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Non-Equilibrium Synthesis of TiO Nanoparticle “Building Blocks” for Crystal Growth by Sequential Attachment in Pulsed Laser Deposition Masoud Mahjouri-Samani, Mengkun Tian, Alexander A. Puretzky, Miaofang Chi, Kai Wang, Gerd Duscher, Christopher M. Rouleau, Gyula Eres, Mina Yoon, John C. Lasseter, Kai Xiao, and David B. Geohegan Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b01047 • Publication Date (Web): 10 Jul 2017 Downloaded from http://pubs.acs.org on July 12, 2017

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Non-Equilibrium Synthesis of TiO2 Nanoparticle “Building Blocks” for Crystal Growth by Sequential Attachment in Pulsed Laser Deposition Masoud Mahjouri-Samani1*, Mengkun Tian2, Alexander A. Puretzky1, Miaofang Chi1, Kai Wang1, Gerd Duscher2,3, Christopher. M. Rouleau1, Gyula Eres3, Mina Yoon1, John Lasseter1, Kai Xiao1, David B. Geohegan1* 1. 2. 3.

Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA Dept. of Materials Science and Engineering, University of Tennessee, Knoxville, TN, 37966, USA Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA

*M. Mahjouri-Samani, [email protected] *D. B. Geohegan, [email protected]

Notice: This manuscript has been authored by UT-Battelle, LLC, under Contract No. DE-AC0500OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, world-wide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes. The Department of Energy will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan (http://energy.gov/downloads/doe-public-access-plan).

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Non-Equilibrium Synthesis of TiO2 Nanoparticle “Building Blocks” for Crystal Growth by Sequential Attachment in Pulsed Laser Deposition Masoud Mahjouri-Samani1*, Mengkun Tian2, Alexander A. Puretzky1, Miaofang Chi1, Kai Wang1, Gerd Duscher2,3, Christopher. M. Rouleau1, Gyula Eres3, Mina Yoon1, John Lasseter1, Kai Xiao1, David B. Geohegan1* 1. 2. 3.

Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA Dept. of Materials Science and Engineering, University of Tennessee, Knoxville, TN, 37966, USA Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN,37831, USA

*M. Mahjouri-Samani, [email protected] *D. B. Geohegan, [email protected]

Non-equilibrium growth pathways for crystalline nanostructures with metastable phases are demonstrated through the gas-phase formation, attachment, and crystallization of ultrasmall amorphous nanoparticles as building blocks in pulsed laser deposition (PLD). Temporally- and spatially-resolved gated-ICCD imaging and ion probe measurements are employed as in situ diagnostics to understand and control the plume expansion conditions for the synthesis of nearly pure fluxes of ultrasmall (~3 nm) amorphous TiO2 nanoparticles in background gases and their selective delivery to substrates. These amorphous nanoparticles assemble into loose, mesoporous assemblies on substrates at room temperature, but dynamically crystallize by sequential particle attachment at higher substrate temperatures to grow nanostructures with different phases and morphologies. Molecular dynamics calculations are used to simulate and understand the crystallization dynamics. This work demonstrates that non-equilibrium crystallization by particle attachment of metastable ultrasmall nanoscale “building blocks” provides a versatile approach for exploring and controlling the growth of nanoarchitectures with desirable crystalline phases and morphologies.

Keywords: Crystallization by particle attachment, nanoparticle building blocks, pulsed laser deposition, molecular dynamics simulation.


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Introduction Crystallization by particle attachment (CPA)1-2 is a fascinating mechanism for crystal growth where nanoscale amorphous or crystalline nanoparticles act as building blocks to form larger crystals by thermodynamically minimizing their collective free energy.3-5 The mechanisms by which nanoparticles build larger crystalline nanostructures by CPA have been investigated primarily by in situ transmission electron microscopy (TEM) of nanoparticles in liquids, and include oriented attachment, recrystallization, or amorphous addition.1-5

The growth

mechanisms of crystalline nanostructures on substrates resulting from the physical vapor deposition of nanoparticles formed in the gas phase are far less studied. Recently, we showed that ‘amorphous’ TiO2 nanoparticles synthesized by condensation in pulsed laser vaporization plumes are thermodynamically metastable precursor particles which can transform into crystalline TiO2 nanostructures with novel phases, such as ‘black TiO2’ core-shell nanoparticles, depending on their post-annealing treatment.6 This phenomenon follows the Ostwald−Lussac law of phases, which states that the pathway to a final crystalline state evolves by passing through increasingly stable states.7-11 Here we utilize the highly non-equilibrium nature of pulsed laser ablation plasma thermalization in background gases to controllably condense and deliver nearly pure beams of these ultrasmall amorphous precursor nanoparticles to explore their crystallization by particle attachment into nanostructures with crystalline phases of different metastability and morphology. Pulsed laser deposition (PLD) is typically employed as a synthesis tool for epitaxial thin film growth primarily for its advantages of stoichiometric transfer of material from target to substrate with non-thermal (~100 eV) kinetic energies, if desired.12-13 For the growth of thin films, background gas pressures during PLD are typically kept below ~10 mTorr in order to maintain the flux as primarily atoms, ions, and molecules to facilitate the layer-by-layer growth of epitaxial thin films with control rivaling molecular beam epitaxy. However, when the background gas pressure is high enough to scatter and confine a sufficient density of the atoms and molecules in the ablation plasma plume, clusters and nanoparticles can form by condensation. When sufficient time and temperature are afforded (e.g., as in a hot oven environment), time-resolved in situ laser spectroscopy and imaging diagnostics have revealed how these nanoparticles and clusters can self-assemble in the gas-phase into larger 3 ACS Paragon Plus Environment

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nanostructures, such as single-wall carbon nanotubes from condensed Ni/Co and carbon nanoparticles.9, 14-19 Alternatively, nanoparticles that have just formed in the ablation plume can be deposited at room temperature to construct mesoporous nanoparticle films for the fabrication of hierarchical electrode architectures (primarily of TiO2, but also WO3, ITO, etc.) for dyesensitized solar cells,20-23 reflective coatings,24 photocatalysis,25-26 and other applications.24, 27-30 In these cases, the nanoparticle films are typically post-annealed to form sintered, hyperbranched, mesoporous architectures of polycrystalline nanowires. However, a few reports have shown that nanoparticle aggregates formed at relatively high pressures (e.g. 10 Torr), in pulsed laser ablation plumes can be delivered as the feedstock for the direct synthesis of crystalline nanostructures on hot substrates, such as single-crystal ZnO nanorods, in a process termed nanoparticle-assisted PLD (NAPLD).31-33 Here, with time-resolved in situ diagnostics, high-resolution electron microscopy observations of nanoparticle crystallization and sintering, and molecular dynamics modeling, we reveal the nonequilibrium conditions for the gas-phase synthesis and sequential PLD of ultrasmall amorphous nanoparticles, and their crystallization by particle attachment at elevated temperatures for the direct formation of crystalline nanostructures. We use TiO2 as a well-known model system and demonstrate how interesting crystalline phases, such as TiO2(B) can be directly synthesized. The pulsed laser deposition of amorphous ultrasmall nanoparticles and their dynamic crystallization by particle attachment demonstrated here is a promising technique to explore the synthesis of nanostructured films with novel phases and morphologies determined by the competition between kinetics and thermodynamics. Results and Discussion Pulsed laser vaporization of a solid target of TiO2 in low-pressure background gases was used to synthesize ultrasmall (2-5 nm-diameter) amorphous nanoparticles of TiO2 by gas phase condensation, under conditions similar to those described recently in our previous work.6 These amorphous ultrasmall nanoparticles (UNPs) can serve as precursors to transform during annealing treatments into larger crystalline nanostructures, such as the core-shell ‘black-TiO2” nanoparticles in that study.6 Here we employ time-resolved in situ diagnostics to adjust the thermalization dynamics of the ablation plume for different background gas pressures and deposition distances in order to not only synthesize gas-phase condensed nanoparticles, but also 4 ACS Paragon Plus Environment

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deposit them selectively on heated substrates in order to explore the dynamic crystallization of larger TiO2 nanostructures.

Figure 1. In situ diagnostics of plume evolution as a function of pressure. (a) In situ ICCD images of the laser ablation plume luminescence at the indicated arrival times of the plume front at a room-temperature substrate 5cm away from the target for different background oxygen pressures (20 ns to 1s exposures, false color, maximum intensity relative to vacuum image provided). (b) In situ ion probe voltage (across 50 Ω resistor) at d = 5 cm as a function of background gas pressure, showing plume splitting. For pressures < 50mTorr the fast component dominates, while two components are evident from 50-100 mTorr, and the slow component dominates for P > 100 mTorr. The 200mTorr curve utilized a 500 Ω resistor and is scaled by 10X to show a weak 3 rd component of the plume. (c-g) Cross-sectional SEM images showing the morphology of the samples deposited at d = 5 cm for various pressures, with dense films transforming to nearly pure nanoparticle architectures after the second component in the ion probe voltage disappears at 200mTorr.

In situ diagnostics, including gated intensified charge couple device (ICCD) array imaging, ion probe time-of-flight current measurements, and laser-induced fluorescence spectroscopy, were employed to adjust the plume conditions for the onset of nanoparticle formation according to previous studies.16, 34 As shown by the ICCD images and ion probe waveforms in Figure 1a in vacuum, the KrF-laser ablation (248nm, 1 J cm-2, 1 Hz) of TiO2 results in a rapidly expanding, forward-directed plasma plume containing atoms, ions, and molecules ejected from the target with high kinetic energies (up to 100 eV for the leading edge of the plasma plume traveling at 2 cm/µs velocity). With increasing background O2 pressure the arrival time of the leading edge of the visible plasma plume to the indicated 5 cm position increases to longer delays, as the plume 5 ACS Paragon Plus Environment

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becomes increasingly more spatially confined due to collisions between the ejected target atoms, ions, and molecules and those of the background gas. These collisions snowplow the background gas molecules forward, forming a weak ‘shock front’ that is visible due to the bright fluorescence resulting from the collisional processes within the confined plasma.35-36 However, a fast, ‘dark’ component of the plume not visible in the images – high kinetic energy atoms and ions that have penetrated the background gas without collisions – is also present during some of these conditions, as revealed by the ion probe current waveforms in Figure 1b. The saturated ion probe current (measured by voltage across 50 ohms, detector area = 1 mm2) waveforms represent a measure of the relative flux (N𝜐) of positive ions penetrating through the background gas at each time. The magnitude of the peak current (and also the overall integrated charge) arriving to the ion probe drops exponentially with pressure at a given distance (and drops exponentially with distance at a given pressure) in accordance with a scattering model with a hard sphere collision cross section,34-35,

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which for the data of Figure 1b yields the value of

1.7×10-16 cm2. As shown in Figure 1b, this ‘fast’ component of the plume arrives with nearly the identical velocity distribution as in vacuum, and eventually separates from the time-of-flight velocity distribution that corresponds to the visible, ‘slowed’ component of the plume. This ‘plume splitting’ phenomenon has been explained in the context of a scattering model, which describes the shape of the two waveforms as composite distributions of target atoms undergoing discrete numbers of collisions enroute to the substrate.34, 37 As the pressure increases in Figure 1, the ion probe reveals that all of the ‘fast’ component (high kinetic energy) ions have disappeared after ~ 100 mTorr, and that between 100-200 mTorr the ions corresponding to the slowed ‘second’ component also have been attenuated to < 0.0002 of the flux arriving in vacuum, and their kinetic energies have been reduced to < 0.04 eV. Corresponding cross-sectional SEM images of the films deposited onto room temperature substrates mounted on the ion probe during these real-time diagnostic experiments are shown in Figure 1c-g. The morphology of the samples changes dramatically and continuously from dense films in vacuum to fluffy, fractal aggregates of nanoparticles at 200 mTorr. With the attenuation of the high kinetic energy species and the emergence of a comparable ‘second’ component of the ion flux in Figure 1b, the films become less dense, rough and columnar as in Figure 1d,e. After the elimination of the ‘fast’ component, and during the elimination of the ‘second’ component 6 ACS Paragon Plus Environment

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between 120-200 mTorr, the films become nearly entirely composed of aggregates of ultrasmall nanoparticles as shown in Figure 1f,g. As shown in the images of Fig. 2a and the distance-time (R-t) plot of the leading edge of the visible plume in Fig. 2b, the forward-directed progression of the visible plasma plume essentially stops after d = 5.4 cm. A linear drag model, 𝑎 = − 𝜐, that relates the plume deceleration 𝑎 to its velocity 𝜐 with a drag coefficient, results in the red dashed curve fit in Fig. 2b. However, the plume, in fact, continues to diffusively expand more slowly in all directions as fit by a more accurate 𝑎 = −𝛼𝜐 2 drag model (blue curve), where 𝜐 = 𝜐0 /(1 + 𝛼𝜐0 𝑡), 𝜐0 is the initial velocity of the visible ablation plasma, and  is the drag coefficient. Following the disappearance of the visible plasma plume, we have previously shown that nanoparticles nucleate and grow in the gas phase from this thermalized, diffusively-expanding plume material.16, 34 Emission spectroscopy measurements confirmed that the weak emission in the long-lived fluorescence observed in Figure 2a at 200 mTorr can be attributed principally to oxide bands, as we have observed previously.16 Over hundreds of microseconds to milliseconds of time, previously we utilized a time-delayed second laser sheet beam to show how ground state oxide species (via laser-induced fluorescence) disappear and nanoparticle aggregates become visible (via Rayleigh scattering) in the thermalized plume region between target and substrate.16 Here, in similar measurements at high laser repetition rates, we were able to observe Rayleigh scattering from gas-suspended aggregated nanoparticles that were organized into bands near the substrate resulting from multiple laser shots, a phenomenon that will be subject of a separate report. However, at the low repetition rates used in the deposition experiments and in single-shot ablation measurements, we were not able to observe detectable Rayleigh scattering from nanoparticle aggregates.

By

focusing the time-delayed probe laser beam (λ=355 nm) very tightly, we were able to observe the individual nanoparticles by fluorescence following their photodissociation, where we observed a shift in the intensity dependence of photoluminescence oxide molecule excitation from linear to quadratic behavior as the time delay between lasers was varied from t = 100 µs to t = 10 ms at our typical deposition distance of d = 5 cm. From these measurements we conclude that under these conditions, gas-phase conversion of the thermalized plume into ultrasmall nanoparticles is completed t > 100 µs after laser ablation, and that the nanoparticles are insufficiently aggregated to scatter detectable laser light.


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Figure 2. Plume conditions for nanoparticle formation and deposition. (a) Intensified CCD-array imaging of visible plasma plume luminescence at the indicated times following laser ablation (20 ns to 5s exposures, false color, normalized peak intensities relative to 14000 counts of 2 s image). Dashed lines at d = 0 (target) and 6 cm (substrate). (b) R-t plot of the plasma plume front (defined as the point of 5% peak intensity). The propagation is best described (blue line) by the 𝑎 = −𝛼𝜐 2 drag model, where 𝜐 = 𝜐0 /(1 + 𝛼𝜐0 𝑡); 𝑅 = 𝛼 −1 𝑙𝑛(1 + 𝛼𝜐0 𝑡), where 𝛼 = 0.91 cm-1 and 𝜐0 = 5.1 cm/s. By comparison the linear 𝑎 = −𝜐 drag model fit (red dashed curve, where R = (𝜐0 /) (1-exp (–t)) = 0.21 s-1) predicts that the plume slows to stop at R = 5.4 cm. (c,d) Cross sectional and top view SEM images of room temperature TiO2 nanoparticles deposited onto a Si substrate at 200 mT oxygen background pressure and 5 cm away from the target. (e,f) TEM images showing aggregates of (~ 3 nm) nanoparticles deposited onto the TEM grids.

These individual nanoparticles that are formed in the diffusively-expanding plume can propagate well beyond it. We find that the nanoparticles can be deposited just past the tip of the visible plume, as shown in Figure 2c, up to several centimeters further away from the target. At room temperature, the nanoparticles assemble in fractal hierarchical structures on substrates at d = 5 cm as shown in Figure 2c-d, very similar to those deposited and post-annealed by others for dyesensitized solar cell electrodes and other applications.20-30, 38 Figure 2e,f show the TEM images (HRTEM, Zeiss Libra 200 MC at an acceleration voltage of 200 kV) of these ultrasmall nanoparticles (~3 nm) formed under these synthesis conditions and deposited on a TEM grid at d = 5cm away from the target. The nanoparticles appear amorphous by selected area electron diffraction as shown in the inset of Figure 2e and exhibit compositional stoichiometry with the TiO2 target as recently reported.6 To understand the mechanism by which pure nanoparticle architectures can be formed beyond the visible plasma plume, we utilize the data of Figure 2, and estimate nanoparticle propagation dynamics using results from our recent studies. Recently, we directly imaged the propagation of much larger (d ~160 nm) Pt nanoparticles through much higher pressure (10-50 Torr) 8 ACS Paragon Plus Environment

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background gases and determined that their deceleration is linear with velocity, in accordance with Epstein drag.39-40 The propagation normal to the target followed x = (𝜐0 /)(1-exp(–t)) where the slowing coefficient, is proportional to pressure, P, and inversely proportional to nanoparticle diameter, D according to (s-1) = 114,000 P(Torr) / D(nm). Utilizing this scaling relationship, we can estimate the stopping distance xf = (𝜐0 /) for nanoparticles launched with velocity 𝜐0 in the present study. From Figure 2b, we calculate the initial nanoparticle velocity assuming that they are formed and propagating with the decelerating plasma at d = 5.4 cm, which is the commonly-accepted “plume range” based on a linear drag model (terminal distance in the red dashed curve). Using the measured velocity according to the quadratic drag model at this distance, we find that 𝜐0 = 370 m/s, from which we can estimate xf(cm)=0.33D(nm)/P(Torr). This predicts that individual 3-nm-diameter nanoparticles condensing in the final stages of plume expansion, and launched at 370 m/s into 0.2 Torr background pressure, can propagate an additional 5.0 cm past the 5.4-cm distance associated with the commonly-associated range of the plume. This explains how relatively pure fluxes of nanoparticles can be deposited beyond the visible plasma plume. This propagation mechanism for individual nanoparticles at lower pressures is quite different than the hydrodynamic propagation within swirling vortex rings at higher pressures wherein nanoparticle aggregates are trapped and can be transported over very long distances.15, 19 Hence, by positioning the substrate just outside the thermalized plasma plume, as detailed above, nearly pure deposits of ultrasmall nanoparticles can be deposited to explore their assembly as ‘building blocks’ for the formation of crystalline nanostructures by particle attachment. The ion probe and imaging diagnostics described above are simple, implementable diagnostics to arrange nanoparticle formation and deposition conditions for different laser fluences, distances, and background gas pressures. At high substrate temperatures, however, it must be noted that the temperature gradient between the heater and the target exerts a thermophoretic force which tends to oppose nanoparticle propagation to the heater.16 Hence, the deposition distance must be carefully chosen to eliminate the high kinetic-energy atomic/molecular component while preserving sufficient nanoparticle momentum to overcome thermophoresis. We found that these as-synthesized ultrasmall nanoparticles can serve as versatile “building blocks” to form a variety of nanostructures and phases when deposited at different deposition 9 ACS Paragon Plus Environment

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temperatures, and background gas pressures. According to the aforementioned understanding of pressure and distance effects on the formation and propagation of nanoparticles during pulsed laser deposition, a background oxygen pressure of 200 mTorr and substrate-to-target distance of 5 cm was chosen as suitable conditions to explore the structural evolution of nanostructured films formed by crystallization of ultrasmall nanoparticles deposited at various substrate temperatures. Figure 3 shows SEM and TEM characterization of representative samples deposited at 400, 600, and 800 oC at a repetition rate of 1 Hz in oxygen. Both the morphology and phase of the structures evolved. At 400 oC, crystalline anatase nanorods were found in roughly similar hyperbranched architectures as the room-temperature deposits of amorphous nanoparticles previously shown in Figure 3a-d. The nanorods were very narrow, with lateral dimensions of ~5-10nm, but long axial dimensions, indicating that crystallization preferentially occurred axially at this temperature, in the direction of the arriving particle flux, but that temperatures were insufficient for lateral crystallization. At higher deposition temperatures (500, 600 oC), however, denser and more crystalline architectures consisting of vertically-standing nanosheets of slightly larger lateral dimensions were found with consistently the TiO2(B) phase, such as those shown in Figure 3e-h. Increasing the temperature still further to 800 oC resulted in the formation of more densely-packed columnar architectures of anatase and rutile TiO2 ranging up to 50-nm lateral dimension, as shown in Figure 3i-l. We interpret the growth mechanisms of these structures from ultrasmall nanoparticles in the context of Crystallization by Particle Attachment (CPA)1-2, which refers to the general phenomenon of crystal formation from particle “building blocks” that can include multi-ion complexes, amorphous nanoparticles, or fully formed nanocrystals. CPA is typically studied in in situ TEM experiments in solution where, for example, anatase TiO2 nanoparticles have been observed to crystallize into larger nanostructures by the process known as Oriented Attachment (OA), where the anatase particles rotate, self-align, and then crystallize with perfect alignment such that the c axis of the anatase particles is oriented along the long dimension of the aggregate after apparent attachment.3-5 However, in our case, the as-synthesized “building blocks” delivered to the substrate are TiO2 nanoparticles formed in the gas phase without additional capping ligands. Observing the real time crystallization of nanoparticles in TEM microscopy


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gives some insight regarding the CPA growth mechanisms responsible for the formation of our nanostructures.

Figure 3. Integration of nanoparticles into nanostructures at higher temperatures. SEM and TEM images with corresponding diffraction patterns of samples resulting from pulsed laser deposition of TiO 2 nanoparticles at d = 5 cm in 200 mTorr O2 as in Figure 2, except at the indicated elevated substrate temperatures. In (a-d), fractal, partially sintered anatase architectures are formed at 400 oC. (e-h) For substrate temperatures of 600 oC, TiO2(B) nanobelts become the dominant crystalline architectures. (i-l) Increasing temperatures to 800 oC resulted in the formation of larger columnar anatase crystals.

As schematically indicated in Figure 4a, ~3 nm nanoparticles arrive to attach and crystallize into larger nanostructures at different temperatures. At room temperature, the amorphous particles attach to one another into hyperbranched architectures determined by diffusion-limited aggregation. However the temperature (and kinetic energy) is insufficient to recrystallize either the individual nanoparticle or the aggregate. As the substrate temperature is increased, the nanoparticles integrate and crystallize into various phases and morphologies. It is currently not yet resolved how the deposition conditions and thermodynamics determine the nucleation and initial growth of a preferred crystalline phase and orientation on the substrate, however, the subsequent nanoparticles appear to template on these already crystallized nanostructures to continue their growth into larger crystallites.


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Figure 4. Schematic and TEM images of nanoparticle crystallization by particle attachment. (a) Schematic diagram showing the arrival of amorphous nanoparticles and different integration pathways at different temperatures. (b) TEM image of the top surface of a crystallized anatase TiO2 nanorod prepared at 400 ºC where the nanoparticles at the top of the crystals (indicated in yellow dashed regions) are not yet completely sintered into the rest of the crystal. (c, d) HAADF-STEM images of this top surface region reveals a sub-20-nm layer of nanoparticles that appear to be co-crystallizing and integrating in registry with the underlying crystal. This nanoparticle arrival, incorporation, and crystallization repeats throughout the deposition and growth process.

Figure 4b-d shows several TEM images of a sample prepared at 400 oC that was cooled by turning off the substrate heater just following the last laser pulse. Examination of the top of the anatase nanorods reveals a very thin (sub-20-nm) layer of ~5nm nanoparticles that are not fully integrated into the nanorod. Atomic-resolution high-angle annular dark field STEM imaging of the nanoparticles shown in Figure 4c,d clearly shows that the very topmost nanoparticles are distinct, yet crystalline (anatase) in atomic registry with the crystal planes of the evolving (anatase) TiO2 nanorod tip. Below them, other nanoparticles are also crystallized in apparent registry with each other, as well as the nanorod. Through observations like these it is currently unclear if the amorphous particles contact the existing nanorod first, and then crystallize – templating their orientation from the existing nanorod – or if they crystallize first in the nearproximity of the substrate and then recrystallize to match the orientation of the existing nanorod. Alternatively, oriented attachment similar to that observed in liquids, where crystalline particles rotate to achieve the nanorod’s crystalline orientation, is difficult to imagine in our case of gasphase aggregation because of the difficulty in overcoming the strong van der Waals forces in the absence of a facilitating liquid solvation layer. Ostwald ripening is another mechanism by which smaller nanoparticles can be rapidly assimilated at elevated temperatures.41-42 In contrast to the sequential nanoparticle deposition and crystallization process, another possible pathway of nanostructure growth is co-deposition and co-crystallization of amorphous nanoparticle aggregates. To gain some understanding of the crystallization and co-crystallization 12 ACS Paragon Plus Environment

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dynamics of amorphous nanoparticles in aggregate form, aggregates of amorphous TiO2 nanoparticles were laser-synthesized and deposited directly onto room-temperature TEM grids to investigate crystallization and sintering dynamics at different temperatures. Starting from room temperature, the grid temperature was rapidly and incrementally raised by 100° to the indicated values and held typically for ~20 min at each temperature while recording the crystallization and sintering processes. Several individual movie frames were extracted are shown in Figure 5a-f. Despite some sample drift, the same region could be viewed throughout the heating process.

(a)

200 oC

(b)

300 oC

(c)

400 oC

500 oC

(e)

600 oC

(f)

700 oC

5nm

(d)

Figure 5. TEM imaging of nanoparticle crystallization and sintering dynamics. TEM images of the crystallization and sintering of amorphous TiO2 nanoparticles that were deposited as aggregates by pulsed laser deposition onto room-temperature TEM grids and consecutively annealed in vacuum for typically 20 minutes at (a) 200°C, (b) 300°C, (c) 400°C, (d) 500°C, (e) 600°C, and (f) 700°C. Arrows follow representative nanoparticles as they crystallize and sinter to form larger crystalline domains while preserving the porous nanoparticle network morphology. Some nanoparticles had crystallized at 200°C, and nearly all had by 400°C. Significant recrystallization and sintering was observed by 600 ºC with the evolution of larger crystalline domains.

The arrows on the images allow the tracking of a few representative nanoparticles as they began to crystallize at temperatures as low as ~200 oC. First, we observed that not all nanoparticles crystallize at temperatures below 400°C, indicating an innate difference between the structures of the amorphous nanoparticles affecting the energy required for crystallization. Secondly, changes in the crystallization of the nanoparticles appeared to occur shortly after ramping to a temperature, and then remain stable. For example, at 200°C, waiting for an hour did not appear 13 ACS Paragon Plus Environment

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to significantly change the crystallization fraction or the relative orientation of the crystallites. Third, at temperatures below 500 oC, the nanoparticles within the chain networks appeared to essentially maintain their shape, size, and orientation. The appearance of sharper facets was evident at 400°C, however, indicating that Ostwald ripening was occurring. Fourth, at 600°C and higher temperatures, larger crystalline grains of common orientation were evident without significant reorientation of the original network architecture, implying recrystallization of some nanoparticle domains during sintering. These data reveal that crystallization of the amorphous TiO2 nanoparticle ‘building blocks’ can be induced at temperatures as low as 200°C, however they suggest that if nanoparticles are aggregated before crystallization, the final crystalline structure should be a porous, interconnected network of misoriented nanodomains within our experimental annealing temperature range (i.e. 400 – 800 oC). This mechanism is the basis for the formation of mesoporous, hyperbranched, polycrystalline electrodes used in dye-sensitized solar cells and other applications. This data supports the conclusion that ultrasmall nanoparticles deposit and become integrated sequentially, and not as aggregates, to enable the evolution of nanostructures at higher temperatures into smooth columnar shapes. A similar integration of amorphous TiO2 nanoparticles into crystalline TiO2 nanorods has been reported during the atomic layer deposition of TiO2 by the group of Wang et al..10, 43 However, in that work, metastable amorphous nanoparticles were observed to grow from within an ALDdeposited ~10-nm-thin amorphous film on an existing crystalline nanorod, and their crystallization was observed within this percolating film to form anatase, then rutile phases according to the Ostwald-Lussac law.10 Nanoparticles within this shell were observed to migrate at higher temperatures (600°C) to the lower energy facets at the top of the nanorod, where they aligned and attached by the process of oriented attachment, a process witnessed for the first time by vapor-phase attachment. The mobility and ability to reorient was thought to have been facilitated by the presence of hydroxyl groups from the H2O cycles in the ALD process, indicating that surface functional groups in addition to the metastable amorphous shell are key parameters to allow the orientation and integration of nanoparticles during vapor phase CPA.10 The gas-phase CPA process described here appears to be very similar, however the metastable nanoparticles are condensed in the background gas within the thermalized laser ablation plume, and are delivered directly to the tip of the nanorod, foregoing the need for migration before 14 ACS Paragon Plus Environment

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crystallization and attachment. Some mobility of the particles may result from their residual kinetic energy, and also acoustic forces imparted by the weak shock front propagating through the snowplowed background gas.

Figure 6. Molecular dynamics simulation of the attachment and crystallization of an amorphous ultrasmall TiO2 nanoparticle. (a) A 3-nm-diameter amorphous TiO2 nanoparticle arrives at a ~10 nm TiO2(B) nanocrystal that simulates the experimental nanorod tip surface. (b) Snapshot of the attachment and crystallization of the nanoparticle after 2 ns at 600°C, where >50% of the Ti atoms have already formed octahedral oxygen bonds.

The time in our process for the nanoparticles to attach, crystallize, and sinter to become integrated within the nanorod is the subject of ongoing experimental and computational studies. Prior work using molecular dynamics methods to simulate the collision of two ~3 nm TiO2 nanoparticles has indicated that particles with different initial phases can attach and recrystallize into a common phase as quickly as several nanoseconds at 1200°C.44-47 However, achieving the necking and coalescence of the two particles within relatively lengthy ~150 ns timeframes required accelerated MD methods at very high temperatures (1527°C).44 To gain insight into the processes and timescales (t) involved in our experiments at much lower temperatures, we performed MD simulations based on the Matsui-Akaogi potential describing interatomic interactions48 for a 3-nm amorphous nanoparticle containing 627 atoms arriving at a TiO2(B) nanocrystal surface (consisting of 19200 atoms) as shown in Figure 6(a) at 600°C. Our MD simulations reveal that the amorphous nanoparticle attaches within t ~10 ps, with some interdiffusion of its atoms into the topmost layer of the TiO2(B) crystal.

Starting from a

metastable amorphous nanoparticle with 8% TiO6 octahedra, the nanoparticle then proceeds to crystallize such that within t~2 ns as shown in Figure 6(b), >50% of the Ti atoms have formed octahedral oxygen bonds. Importantly, the contact with the substrate crystal induces 30% more octahedral bonds than for nanoparticles that were isolated and exposed to the same temperature increase. Although demonstrated in one specific configuration, the results clearly indicate that at 15 ACS Paragon Plus Environment

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the relatively low temperatures involved in our experiments, amorphous nanoparticles can rapidly attach and achieve significant crystallization induced by the substrate within nanoseconds. To estimate the time scales for integration of the attached nanoparticles into the nanorod at our temperatures, which we refer to as sintering, experimentally we note that we could find crystalline nanoparticles that had not integrated fully into the nanorod tip following PLD at 1 Hz at 400°C, as shown in Fig. 4. However for depositions at 600°C temperatures and above, we found smooth nanorod top surfaces for 1 Hz laser repetition rates indicating that nanoparticles become sintered and incorporated in those timescales. Both Seto et al.49 and Kobata et al.50 experimentally studied the gas-phase sintering of TiO2 nanoparticles in floating aggregates, and derived scaling relationships based upon a surface diffusion neck growth model51 from their data that are used today for comparison with results of MD simulations.44 Sintering was found to follow an Arrhenius temperature dependence, and the sintering time for particles was found to scale as d4, where d is the nanoparticle diameter.49-50 Using the experimental relationship from Seto summarized in Ref.44 (where it should be noted that, based on MD simulations, it is predicted that smaller nanoparticles deviate from that relationship to sinter much faster than the 10-100nm particles used in Ref.49), we can estimate upper limits for the sintering time of the 3nm diameter TiO2 nanoparticles in our study of 0.1 ms at 800°C, 70 ms at 600°C, and 2100 s at 400°C. These estimates support the hypothesis that nanoparticles deposited at 600°C and 800°C can sequentially attach and sinter into the evolving nanorods within the 1000 ms interpulse arrival time of particles in our PLD experiments, in agreement with our observations. However, also in agreement with our experiment, due to the steep Arrhenius behavior, the nanoparticles at 400°C are predicted to not become fully integrated at the top of the nanorods during this time, yet may sinter to be completely integrated nearer the bottom of the nanorods during the ~ 1800 s time of the deposition. Note also that based upon the d4 dependence of the sintering time, 10-nm nanoparticles should sinter ~120 times more slowly at a given temperature than ultrasmall (~3 nm) nanoparticles. As the ultrasmall nanoparticles sinter within the aggregate chains to form much larger crystallites, their evolution slows to achieve a terminal particle diameter, as in Figure 5. This explains why


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sequential deposition of individual, ultrasmall nanoparticles as “building blocks” is advantageous over the deposition of aggregates. In conclusion, the conditions for pulsed laser deposition (PLD) can be adjusted to selectively form and deliver ultrasmall (~3-5 nm) amorphous TiO2 nanoparticles in the gas-phase as “building blocks” for dynamic crystallization into nanostructured films with metastable phases and interesting morphologies. The nanoparticles condense in the decelerating laser plasma plume due to background gas collisions, and can be deposited at distances beyond the atomic and molecular component of the plume due to the longer ballistic range predicted for individual nanoparticles. The in situ plume diagnostics, TEM observations of nanoparticle sintering, and associated modeling support the conclusion of crystallization by particle attachment (CPA) from sequentially-deposited, individual, ultrasmall nanoparticles as the mechanism for the growth of crystalline nanostructures at the 1 Hz repetition rates and temperatures used in these experiments. Several CPA mechanisms are possible for the rapid integration of the UNPs into the crystalline nanostructure, depending on whether the individual amorphous UNPs crystallize during the penetration of the gas-phase thermal gradient near the heated substrate surface. These range from addition of amorphous nanoparticles and crystallization from the growing nanocrystal surface, to oriented attachment, Ostwald ripening, and non-oriented attachment of crystalline nanoparticles followed by recrystallization. The adjustment of conditions in PLD for the deposition of UNPs as ‘building blocks’ of crystalline nanostructures appears quite general, and can be used to explore the synthesis of not only a variety of crystalline oxide nanostructures, but also other materials, such as two-dimensional metal chalcogenides.52-53 This report indicates the promise of exploration and exploitation of ultrasmall nanoparticles as intermediate “building blocks” in the formation of nanostructured films with novel phases.

METHODS Pulsed Laser Deposition and In Situ Diagnostics. Gas phase condensation in PLV was use for synthesizing TiO2 UNPs as building blocks for the CPA process (Figure S1 in Supporting Information). A 2 inch TiO2 pellet (Praxair Company), composed of rutile powder, was used as the ablation target in this experiment. A pulsed KrF (248 nm, 1 J cm-2, 1 Hz) laser was used for 17 ACS Paragon Plus Environment

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the ablation of the targets in oxygen background gas to produce a spatially confined laser ablation plasma plume that thermalized over a 5 cm stopping distance, and resulted in the synthesis of UNPs with a diameter of 2-5 nm. A 1-inch-diameter stainless steel heater (HeatWave Labs, Inc.) was placed d = 5cm away and parallel to a 5-cm-diameter TiO2 bulk crystalline in a cylindrical chamber (50 cm inner diameter, 36 cm tall) to deposit the nanoparticles at various substrate temperatures. Gas pressure was controlled with a mass flow controller (Ar 99.995%, 0-500 sccm) and a downstream throttle valve.

The heater was

controlled by an integrated thermocouple to ± 2°C. The plume imaging was performed with an ICCD camera system (Princeton Instruments, PI-MAX) with variable gating (minimum 5 ns). The f4.5 Nikon camera lens was positioned 46 cm away from the plume center, outside the chamber, looking at the plume through a 2-inch x 8-inch Suprasil window. Exposures time and fstop were varied from 5 ns to 5 ms, and 32 to 4.5, respectively. The ion probe time-of-flight waveforms (voltage measured across 50 ohms, detector area = 1 mm2) were acquired as a function of pressure and distance during the ablation process. TEM Characterization of Nanoparticle Sintering. For observation of crystallization dynamics in TEM, amorphous TiO2 nanoparticles were laser-synthesized and deposited directly onto lacey carbon TEM grids at room temperature, forming aggregated nanoparticle networks. The grids were then transferred onto a heating stage and inserted into the TEM for high-resolution electron microscopy (FEI Titan S aberration-corrected TEM-STEM, acceleration voltage of 200 kV). The crystallization and sintering dynamics of a small cluster of nanoparticles that were overhanging an empty grid hole were observed in situ at different temperatures using video capture at 25 frames/second and 8ms exposure time for each frame. The high-resolution TEM of nanoparticles deposited by PLD at high temperatures and then cooled to room temperature was performed with a Zeiss Libra 200 MC TEM at an acceleration voltage of 200 kV, and also with a fifth-order aberration-corrected STEM Nion UltraSTEM 200. Raman Spectroscopy and XRD Characterizations. A Renishaw Raman spectroscopy system (532 nm excitation source focused onto the samples through a 100X, 0.8 NA objective lens) was used to obtain the Raman spectra of the samples. An X-ray diffractometer (Panalytical X’Pert MPD Pro) with Cu-Ka radiation (l =1.54050Å) was used to acquire XRD patterns. (Figure S2 in Supporting Information) 18 ACS Paragon Plus Environment

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ASSOCIATED CONTENT Corresponding Author *To whom correspondence [email protected]

should

be

addressed:

Email:

[email protected];

Notes The authors declare no competing financial interest. Supporting Information Available: Schematic representation of the PLD setup, XRD data, Raman spectra, and room temperature TEM images of the nanoparticle. This material is available free of charge via the Internet at http://pubs.acs.org. Acknowledgements Synthesis science including PLD, in situ plume diagnostics, XRD, TEM analysis, SEM and AFM studies and MD simulations (MMS, MT, DBG, CMR, AAP, GD, KW, KX, GE) were supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences (BES), Materials Sciences and Engineering Division and performed in part as a user project at the Center for Nanophase Materials Sciences, which is a DOE Office of Science User Facility. The authors gratefully acknowledge the assistance of A. Lupini on the Z-STEM measurements. This research used resources of the National Energy Research Scientific Computing Center, a DOE Office of Science User Facility supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. Author contributions M.M-S. performed the synthesis experiments. M.T, M.C. and G.D. carried out the TEM studies and analysis. A.A.P., D.B.G., M.M-S. participated in spectroscopic diagnostic experiments. M.Y. and J.L. performed MD simulations and analysis. All the authors contributed to data analysis, discussed the results, wrote and commented on the manuscript.


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Figure 1 177x104mm (300 x 300 DPI)

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Figure 3 177x105mm (300 x 300 DPI)


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