewer for the careful review. We have made the corresponding corrections in the revised manuscript.\nReviewer #2 (Remarks to the Author):\nThis manuscript reports the 3D printing of functional glass materials: a low-temperature post-functionalization strategy for nanoporous glass (NG). This strategy addresses the critical contradiction of \"high-temperature sintering-induced quantum dot (QD) degradation\" that arises during the integration of QDs with 3D-printed glass. Leveraging the intrinsic properties of nanoporous glass, the authors propose an approach to realize the in-situ growth of diverse QDs through spatial and chemical dual nano-confinement effects. Furthermore, the manuscript employs a comprehensive set of characterization techniques to fully validate the underlying mechanism and provides compelling demonstrations of potential applications. I recommend this work for publication in Nature communications after the authors address the following issues.\nResponse: We sincerely appreciate the reviewer’s positive comments and constructive suggestions. We have made the following revisions point by point according to these comments.\n1. The authors did not provide any refractive index data for nanoporous glass, metal-ion doped NG (M-NG), and 3D-printed QDs-functionalized glass. Since refractive index is critical for the optical behaviors (e.g., light coupling, waveguiding) vital to the proposed optical applications, these data should be given.\nResponse: We thank the reviewer for this valuable comments. We have measured the corresponding refractive index according to your suggestion, as shown in Fig. R6 (Supplementary Fig. 7). In the wavelength range of 300-1600 nm, pure NG exhibits a refractive index of 1.351-1.371, which is slightly lower than that of commercial fused silica glass (~1.45 in the visible range) due to its nanoporous structure. After doping with Pb, the refractive index of M-NG (Pb-NG) is increased to 1.364-1.381 due to the high electronic polarizability of Pb. The refractive index of 3D-printed QDs-functionalized glass is further increased to 1.367-1.401 after QD growth. These results and discussion have been added in the revised manuscript (Page 4, Line 128-129) and revised SI.\nFigure R6. Refractive index of NG (Al: Si = 12:88), M-NG (Pb: Al: Si = 2: 12:86), and 3D-printed QDs-functionalized glass (CsPbBr3 QDs), respectively.\n2. When describing the PL stability of PQD (perovskite QD) glass, the manuscript refers to \"ambient condition\" (Fig. 3i) but does not define its specific parameters-especially relative humidity (RH), a key factor which influences the stability of perovskite. What is the RH level of the \"ambient condition\"? In addition, can the PQD glass maintain its PL intensity and structural integrity under high-humidity environments (e.g., 80% RH or 95% RH)?\nResponse: We sincerely thank the reviewer for the valuable comments. We have explicitly defined the “ambient condition” as 25°C and 50% relative humidity (RH) in the revised manuscript. Additionally, we have conducted PL stability tests under high-humidity conditions of 80% RH, as shown in Fig. R7a (Supplementary Fig. 41a). PQD glass retains 74% of its initial PL intensity after 600 hours at 80% RH, while pure QDs only retain 15% after 117 hours. The significant improvement of stability can be ascribed to the protective glass matrix. The insets in Fig. R7a show the pictures of PQD glass before and after 600 hours, demonstrate the PQD glass maintains its structural integrity. Correspondingly, the XRD patterns of PQD glass can maintain the perovskite structure after 600 hours at 80% RH (Fig. R7b, Supplementary Fig. 41b). These results confirm the excellent stability of PQD glass even under high-humidity environments. These results and discussion have been added in the revised manuscript (Page 9, Line 299-303) and revised supporting information.\nFigure R7. (a) Time-dependent PL intensity of the PQD glass and pure CsPbBr3 QDs under 80% RH. The PQD glass maintains 74% of its initial PL after 600 h, whereas the pure QDs sharply degrade to 15% after 117 h. Insets: photographs of the PQD glass before and after the 600 hours test. (b) XRD patterns of PQD glass under 80% RH.\n3. In the photocatalytic CO2 reduction experiments (Fig. 4), the authors attribute enhanced performance of complex 3D structures (e.g., channel-like microarray) to \"enhanced photon scattering and surface area.\" However, they do not distinguish whether the improved activity stems from the porous nature of the glass matrix (e.g., mass transfer of CO2/O2 through pores) or solely from the increased surface area of the 3D architecture.\nResponse: We thank the reviewer for this insightful comment. We agree that both the porous nature [Nature Catalysis, 2020, 3, 921] and increased surface area of the 3D architecture of the glass matrix could enhance the catalytic performance. However, in our work, all the porous nature of the complex 3D structures are identical, which is confirmed by the N2 adsorption/desorption measurements on the 3D printed QD functionalized glass with different architectures (planar, #1, #2, #3) in Fig. R8. Therefore, the enhanced photocatalytic performance is due to the increased surface area of the 3D architecture, rather than the porous nature.\nFigure R8. (a) N2 adsorption isotherms at 77 K, and (b) pore size distribution for 3D printed QD functionalized glass with different 3D architectures: planar, #1, #2, and #3.\n4. The manuscript claims \"voxel-level control over the QD bandgap\" but does not specify the voxel size of the 3D-printed structures (e.g., lateral resolution × layer thickness) or how QD uniformity is verified across individual voxels. What is the voxel dimension achieved via DLP printing?\nResponse: We thank the reviewer for the insightful comment. We printed voxel arrays on a quartz substrate using a DLP printer. The voxel size realized for the printed QD functionalized glass is 180.0 × 180.0 × 5.6 μm (Figs. R9a and 9b, Supplementary Figs. 10a and 10b). For the QD uniformity in individual voxels, we performed PL mapping on a single voxel, as shown in Figs. R9c and 9d (Supplementary Figs. 10c and 10d). The results demonstrate the uniform distribution of QDs within the voxels. These results and discussion have been added in the revised manuscript (Page 4, Line 138-139) and revised SI.\nFigure R9. (a) SEM images of top-view and (b) side-view of the printed QD functionalized glass voxel array, showing an edge length of 180.0 μm and vertical thickness of 5.6 μm. These results demonstrate that the voxel dimension is 180.0 μm × 180.0 μm × 5.6 μm. (c) Optical image and (d) PL mapping of the printed PQD glass voxel, confirming the homogeneous distribution of QDs within the voxel.\n5. For the in-situ QD growth process, the authors mentioned that M-NG is \"saturated in QDs precursor solution\" but did not report the influence of precursor concentration on the QD growth and the light emitting performance. Will the increase of precursor concentration lead to excessive QD aggregation, and if so, what is the optimal concentration window to maintain monodisperse QDs?\nResponse: We thank the reviewer for this insightful comment. To investigate the influence of precursor concentration on the QD growth, we have performed TEM of the 3D printed QD functionalized glass synthesized by precursor solution with different concentrations (5, 10, and 15 mg/mL). As shown in Figs. R10a-c, the QDs show excellent monodispersity without observable aggregation. The corresponding size distribution (Fig. R10d) shows the QD sizes remain consistent across these different concentrations. The monodispersity of QDs is attributed to the uniform dispersion of Pb precursors and the spatial nanoconfinement effect provided by the M-NG matrix. The former ensures a homogeneous supply of Pb precursors, and the latter, where the pore size physically defines the QD dimensions, effectively prevents aggregation and renders QD growth independent of precursor concentration variations.\nAdditionally, we conducted corresponding PL measurements to demonstrate the light emitting performance of QD functionalized glass with different concentration precursor solutions. The PL spectra in Fig. R11 show that there is almost no difference in the emission across the three precursor concentrations, indicating that precursor concentration does not have a noticeable effect on the light emitting performance of the QDs. Even at the high precursor concentration of 15 mg/mL, the emission peak remains at 490 nm, indicating the monodispersed quantum-confined QDs rather than aggregated ones (for CsPbBr3 bulk material, the emission peak is around 534 nm), in agreement with the TEM results. We observed the highest PLQY of 82% at a CsBr concentration of 5 mg/mL, as shown in Fig. R11b., indicating this is the optimal concentration.\nFigure R10. TEM images of the 3D printed QD functionalized glass synthesized by precursor solution with different concentrations: (a) 5 mg/mL, (b) 10 mg/mL and (c) 15 mg/mL of CsBr solution. (d) QDs size distribution for different precursor concentrations.\nFigure R11. (a) Normalized PL spectra and (b) PLQY of 3D printed QD functionalized glass synthesized by precursor solution with different concentrations.\n6. There is a spelling error in Line 102 of Page 3 in the manuscript, which should be corrected to Zn(NO3)2.\nResponse: We thank the reviewer for their careful review. We have corrected this spelling error."}, {"round": 2, "reviewer_comments": {"reviewer_1": "As the authors have addressed all my concerns, I suggest this work can be accepted for publication by Nature Communications.", "reviewer_2": "The author has addressed all my concern, and I believe the manuscript is now ready for publication."}, "author_response": ""}]

Decision letter