Metal/organic-molecule interactions underpin many key chemistries but occur on sub-nm scales where nanoscale visualisation techniques tend to average over heterogeneous distributions. Single molecule imaging techniques at the atomic scale have found it challenging to track chemical behaviour under ambient conditions. Surface-enhanced Raman spectroscopy can optically monitor the vibrations of single molecules but understanding is limited by the complexity of spectra and mismatch between theory and experiment. We demonstrate that spectra from an optically generated metallic adatom near a molecule of interest can be inverted into dynamic sub-Å metal-molecule interactions using a comprehensive model, revealing anomalous diffusion of a single atom. Transient metal-organic coordination bonds chemically perturb molecular functional groups > 10 bonds away. With continuous improvements in computational methods for modelling large and complex molecular systems, this technique will become increasingly applicable to accurately tracking more complex chemistries. Tracking single molecule movements is a challenging task, but highly desired for applications and fundamental studies. Here the authors reconstruct the sub-angstrom relative movements of a molecule interacting with a metal adatom, by measuring its vibrational spectrum in a self-assembled monolayer, continuously modified by the adatom in a nanoparticle-on-mirror construct. On this same scale, metal-organic interactions underpin many key chemistries from catalysis for feedstocks4 to anti-bacterial surfaces5 and green initiatives6. To extract properties at these atomic length scales, myriad imaging techniques have been developed including electron8, atomic-force9, scanning-tunnelling11, and tip-enhanced Raman13 microscopies. Although these are capable of visualising individual atoms, a technique is still required to allow single atom dynamics to be studied at relevant timescales and at ambient conditions. This impedes progress in fields from molecular electronics14 to photocatalysis15 and electrochemistry16. Such a technique is essential to understand how atoms explore their energy landscape and interact with their local chemical environment and to facilitate observations of chemical reactions at their operating temperature.The push for energy-efficient electronics and growing needs for sustainable manufacturing are driving the miniaturisation of devices and the development of functional nano-components, with promising technologies now incorporating atomic-size features17. Single metal adatoms stabilised on the surface of these plasmonic nanocavities further localise the confined optical field to sub-nm3 effective volumes termed picocavities19. This allows single molecules to be optically isolated while the resulting extreme field gradients alter the spectrosopic response, allowing its spectral fingerprint to be distinguished from the ensemble. Such picocavities, which are produced stochastically under laser irradation of the nanocavity, are stable at room temperature depending on nearby chemical species20.By placing molecules within the sub-wavelength-confined optical fields of nanostructured plasmonic metals, the enhancement of scattering and absorption opens up new methods of non-destructive and real-time optical interrogation. In nanometre-scale gaps between plasmonic materials, the enhanced field strengths enable optical spectroscopy of fewer than a hundred molecules, in contrast to the trillions observed with classical optics21 model allows the experimental spectra to be inverted into sub-Å relative movements of the molecule and adatom. We discuss also the limitations of this approach and highlight new theory needing development.Here, the chemical influence of a low-coordination gold adatom on neighbouring organic molecules is probed by the tightly-confined light, revealing their mutual interaction. This is experimentally achieved using surface-enhanced Raman spectroscopy (SERS) to reveal how the vibrational modes of the molecule are continuously modified by the presence of the adatom Fig. . These md < 2 nm from a gold mirror  and the width and refractive index (ng) of the gap as 3. The resulting field enhancement offsets low Raman scattering cross-sections allowing for inelastic scattering of light to probe the vibrational modes of any material placed within this volume. Using a self-assembled molecular monolayer (SAM) between the nanoparticle and mirror provides robust Raman signatures due to its close-packed nature22. Light-induced picocavities can then interact with these molecules. The cyanobiphenyl-4-thiol (NC-BPT) molecular SAM used here  construct consisting of a gold nanoparticle positioned ror Fig. 17. Upon ere Fig.  containstem Fig.  consiste20. Because only one red-shifted C≡N line is typically seen, this suggests it is from a single molecule18, as supported by its narrower linewidth and spectral wandering. Picocavity events in which a modified C≡N stretch is not observed are considered to be atomic features appearing on the gold mirror and are not explored here20. We emphasise that such transient SERS peaks from adatom atomic features are pervasive in all nanogap nanocontructs24.During a picocavity event, the strong optical field gradient over a single molecule next to the Au adatom alters the effective Raman selection rules, enabling previously weak modes to become visible and modifying peak ratios. Most obvious is a distinct adatom modified C≡N stretch observed at lower Raman shifts than expected  is 0.3 for the Au-N coordination bond while the C≡N bond order decreases from 3.0 to 2.7. Optimising the position of the gold atom with respect to the molecule allows DFT to accurately reproduce the experimental results  and only ~15% form on the nanoparticle facet20. Once formed, they are probed with 25 μW irradiation for 60 consecutive spectra with 1 s integration time. Among trajectories recorded for 87 individual picocavities, the energy dynamics of the perturbed CN vibration displays both continuous and discrete changes on these timescales  is observed to range over 400 cm−1 but occurs most often around −1 , follow an exponential distribution with an average wandering rate −1 s−1  and the plane of the upper phenyl ring (xz plane) with the N atom as the origin  rapidly increases in energy as dNAu decreases, generating sets of anti-crossings with successively higher energy vibrations  as for the N−Au Raman in Fig. dNAu coordinate, a thermal fluctuation of 24 pm is expected, equivalent to changes in νCN of 40 cm−1 . Again, the vibrational modes are strongly modified by the Au adatom as its valence orbitals overlap with the π-orbitals of the ring giving a 0.1 eV local energy well at dCAu = 2.2 Å with a MBO of 0.5 for the C−Au bond  upon adatom perturbation Fig. . Without−1 V−1 28. This effect can be replicated in DFT by replacing the adatom with a point charge, but magnitudes as large as 0.5e− change the vibration energy of this bond by only ≤6 cm−1  cm ± 5 cm−122 this would require the adatom to traverse an unfeasibly large distance. Occasionally, the CN-picocavity scans display split C≡N vibrational lines. While not explored in depth here, DFT suggests that these could be caused by the hybridisation of C≡N vibrations from two molecules coupling through mutual interaction with a single adatom . The net result is that no scaling factor for each vibrational mode can place the experimental frequencies within the range of possible frequencies supplied by the DFT for all times . By maximising the averaged Spearman correlation, this forces the trend of each peak perturbation to be captured simultaneously while allowing the correlations to be nonlinear. To test the reproducibility and uniqueness of the solution, five random initial trajectories are maximised in this way, each returning very similar optimised trajectories with averaged correlations ranging from 0.952 to 0.972 , shows its predicated dynamics from DFT is well anti-correlated to the extracted experimental trace . While the final Spearman correlation degrades as expected, the two site feature persists   giving α = 0.87 ± 0.03 and Γ = (0.92 ± 0.03) Å2 s−1. This confirms that the sub-atomic thermally-diffusive motion experiences extra hindrance for √MSD < 3 Å , which likely comes from the directionality of the bonding involved. This adatom diffusivity agrees in magnitude with previously extracted values for copper adatoms on gold nanocrystals by X-ray diffraction33. We point out that using an overdamped Ornstein−Uhlenbeck model, the measured diffusivity −21 m2 s−1k ~ 10 J m−2, this suggests that the fluctuation correlation time τ ~ 100 ms. The large adatom damping −2 s thus yields slow dynamics.Over short timescales where the bounded nature of adatom confinement has little effect, the adatom dynamics show anomalous sub-diffusion Fig.  with a mπ orbitals of the molecule below and the Au facet above . For BPT, picocavity lines are chosen that appear near 1500 cm−1. To compare these to DFT, dNAu is scanned to reveal the energy required to shift the CN vibration frequency  in ethanol  for 18 h. The sample was then rinsed with excess of ethanol and blown dry using nitrogen.Gold nanoparticles were purchased from BBI Solutions (80 nm citrate capped) and stored at 5 °C. All chemicals were purchased from Sigma Aldrich and used as received. Template-stripped gold was prepared by cleaning a silicon wafer with Decon 90 and rinsing twice with water, ethanol, and isopropyl alcohol, and blown dry using nitrogen. A 100 nm thick layer of gold was deposited at a rate of 1 Å∙sFor nanoparticle deposition, 100 µL of colloidal gold suspension  was mixed with 5 µL of a 20 mM solution of KCl and allowed to mix for 30 s. A 30 µL droplet was deposited on the SAM functionalised gold substrate and left for 30 s, after which the droplet was rinsed off with an excess of deionised water and blown dry using nitrogen.NPoM constructs were imaged using an Olympus BX51 microscope in dark field configuration using a 100x Zeiss Ex Epiplan-Neofluar 0.9NA objective lens and a Lumenera Infinity 2 CCD. A Prior Scientific motorised stage was used to centre nanoparticles. A 632.8 nm Integrated Optics SLM VBG diode laser focussed to a diffraction-limited spot was used to excite Raman scattering, with scattered light imaged onto an Andor Newton 970 EMCCD through a HORIBA Triax 320 spectrometer using a 600 l/mm diffraction grating with 750 nm blaze. The elastically scattered light was removed using two Thorlabs 633 nm notch filters and laser power on the sample was measured using a Thorlabs PM16-121 power metre.21. In these computations, the B3LYP34 hybrid density functional, and D3 dispersion correction with Becke−Johnson damping35 were employed with the Def2TZVP basis set37. The adatom was modelled as a single gold atom, while the gold mirror was modelled as one or two gold atoms attached to the sulfur atom. The number of gold atoms was chosen for each case to retain the singlet electronic state of the system. Unrestricted self-consistent field calculations were performed to ensure electronic stability for the wavefunction. For more details see Supplementary Notes All DFT calculations were carried out with Gaussian09 Revision E programme suiteSupplementary InformationDescription of Additional Supplementary FilesSupplementary Movie 1


In modern 3D microscopy, holding and orienting arbitrary biological objects with optical forces instead of using coverslips and gel cylinders is still a vision. Although optical trapping forces are strong enough and related photodamage is acceptable, the precise (re-) orientation of large specimen with multiple optical traps is difficult, since they grab blindly at the object and often slip off. Here, we present an approach to localize and track regions with increased refractive index using several holographic optical traps with a single camera in an off-focus position. We estimate the 3D grabbing positions around several trapping foci in parallel through analysis of the beam deformations, which are continuously measured by defocused camera images of cellular structures inside cell clusters. Although non-blind optical trapping is still a vision, this is an important step towards fully computer-controlled orientation and feature-optimized laser scanning of sub-mm sized biological specimen for future 3D light microscopy. Optical manipulation of large objects is challenging as optical trap positions are blindly chosen. Here, the authors present off-focus interferometric tracking, which localises optimal grabbing positions with increased refractive index, by analysing the beam deformations of several holographic optical traps. When we buy fruits in a super-market, we often grab, rotate and squeeze the product with our hands to investigate it from all sides. While this process is relatively easy in a macroscopic world, a precise control of object orientation is still a big challenge in the microscopic world. Especially in 3D microscopy, this so far non-existing feature would be very beneficial for a visual object inspection.1 or light-sheet microscopy2. Without toxic clearing techniques3, cell clusters, small plants or embryos with sizes from tens of µm to few mm are often so thick that illumination light and/or fluorescence light is absorbed or scattered so strongly4 that observation from different directions is often the only way out5. Embedding the specimen in soft gels (e.g. agarose) prevents unwanted motion or diffusion in liquid environments. One-axis rotatable gel cylinders can make alignment procedures complicated, hinder object growth and make multi-object investigations nearly impossible.The behavior of cells on a glass coverslip is often significantly different to cells embedded in a natural matrix of adjacent cells. Therefore the preparation, handling, and the investigation of cell communication and responses to stimuli require more advanced manipulation for imaging methods, such as e.g. two-photon microscopy6. Optical forces, adequately distributed across the specimen as sketched in Fig. 7 are the most flexible tool to reorient and rotate objects without mechanical contact like invisible fingers9, However, reaching several local potential minima is not possible with multiple optical traps and most objects, which are hardly deformable - leading to blind or frustrated trapping.Repositioning and reorientation of objects with electromagnets can be achieved by adding micro-magnets to millimeter-sized specimen either directly or inside rotatable gel spheres embedding the biological object11, which can also reorient single bacteria during conventional fluorescence12 or super-resolution fluorescence13 imaging. Based on a balance of scattering forces, objects were shifted by four beams to allow multiview imaging14 or by two fiber-emitted counter-propagating beams, where additional cell rotation was achieved by rotating the asymmetric beams15. As an alternative to holographic trapping, time-multiplexed point traps have allowed to hold and manipulate deforming bacteria16 or single microtubules17 in combination with fast super-resolution coherent imaging. Large specimen, with tens of micrometers in diameter, have been trapped and oriented by using low focused, counter-propagating laser beams enabling a large working range in combination with acoustic forces and conventional imaging18 or with light-sheet fluorescence imaging19.Several approaches have been pursued to trap, hold and orient biological objects with optical forces to improve imaging. Single cells or bacteria have been rotated by multiple-point holographic optical traps during brightfield imaging20, for a position clamp and single-axis rotation21 or counter-propagating beams22. Using time-multiplexed optical traps and electro-optical tunable lenses, Tanaka et al.23 realized optical multiple-force clamps for 3D rotational control of diatoms and their fragments. Surface relief imaging approaches used dynamic optical traps to scan particles across a structured surface requiring advanced feedback mechanisms to control the position of either small spheres24 or complex-shaped styluses using force feedback25. An advanced method to rotate small objects performs a 3D tomographic scan to reconstruct the 3D refractive index (RI) distribution of cells of arbitrary shape26. From this, a 3D coherent light distribution is calculated and generated holographically in real time, which exerts optical torques in different directions. However, because of the intensity gradients, which are manifold weaker than those of point traps, only small objects such as a few cells can be rotated.However, in all these approaches, highly precise 3D control in position and orientation has not been possible, mainly because of a missing feedback, also called force clamp. This was partly achieved in early approaches with elongated particles (diatoms) with optimized video controlThe driving optical forces come to their limits for increased gravitation and friction forces—counteracting the rotation of cell clusters or small embryos. Figure Thus, non-blind optical trapping means that an efficient trapping position within the RI distribution is found by any readout signal to enable stable object displacements and small rotation steps. Non-blind optical trapping with point tweezers is still a vision and might be achieved within the next 5–10 years. It is of great interest in modern 3D microscopy to realize a purely optical position and orientation control of larger cell clusters or small embryos, i.e., of arbitrarily shaped specimen, using force-optimized grabbing positions for each optical trap. The localization of such local RI changes  optical trapping applications and provides nanometer precision in 3D at megahertz sampling rates using quadrant photo diodes (QPDs)29. However, BFP tracking is typically limited to one particle or structure at a time (time-multiplexing)31, since the signal is recorded in the Fourier plane of a detection lens (DL), where the scattered fields from all particles overlap. Usually assignments to different positions is impossible for holographic traps32. However, BFP interferometry can also be used to scan asymmetric structures, such as cell membranes34 or even for shape tracking. In principle, interferometric tracking is also possible with a QPD in the focal plane35, but requires spatial filtering for fixed focus positions for parallel trapping of several particles36.This localization problem resembles 3D particle tracking techniques, often using interferometric approaches and measuring phase differences between scattered and unscattered light. Important developments are summarized: The most prominent interferometric tracking technique is back focal plane (BFP) interferometry37. In reflection mode, sensitivity and the signal-to-noise ratio can be strongly increased with interferometric scattering (iScat)38 to track several nanometer-sized particles typically in lateral directions.Most other tracking techniques based on interferometry are holographic microscopy methods using a plane reference wave. Here, a single brightfield image from a camera enables 3D position tracking of several beads (in transmission mode), which works well, if particles are well-defined and not too close to each otherIn this work, we show how to sense interferometrically and in parallel local RI changes at several 3D positions within a 3D object—only by scattered laser light. Our technique, called off-focus interferometric (OFI) tracking, measures positions of particles  by analyzing several interference patterns from different foci distributed over a camera. We present a local intensity gradient scheme—similar to that of QPD tracking—which we apply to beam cross-sections, which are inherently off-focus for 3D distributions of foci. We perform advanced computer simulations to demonstrate that individual RI changes inside a focus volume can be retrieved even in the middle of clusters of scattering spheres.b is characterized by sufficiently strong optical gradient forces 39 and with vacuum speed of light c.Non-blind optical trapping is complicated, because on the one hand counteracting gravitational and frictional forces often outbalance the 3D optical forces. On the other hand, conventional imaging does not provide sufficient object information as feedback for the optical traps to find the optimal grabbing position, which is not automatically the edge of an object, as often assumed. This optimal position λ =  1.06 µm). The laser light, scattered at multiple objects in multiple foci, is projected telecentrically by a DL  and other achromatic lenses (L) onto a camera , which is conjugated to the focal plane . The lateral magnification is M = 20, defined by the ratio of the focal lengths of L and DL. As outlined exemplarily in Fig. N = 3 traps at different 3D positions, N laterally displaced intensity distributions n ≤ N). Hence, the camera records N focal spots at axial positions zdn·M², which are typically all off-focus. The rectangular inset on top defines the particle’s 3D trapping position tn relative to the center of the reference FP plane .The setup for OFI tracking requires a computer-controlled holographic trapping unit and a camera with sufficient bit-depth to record several defocused laser foci. These components were added to a commercial inverted microscope  and are sketched in Fig. ane Fig. . The scan(bn) and thereby optical forces  can be N < 10 and an adequate camera position at zd. Figure zdn behind the geometrical focus . The incident electric field Ei (including wavefronts) is shown in blue, the scattered electric field b  is sketched in green. The resulting off-focus interference (OFI) intensity distribution N deformed beams at the camera in the (conjugate) plane zd readsnm).In our simplified model, scattering between the particles is neglected, i.e., θs vanishes (θs = 0), whereas θs > 0 for larger particles.The position-dependent difference between the k-vectors of the incident and the scattered wave, bz with the Gouy phase difference 29.The interference intensity zd = 20 µm along x are sketched in Fig. b = ·μm is shown in Fig. ns = 1.47. The vacuum wavelength was λ = 1.06 µm, the RI of water nm = 1.33 and the numerical aperture of the focusing lens NA = 1.2.The three intensity profiles b can be analyzed by integrating the intensity n  and that the pupil planes of both lenses ensure collimated illumination and detection, respectively.The basic requirement for correct instrument alignment and precise experiments is that the OL and the DL have a confocal alignment , which is the object plane , without knowing the optimal grabbing positions. As illustrated in Fig. N = 6 optical traps were distributed equally across a spherical surface with 70 µm in diameter  Fig. , each wins = 1.47 and 1 µm in diameter was easily possible, which has become a standard task over the last decade. While stable optical trapping is not a big challenge, it is more difficult to achieve precise, linear and orthogonal particle tracking in 3D. Therefore, we measured the detector responses for orthogonal particle displacements through each of the six laser foci in a meander-like scan as outlined in Fig. Parallel trapping of six glass beads with RI n = 5th trap with axial displacement zd5 = 20 µm and lateral bead displacements b = ·µm. By subtracting the intensity distributions b = ·µm become visible in the two plots of Fig. bx = 0). A horizontal line scan zd = +20 µm were performed and are displayed in blue colors in Fig. zd. The corresponding line scans are shown in the two graphs for lateral and axial particle displacements in Fig. bj = 0 .To illustrate the optical principle and to allow comparisons to the theoretical model, Fig.  in Eqs.  and 4),,\documenz40. However, this is different for a focused wave, when the sphere size and phase perturbation increase relative to the decreasing beam diameter. We performed computer simulations using the beam propagation method (BPM)40 to investigate whether it is possible to measure the characteristic phase change When a plane wave propagates through a cluster of beads, the phase perturbation entclass1pt{minimab =  generates diffraction patterns zd = −4 µm. Two such off-focus camera signals from beam positions  and  are shown in green and orange frames. We performed more than 15,000 BPM simulations for different focus positions within a 50 µm cluster and for different detector positions. Figure This computer experiment is illustrated in Fig. b that generates contrast, e.g., in brightfield imaging with intensity 39, but not for off-focus planes (defocussed by zd = 20 µm).A spherical bead represents a structure with well-defined shape and RI change zd = 20 µm). First, we scanned a small cluster of unlabeled RH30 cells attached to a coverslip through one of the laser foci. The cross-section of the distorted beam on the camera was analyzed according to the steps explained in the theory section to obtain the signal maps Here, we demonstrate in a second type of experiments that OFI scanning allows to acquire also high-resolution maps of refractive indices although the trap is clearly off-focus :41Here, mentclass2pt{minimzd = 20 µm). In consequence, it must be possible to perform OFI scans of cell clusters also in differently oriented planes or along arbitrary trajectories opening the option to scan along specific volume areas of interest, such as the edges of cells at the periphery of the cell cluster.Figure In a third type of experiment, a 70 µm large cell cluster remained stationary and three laser foci—now displaced by continuously varying holograms of the SLM—scanned three cells at the edge of the cluster as illustrated in Fig. The motivation of this study was to evaluate a technical approach based on optical forces that allow to flexibly hold and rotate biological specimen of a few 100 µm in size for taking 3D images —without using coverslips or gel cylinders for object mounting. It has been shown in several studies that single cells can be well rotated in three directions using optical tweezers at laser powers of a few milliwatts. However, larger specimen, consisting of hundreds or thousands of cells are difficult to manipulate because of increased friction and gravitational forces—especially for rotations out of the image plane and for blind optical forces Fig. .x-axis, i.e., out of the image plane is possible within 140 s using 3 optical traps  in distances zd = ±5 µm … ±50 µm away from the laser focus, from which a center-of-intensity signal N x100 scattering patterns from N beams could be analyzed to find the optimal beam deflection and grabbing position, which will be one of the next steps to do.The OFI principle is different to BFP trackingion Fig. . This noThere are two typical working modes: (i) OFI is used in tracking mode, when small structures move within the stationary laser focus volume. (ii) OFI is used in imaging mode, when the laser focus is scanned across a stationary structure.S(t) of the same small structure having a constant RI  are recorded. OFI tracking combines the advantages of holographic video tracking37 and BFP interferometric tracking29. The first technique uses a planar reference wave and a camera and hence can track several particles in parallel. The second uses a focused laser beam and a QPD, but does not require the knowledge of any optical properties of the particle, especially when thermal noise calibration is applied. As shown in Fig. λ laterally, 2–4 λ axially) is of similar length as in BFP tracking, provided that the defocused distance zd is not too small . Although larger off-focus distances provide longer linear detection ranges , this goes along with larger cross-sections on the camera, leading to an overlap between adjacent beams, which can complicate postprocessing and analysis. However, as shown by the adjacent positions of optical traps n = 1 and n = 6 in Fig. Sxy and Sz, change inversely with the beam diameter, which is known from BFP tracking42. Here, analyzing larger  beam cross-sections does improve linearity and sensitivity for lateral  tracking. This would be also one of the future steps to go with OFI tracking.In the OFI tracking mode for a static focus, the time trajectories Rs > M²·10 µm . This enables a sufficiently precise and flexible tracking of several particles in parallel, when computer-holographic focus control is favored, relative to time-multiplexed approaches23. The holographic trap control brings us to the second OFI mode, which can be used in connection with the multiple tweezers based rotation of cell clusters and other biological objects, which are often larger than 100 µm.Hence, both the position of the camera plane and the distribution of laser foci in 3D space should be chosen such that defocused spots have a radius  µm Fig.  and do nzd. A confocal scanning microscope, for comparison, typically collects incoherent fluorescence light at each focus position using sensitive point detectors, such that the fluorescence intensities Ifluo(b) are mapped to an image.In the OFI imaging mode for static objects, the position signals OFI uses coherent light from a single laser, which propagates through inhomogeneous material, suffering from multiple phase perturbations, and measures a differential signal. Whereas the phase front of a beam with a large cross-section is only slightly distorted by the scatterers, the same type of scatterer in the beam’s focal region (minimum cross-section) generates a large phase perturbation and a stronger redirection of momentum vectors . In consequence, the spatial change in beam intensity is more pronounced, which can be well detected on an off-focus camera plane as illustrated by the computer experiments in Fig. S(b) can be plotted as a vector, displaying both the strength and the direction of the RI changes zd, also oblique scans can be performed without requiring any refocusing as illustrated in Fig. In Fig. 44, which is however typically measured in the BFP. Here, the reflection of the mean momentum vector permits direct measurement of the momentum transfer, and hence the force, applied to the trapped object—assuming that all the scattered light can be collected. Although measuring absolute optical forces would be another interesting future investigation, the main purpose of this study was to identify a strategy to detect relative changes in optical forces as described by relation (7), where The OFI approach is related to direct optical force measurements based on a shift of the center-of-intensity position45. OFI imaging might have potential applications in fast object screening of especially biological specimen, consisting of hundreds to thousands of cells. Here, OFI could provide a distribution of refractive indices  relative to the fluorescence distribution, which could be imaged in parallel.Further scanning experiments with different laser illumination wavelengths need to be performed to investigate the influence on scattering strength, image contrast and image resolution. Alternative approaches to apply OFI also for high-quality 3D coherent imaging might be to vary, e.g., the illumination angle and the degree of coherenceMotivated by the vision to stably hold, move, and 3D orient large biological specimen such as cell clusters and small embryos in the future purely by optical forces, we have gone an important step in this direction with the present study. We have shown that by OFI label-free imaging of RI changes is amazingly robust against multiple scattering, such that this technique is not only useful for the localization of grabbing positions for optical traps, but also for 3D imaging. As the basic requirement for this long-term tweezing challenge, we demonstrated on the one hand the optical force driven out-of-focal plane rotation of a 70 µm large cell cluster, and on the other hand single off-focus camera shots to recover the 3D gradient force distribution within cell clusters. Our future goal is to let multiple feedback-controlled laser foci find their optimal grabbing positions at positions of high RI gradients on their own—similar to our fingers trying to rotate a very fragile or spiky object, with each finger continuously sensing the pressure applied to the object Fig. . We wantRH30 cancer cells, RH30 human rhabdomyosarcoma (DSMZ Braunschweig – ACC 489) were cultivated in RPMI 1640 medium with 10% FCS . Cells growing adherently are typsinated with Trypsin/EDTA-solution for counting and passaging.Serial dilutions with different cell concentrations were performed in 50 ml cubes with, e.g., 50,000 cells in 5 ml medium (corresponding to 200 cells in 20 µl). Further dilutions 1:2, 1:4, and 1:8 (=25 cells in 20 µl) were performed.2 concentration of 5% until the cell clusters inside the hanging droplets have grown to the desired size.A black cross mark on the lid of a 6 cm culture plate separates four different regions. Several 20 µL droplets with different cell concentrations are pipetted on the lid, which is then carefully turned upside down to be placed on the lower part of the culture plate. 2.5 mL phosphate buffer are added to the plate for sufficient humidity in the local atmosphere and to prevent medium up-concentration. Cells are cultivated at 37 °C the COFurther information on research design is available in the Peer Review FileDescription of Additional Supplementary FilesSupplementary Movie 1Supplementary Movie 2Supplementary Movie 3Supplementary Movie 4Supplementary Movie 5Supplementary Movie 6Reporting Summary


Copper (Cu) nanoparticles (NPs) have received extensive interest owing to their advantageous properties compared with their bulk counterparts. Although the natural oxidation of Cu NPs can be alleviated by passivating the surfaces with additional moieties, obtaining non-oxidized bare Cu NPs in air remains challenging. Here we report that bare Cu NPs with surface excess electrons retain their non-oxidized state over several months in ambient air. Cu NPs grown on an electride support with excellent electron transfer ability are encapsulated by the surface-accumulated excess electrons, exhibiting an ultralow work function of ~3.2 eV. Atomic-scale structural and chemical analyses confirm the absence of Cu oxide moiety at the outermost surface of air-exposed bare Cu NPs. Theoretical energetics clarify that the surface-accumulated excess electrons suppress the oxygen adsorption and consequently prohibit the infiltration of oxygen into the Cu lattice, provoking the endothermic reaction for oxidation process. Our results will further stimulate the practical use of metal NPs in versatile applications. A growth strategy based on surface-accumulated excess electrons enables the realization of non-oxidized bare copper nanoparticles stable under ambient conditions. This inevitable formation of Cu oxides under ambient conditions deteriorates the metallic properties of Cu, and the process is accelerated in Cu nanoparticles (NPs) with a large surface area, which have the potential to be a low-cost alternative to precious gold and silver analogues in various applications2. Thus, in the synthesis of Cu NPs, such as by wet chemical synthesis and solid-state thermal treatment , the artificial manipulation of surface structure via additional post-treatments is a mandatory process. One example is surface passivation with thiolate ligands to retard oxidation5. Another process, based on chemical etching6, thermal heat treatment7 and photoreduction8, is to eliminate the naturally formed Cu oxide moieties on the surfaces of Cu NPs. Although non-oxidized Cu NPs can be synthesized by the in situ reduction of Cu precursors in the form of composites with an epoxy matrix9, there is no practical way to autonomously prevent the oxidation of bare Cu NPs in air.Copper (Cu) has an electron configuration of [Ar]310, electrons from a metal migrate to the surface and bind with the adsorbed oxygen molecules, forming oxygen anions. Then, metal cations combine with the oxygen anions due to the surface dipole between them, and metal oxides form at the surface. This oxidation mechanism, thus, implies that it is possible to prevent the oxidation of a metal by preventing the release of electrons from metal atoms or by hindering the adsorption of oxygen. Indeed, the classic cathodic protection technique prevents metal corrosion by supplying additional electrons to realize a negatively charged state with excess electrons, which can suppress the ionization of metal into the cation12. Applying this strategy to metallic NPs can be a radical solution to the naturally occurring oxidation problem of Cu NPs without artificial manipulations of the surface structure. However, it is difficult to apply the cathodic protection technique for metal NPs in ambient conditions due to non-conducting air and the impractical targeting of each NP to supply currents. To realize bare Cu NPs having excess electrons for a negatively charged surface state, we conceived a rational one-step synthesis route that enables simultaneous growth and electron transfer, in which Cu NPs are grown on the surface of a support having an extraordinary electron-donating ability. Here we successfully demonstrate non-oxidized Cu NPs in air over several months, without a surface layer of Cu oxide even on the monoatomic scale, by employing inorganic electrides as electron-donating materials. Notably, no additional post-treatments or surface passivation processes are required.According to the Cabrera–Mott theory for oxidation of metals15). The most distinct characteristic of electrides is a low work function16 that allows the effective transfer of anionic electrons to a material with a higher work function. It has been demonstrated that anionic electrons are transferred to chemically and physically adsorbed materials such as ruthenium particles and carbon nanotubes, showing enhanced catalytic and emission properties18. In this context, we grew Cu NPs on the surface of [Gd2C]2+·2e− electride with an extremely low work function (~2.8 eV)20 and successfully obtained bare Cu NPs having excess electrons that are accumulated on the surface, exhibiting ultrahigh oxidation resistance under ambient conditions.Electrides, which are used as a support in the present synthesis of bare Cu NPs, are ionic crystals that exhibit exotic properties based on excess anionic electrons occupying a structural space of ~0.4 nm  image of a Cu NP on the completely oxidized [Gd2C]2+·2e− electride is shown in Fig. d-spacing value can be observed in the cubic Cu2O {200} plane, a chemical analysis was performed by electron energy loss spectroscopy (EELS) to verify the absence of Cu oxide moieties at the surface. It is apparent that all the EELS data obtained for the Cu NPs exhibit the same energy-loss near-edge structure of Cu L-edges as that of Cu metal without white lines (arrows) of Cu oxides26, demonstrating that the surfaces of Cu NPs are non-oxidized in air  than that for Cu foil (932.6 eV), which imply a higher electron density in Cu NPs  than that of Cu metal (4.5 eV)27  and scanning tunnelling spectroscopy (STS). Both histograms obtained from KPFM and field-emission resonance spectrum measured in the STS show substantially reduced work-function values (ϕ) for Cu NPs on [Gd2C]2+·2e− (~3.2 eV on average), which are much lower than 4.5 eV for Cu metal  measurements show lower Cu 2NPs Fig. . This ne metal 4. eV27   from the surface  and the bulk , implying no oxidation at the surface area. From the high-resolution ADF STEM observations, which avoid the Fresnel fringes at the surfaces appearing in the HR-TEM images . Considering the theoretical and experimental studies on the AES data of Cu metal36, we analysed the AES data of our non-oxidized Cu NPs and compared with those of Cu metal foil having 1-nm-thick Cu2O  –EELS chemical mapping and atomic structure observations Fig. , provingn″) Fig. . The EELges Fig. , the atoges Fig. . This peu2O Fig. , Cu2O poder Fig. , bottom Q) is −2.5e , which is also verified by rigid-band approximation  and charged (Q = −2.5e) models, considering the morphology effect of Cu NPs38.Finally, using first-principles calculations (Methods), we explain how bare Cu NPs are not oxidized in air, validating our experimental results. First, we identified that the transferred excess electrons are accumulated at the surface, where the number of excess electrons on the surface Cu atoms is estimated by matching the calculated work-function value (3.25 eV) from theoretical models having different excess electrons to the average of the measured values of non-oxidized bare Cu NPs by KPFM (~3.2 eV) and STS (~3.3 eV) measurements Fig. . It was ads) on the Cu atom at the surface  and Oads infiltrates into the lattice to form an oxide   to −4.4 eV (SE)  or 2.0 eV (SE). The much smaller activation energy for infiltration at the SE is consistent with the previous observation in conventional Cu thin films, where the edges of atomic steps are most vulnerable to oxidation33. However, the overall reaction is exothermic, and the persistent presence of Oads threatens the breach of the surface Cu layers when exposed to ambient conditions. Unlike the stabilized Oads in the conventional oxidation process, the negatively charged surface makes Oads much more unstable, with adsorption energies of −1.1 eV (FS) and +0.3 eV (SE)  and 4.7 eV (SE) in the charged system compared with those of the neutral system. This relatively unstable bonding between Oads and charged Cu surfaces originates from the large number of surface-accumulated excess electrons, most of which remain at the surface even after transferring electrons to O atoms. These coexisting negative charges lead to Coulomb repulsion between negatively charged Cu surface and oxygen anions, strongly suppressing the adsorption of oxygen anions on the Cu surface and consequently prohibiting the infiltration of Oads into the bulk lattice.Figure ads into the Cu lattice increase from 3.2 to 3.6 eV (FS) and from 2.0 to 2.2 eV (SE) for neutral and negatively charged systems, respectively. It should be noted that the increase in adsorption energy by 3.6–4.7 eV in the negatively charged system predominates the infiltration process with a relatively small energy difference of 0.2–0.4 eV, indicating that the extraordinary oxidation resistance of the charged Cu NPs can be attributed to the severe suppression of oxygen adsorption on the surface. Therefore, in contrast to the general oxidation process through the dipole formation between Cu cations and oxygen anions, a substantial amount of energy is needed for the oxygen anions to be chemisorbed and to pass through the negatively charged outermost Cu layer, developing the endothermic reaction for oxide formation in Cu NPs with surface-accumulated excess electrons. The schematic of whole process is illustrated in Fig. The insets of Fig. 12 at the nanoscale. Using the present methodology, not only Cu NPs but also silver NPs maintain non-oxidized surfaces under ambient conditions  and graphite at a molar ratio of 2:1. To obtain high-quality single-phase [Gd2C]2+∙2e− and ensure homogeneity, we repeated the cooling and re-melting process three times. After the melting process, the polycrystalline [Gd2C]2+∙2e− ingot was moved into the glove box to grind the oxidized surface. Then, the fresh ingot was pulverized into a fine powder for producing Cu NPs.All the synthesis processes were carried out in glove boxes filled with high-purity argon gas  to minimize the oxidation of raw materials. Polycrystalline [Gd2C]2+·2e− powder (0.500 g) and dried Cu(ii) acetate (0.015 g) were mixed with n-heptane in an agate mortar in an argon atmosphere and transferred to a quartz tube for heating at 150 °C under a dynamic vacuum for 24 h. Conventional Cu NPs without a capping agent were prepared as a reference sample by the thermal reduction of Cu(ii) formate complex (CuF) with 2-amino-2-methyl-1-propanol (AMP)41. A CuF–AMP complex was obtained by mixing anhydrous CuF and AMP at a molar ratio of 1:2. The mixture was dispersed in methanol and heated to 180 °C for 4 h in a vacuum oven. Finally, Cu NPs were collected as the final product by drying in a vacuum. Commercial Cu NPs ) prepared by the electric explosion method42 are used . For demonstrating the practical applications of negatively charged Cu NPs, we also developed a simple and efficient method for preparing the negatively charged Cu NPs via a wet chemical process that enables the complete separation of Cu NPs from the electride. Cu precursors  were reduced by excess electrons of [Gd2C]2+·2e− electride chips (1.0 mmol) in anhydrous 1-hexanol (5 ml) at 80 °C under mild stirring, and then [Gd2C]2+·2e− electride chips were removed by using a magnet. This method uses no additional stabilizing agent, which is essential for reducing the surface energy of nanoparticles in conventional synthesis methods of metal NPs. Furthermore, the product was easily separated using a magnet, obviating the need for post-treatment. The collected Cu NPs were washed by centrifugation and tip sonication  and applied to a QUANTIFOIL 300 mesh gold holey carbon grid (Quantifoil Micro Tools) in an argon-filled glove box and then further dried in a vacuum chamber. For Cu NPs with [Gd2C]2+·2e− electride, before TEM observations, as-prepared Cu NPs with [Gd2C]2+·2e− electride samples were exposed to ambient air for 5 min to trigger the oxidation of [Gd2C]2+·2e− electride for excluding any possible charge transfer between Cu and electride support. To investigate the oxidation behaviour of Cu NPs, TEM grids prepared as above were transferred into uncapped glass vials and exposed to ambient air at room temperature during designated time periods. HR-TEM, energy-dispersive X-ray spectroscopy and STEM–EELS measurements were performed using a Cs-corrected JEM-ARM200F probe (JEOL) operated at acceleration voltages of 200 and 300 kV. For EELS mapping, the sample drift during acquisition was compensated by tracking the position of the reference atom assigned at the beginning of the acquisition. EELS measurements were carried out in the STEM mode using the same microscope equipped with a Gatan imaging filter (GIF) detector . The spectrometer was set to an energy dispersion of 0.2 eV per channel, and the energy resolution at the zero-loss peak is 0.9 eV.All the samples were dispersed in a non-polar solvent (2C]2+·2e− electride with Cu NPs was investigated by X-ray diffraction (XRD) using an X-ray diffractometer (Rigaku Smartlab) with monochromatic Cu Kα1 line (8.04 keV). The XRD patterns of the as-prepared Cu NPs on [Gd2C]2+·2e− were measured in a plastic dome-type stage filled with argon gas to avoid oxidation during measurements. Time-dependent XRD patterns were obtained by repeating the measurement with the same samples after exposure to ambient air at room temperature.Scanning electron microscopy (SEM) images of Cu NPs were obtained using a field-emission scanning electron microscope . The crystal structure of 2+·2e− electride ensures the reliability of the extremely low work function of negatively charged Cu NPs.The surface work function of the samples deposited on a Au-coated Si substrate was measured by an atomic force microscope  equipped with a sealed electrochemistry cell filled with argon gas. The KPFM experiments were conducted using Ti/Ir-coated Si probe  with a force constant of 2.8 N m–10 torr at room temperature using a commercial STM instrument . An electrochemically etched tungsten tip was used after removing oxides by electron bombardment in the ultrahigh vacuum chamber. The tips were calibrated by measuring the reference spectra on HOPG to avoid tip artifacts. A tunnelling bias was applied to the sample. Further, z–V spectroscopy was done by sweeping the tunnelling bias (V) and monitoring the tip height (z) when the feedback loop is on.STM/STS measurements were performed in an ultrahigh vacuum chamber with a base pressure of ~1.0 × 1044 with the projector augmented-wave method; a generalized gradient approximation in the Perdew–Burke–Ernzerhof form is adopted for the exchange–correlation functional45. The electron wavefunctions were expanded in a plane-wave basis set with a cutoff energy of 500 eV, and the Brillouin-zone integration for the slabs was performed using a 5 × 5 × 1 Monkhorst–Pack k-point grid46. The climbing-image nudged elastic band method47 was used to calculate the activation energy of oxidation, and all the atoms are fully relaxed until the residual forces on each atom are less than 0.01 eV Å–1. The Cu substrate was represented by slabs of six layers with the theoretical equilibrium lattice constant. A vacuum length of 25 Å was used, and the bottom two layers of the slab were fixed in their bulk positions. For surface energy calculations, the fcc Cu(111), Cu(100) and Cu(110) systems consist of (4 × 4),  is computed as follows: ΔρQ)( = ρQ)([Cu + O] – (ρ(0)[Cu] + ρ(0)[O]), where ρQ)([Cu + O] is the charge density of Cu NP with Oads on its surface and ρ(0)[Cu] and ρ(0)[O] are the charge densities of the neutral Cu substrate and isolated O atom at the position of Oads, respectively. The rigid-band approximation was adopted to verify the calculational results for correlation between the excess electrons and work functions. The work-function change of the Cu(111) slab after adopting the excess electrons can be defined as Δϕ = ϕexcess – , where ϕexcess and ϕneutral are the work functions of the excess and neutral Cu slabs, respectively, and ΔEF is the variation in the Fermi level of the system due to excess electrons. The ΔEF values were extracted using the density of states of the neutral Cu(111) slab  on a glass substrate and aged for 12 h at room temperature. A curing process was conducted at 250 °C for 60 min in an inert atmosphere. The electrical conductivities of the electrodes were measured using a four-point-probe in-line method with a current source (Keithley 6221) and a nanovoltmeter (Keithley 2182A)50. The distance between the probes is 1 mm. The volume resistivity (ρ) was then calculated using the formula ρ = C(V/I)t, where C is a geometric correction factor, V is the measured voltage, I is the supplied current and t is the thickness of the electrode determined using a stylus profilometer (Alpha-Step D-100). The geometric correction factor (C = 3.575) of the volume resistivity is applied as discussed elsewhere51.For the preparation of conductive Cu NP inks, powder of bare Cu NPs synthesized by the wet chemical method are added into the dispersion of PVP  in isopropyl alcohol. PVP acts as an adhesive at the interface between the Cu NPs and glass substrate. The mass ratio for the Cu, adhesive and solvent is fixed to 5.0:0.6:94.4. The mixture is then tip sonicated for 15 min (560 W) to ensure homogeneity. A powder of PVP-capped commercial Cu NPs is used as the control material in the same procedure. Further, 40 μl of the prepared conductive Cu NP inks were dropped onto a predefined mould (20.00 × 5.00 × 0.05 mmAny methods, additional references, Nature Research reporting summaries, source data, extended data, supplementary information, acknowledgements, peer review information; details of author contributions and competing interests; and statements of data and code availability are available at 10.1038/s41565-021-01070-4.Supplementary InformationSupplementary Figs. 1–21 and references.


The process instabilities intrinsic to the localized laser-powder bed interaction cause the formation of various defects in laser powder bed fusion (LPBF) additive manufacturing process. Particularly, the stochastic formation of large spatters leads to unpredictable defects in the as-printed parts. Here we report the elimination of large spatters through controlling laser-powder bed interaction instabilities by using nanoparticles. The elimination of large spatters results in 3D printing of defect lean sample with good consistency and enhanced properties. We reveal that two mechanisms work synergistically to eliminate all types of large spatters: (1) nanoparticle-enabled control of molten pool fluctuation eliminates the liquid breakup induced large spatters; (2) nanoparticle-enabled control of the liquid droplet coalescence eliminates liquid droplet colliding induced large spatters. The nanoparticle-enabled simultaneous stabilization of molten pool fluctuation and prevention of liquid droplet coalescence discovered here provide a potential way to achieve defect lean metal additive manufacturing. Defects induced by process instabilities in metal additive manufacturing limit its applications. Here, the authors report controlling laser-powder bed interaction instabilities by nanoparticles leads to defect lean metal additive manufacturing. The high spatial resolution stemming from the small focused beam size (about 50–100 µm) gives LPBF the capability to manufacture metal parts with complex geometries unachievable by conventional manufacturing routes2, as shown in Fig. 5. However, the focused laser heating of a powder bed creates severe process instabilities, which causes the formation of various defects11, as schematically shown in Fig. 9. The recoil pressure created by the vapor jet pushes the melt surface downward to form a vapor depression 12; the high-speed upward vapor flow of the vapor jet ejects powders and liquid droplets away to form spatters and induces ambient gas flow toward the laser beam to cause powder entrainment.Laser powder bed fusion (LPBF) uses a focused high-energy laser beam to selectively melt thin layers of metal powders to directly convert a computer-aided design model to a part13, the nonuniform energy absorption leads to nonuniform vaporization that causes nonuniform recoil pressure on the melt pool surface (liquid-gas interface). The nonuniform recoil pressure causes the fluctuation of the liquid-gas interface, which further causes the fluctuation of laser energy absorption and vapor pressure. The mutually supported energy absorption fluctuation and liquid surface fluctuation cause strong laser-powder bed interaction instabilities, e.g., melt pool/vapor depression fluctuation and gas-flow-driven spatter colliding. The vapor depression fluctuation induced liquid breakup14 from the melt pool and the gas-driven spatter colliding15 can cause the formation of large spatters (here defined as the spatter with the size larger than 100 µm).Due to the strong dependence of the laser absorptivity on the incident angle18, because it could cause critical processing faults , and defects in the printed parts . The part quality inconsistency induced by unpredictable defects is the most prominent barrier for the widespread adoption of the LPBF in various industries, especially for mission-critical applications26.The stochastic formation of large spatters is the major cause of unpredictable defect formation in the LPBF process and a big challenge for quality control15. Eliminating the stochastic formation of large spatters remains a challenge.Previous efforts on optimizing processing conditions could alter the amount of spatter, but could not eliminate the large spatter, because tuning the processing parameter cannot change the intrinsic nature of the localized interaction of laser with powder bedHere we show the elimination of large spatters by using nanoparticles to control the laser-powder bed interaction, which leads to 3D printing of defect lean samples with good consistency and enhanced properties. Our in situ high-speed synchrotron x-ray imaging experiments reveal that nanoparticles eliminate all types of large spatters by simultaneously stabilizing molten pool fluctuation and controlling liquid droplet coalescence. The method and mechanisms we discovered here provide a potential way to achieve defect lean metal additive manufacturing.28. To confirm that the large spatter elimination is not caused by the change of the vapor depression depth and is not just limited to a certain processing parameter, we conducted in situ x-ray imaging experiments in a wide range of processing parameters. The nanoparticle-enabled elimination of large spatters was observed under all the processing conditions studied. Figure We demonstrated the nanoparticle-enabled elimination of large spatters in Al6061+4.4vol.%TiC nanoparticles system . The scanning electron microscopy (SEM) images of the as-printed surface after removing the loose powders show the absence of large spatters on the surface of the Al6061+4.4vol.%TiC sample Fig. ; in contple Fig. . The loo30. A defect lean as-printed sample was successfully obtained.The elimination of large spatters leads to a dramatic reduction of defects in the printed sample. Surface profile measurements show that the surface roughness (Ra) was reduced by 90%, from 20 ± 3 µm in Al6061 to 2.1 ± 0.2 µm in Al6061+4.4vol.%TiC, and the maximum height difference  was reduced by 89% from 134 ± 20 to 15 ± 2 µm Fig. . X-ray iTensile testing results show that the as-printed Al6061+4.4vol.%TiC achieves a tensile strength comparable to, and a tensile elongation 33% higher than, that of the wrought Al6061 after T6 heat treatment; the heat-treated Al6061+4.4vol.%TiC exhibits a tensile strength of 123 ± 13 MPa (42%) higher than that of the wrought Al6061 and still maintains a reasonable elongation of 7.9% ± 0.3% Fig. . The wroTo unravel the mechanisms underlying the nanoparticle-enabled large spatter elimination, we conducted an in-depth in situ high-speed x-ray imaging study. We discovered two mechanisms that work synergistically to prevent large spatter formation.16, which is similar to the splash of water. However, in Al6061+4.4vol.%TiC, no liquid breakup (or even no liquid protrusion) was observed  is mitigated. Without sufficient liquid momentum, the inertial pressure  cannot overcome the surface tension induced capillary pressure to cause spattering.The analysis of the fluctuation of the vapor depression depth and width shows that the absence of liquid breakup can be attributed to the stabilization of the vapor depression fluctuation Fig. . We hypoThe second mechanism we discovered is that the nanoparticles prevent the liquid spatter coalescence during colliding, which leads to the elimination of the colliding induced large spatters. Powder colliding happens frequently during LPBF process due to the intensive and chaotic gas flow around the laser-powder bed interaction area. As shown in Fig. 32:σ is the surface tension of liquid Al6061, θ is the contact angle of Al6061 on TiC, c is a constant related to coverage fraction of nanoparticles on spatter surface (the ratio of projection areas of the nanoparticles to the surface area of the spatter), z is a constant related to the number of nanoparticle layers between the two liquid spatter surfaces. For closely packed double-layer nanoparticles at the interface of the two colliding droplets  because the formation of the large spatter may need coalescence of multiple spatters. Even though the very hot spatters with an initial temperature higher than 1420 K can coalesce during the first colliding, they may cool down to below 1420 K to avoid the coalescence in the subsequent colliding.The reduction of capillary pressure barrier at high temperature is critical for powder incorporation into the melt pool to form a build track. In the laser interaction zone, the powders can be heated to a very high temperature. This high temperature could eliminate the capillary pressure barrier in the melt pool region to allow the incorporation of nanoparticle-coated powders to enter the melt pool. We hypothesize that the temperature-dependent capillary pressure barrier is crucial for enabling spatter control and allowing powder incorporation into the melt pool at the same time.In summary, we discover and demonstrate the control of laser-powder bed interaction instabilities by nanoparticles through simultaneously stabilizing molten pool fluctuation and preventing liquid droplet coalescence, which leads to the elimination of large spatters and printing of defect lean samples with good consistency and enhanced properties. Nanoparticle-enabled control of laser-matter interaction instabilities provides a viable path for achieving defect lean metal additive manufacturing.The Al6061 powders (17–60 µm) were purchased from Valimet (USA). The Al6061 substrate (for x-ray imaging experiment with powder layer) and bare substrate (for x-ray imaging experiment without powder layer) were cut from the commercial Al6061 plate . The Al6061+TiC powders were prepared by planetary ball milling of Al6061 powders and TiC nanoparticles . The Al6061+TiC substrate (for x-ray imaging experiment with powder layer) and bare substrate (for x-ray imaging experiment without powder layer) were cut from the as-printed Al6061+TiC samples. The as-printed Al6061+TiC samples were fabricated by LPBF of the ball-milled Al6061+TiC powders. During LPBF, a self-designed LPBF system was utilized, which includes a continuous-wave (CW) ytterbium fiber laser , a galvo scanner , and an experimental chamber filled with argon gas. The self-designed LPBF system was used because it can handle the non-spherical powders with low flowability . More detailed information about powder morphology, material composition, and sample preparation procedure is shown in Supplementary Figs. 36. A CW ytterbium fiber laser  and a galvo scanner  were integrated to perform single track laser melting experiment (laser beam size (D4σ) of 90 µm, laser power of 208–416 W, scan speed of 0.4 m s−1). During laser scanning, the synchrotron x-ray beam (exposure time of 1 µs) penetrated through the samples. The transmitted signal was captured by a detection system at 25 kHz for characterizing spattering behavior to acquire a larger field of view, and at 50 kHz for capturing vapor depression dynamics.High-speed high-resolution x-ray imaging  was performed to capture the spatter, vapor depression and melt pool dynamics during laser scanning−1). The powder layer thickness is 100 µm. Imaging was performed at 25 kHz frame rate with an exposure time of 4 μs. The resolution is 4 μm. The view angle is 15° away from the normal direction of the powder bed surface. SugarCUBE Ultra illumination system  was used to illuminate the powder bed during the imaging experiment.A high-speed visible light camera  with microscope optics  was used to capture the spatter dynamics during LPBF process from the top view. A CW ytterbium fiber laser  and a galvo scanner  were integrated to perform single track LPBF experiment (laser beam size (D4σ) of around 250 µm, laser power of 500 W, scan speed of 0.2 m sThe nanoparticles in the as-printed Al6061+4.4vol.%TiC were characterized by SEM . To reveal nanoparticles clearly, the sample was tilted by 52°. Before SEM, the polished samples were firstly cleaned by low-angle ion milling  and then slightly etched by gallium ions . Electron backscatter diffraction  analysis was performed to characterize the size, morphology and orientation of grains in the as-printed Al6061 and as-printed Al6061+4.4vol.%TiC with the step size of 3 and 0.2 µm, respectively. The artifacts at the grain boundary  were cleaned by a grain dilation cleanup process using OIM analysis software. SEM  was utilized to characterize the distribution of TiC nanoparticle on the surface of the ball-milled Al6061+4.4vol.%TiC powder.−4 s−1. At the beginning of the test, the extensometer  was used to measure/control strain. At a strain of 1%, the extensometer was removed, and the strain rate was further read out/controlled by the crosshead. The tensile specimen with a size of 9.5 mm (length) × 4.1 mm (width) × 1 mm (thickness) and a gauge section of 3 mm (length) × 1 mm (width) × 1 mm (thickness) developed for additively manufactured metal in reference37 (MT2) was used.Tensile tests were performed using an Instron universal testing machine  with a strain rate of 2.5 × 10The T6 condition heat treatment on Al6061+4.4vol.%TiC was conducted using a muffle furnace under argon protected environment. The samples were first solutionized at 530 °C for 1 h, then quenched with water at 15 °C, finally aged at 160 °C for 18 h.38, as schematically shown in Supplementary Fig. −2 g corresponding to a sphere with a diameter of 2.2 mm using a laser power of 150 W, laser beam size (D4σ) of 250 µm, and heating time of 2 s. After melting, the linear solenoid was immediately triggered to accelerate the ring, which forces the liquid droplet to flow through the ring and introduces the initial deformation. The oscillation of the droplet after leaving the ring was captured by a high-speed visible light camera  at a frame rate of 10 kHz. The frequency and amplitude of the oscillation are used to calculate the surface tension and viscosity, respectively, as detailed in Supplementary Note The surface tension and viscosity were measured based on the oscillating droplet methodam size Dσ of 250 Supplementary InformationDescription of Additional Supplementary FilesSupplementary Movie 1Supplementary Movie 2Supplementary Movie 3Supplementary Movie 4Supplementary Movie 5Supplementary Movie 6


The refractive index contrast (∆n) for 100 kHz WGs is ~ 1 × 10–4, while for 5 MHz, ∆n increases to ~ 5 × 10–4. The WGs in the thermal regime are less effected by self-focusing and are larger in dimensions with reduced propagation losses. For the 1 MHz repetition rate thermal writing regime we report a low propagation loss WG (0.2 dB/cm) and demonstrate laser operation with slope efficiencies of up to ~ 28%.To support the growing landscape of near to mid-IR laser applications we demonstrate a range of low propagation loss femtosecond laser (FSL) written waveguides (WGs) that have achieved guided-mode laser operation in a rare earth (RE) doped lead-germanate glass. The WGs are fabricated in both the athermal and thermal FSL writing regimes using three different pulse repetition frequencies (PRF): 100 kHz ; 1 MHz; and 5 MHz . The lasing capability of Yb These WGs can realize lasers that are ideal sources for photonic networks because of their small cavity size, moderate average power (~ 100 s of mW), and long energy storage lifetimes2. FSL allows for writing embedded low propagation loss symmetrical structures within the transparent materials (glasses) by tight focusing of the ultrashort laser pulses beneath the surface of the material.Direct inscription of waveguides (WGs) into laser gain materials using a femtosecond laser (FSL) is a fast fabrication technique that allows complex photonic functionalities to be integrated into a single device7. This includes longer wavelength transmission of germanates into the infrared region, competitive thermal, chemical and mechanical strength, medium phonon energy (~ 800 cm−1), and high refractive index compared to the widely researched silicates8 and fluorides9 . Despite the above-mentioned properties, FSL-based WG writing in germanates for laser development has only been minimally investigated. Germanate is a good candidate for further research to identify the suitable FSL parameters that can introduce low propagation loss guiding structures and high laser slope efficiencies in near to mid-IR regions.Germanate glasses are fascinating hosts to achieve new laser operating regimes as this glass provides a good balance of properties required for efficient laser operation in the short to mid infrared region3+ doped lead-germanate glass GeO2–PbO–Ga2O310, however the measured propagation loss was high at 4.8 dB/cm. Operating at a higher PRF  resulted in bright guiding structures in a fluorogermanate glass possessing a low propagation loss of ~ 0.7 dB/cm11. In other work, single line WGs inscribed in a widely investigated Barium Gallo-Germanate (BGG) glass resulted in propagation losses of as low as 0.5 ± 0.1 dB/cm for WGs inscribed using 250 kHz, 70 fs laser pulses12. Stress-induced WG writing (double line) has also been achieved in germanate by a low PRF FSL (4 kHz) in Er3+ doped GeO2–PbO–Ga2O3 resulting in 2 dB/cm of propagation losses13 i.e. the propagation losses reduced by a factor of 2 with double line writing compared to single line as in10. Similar stress induced (double line) WGs include a Nd3+ doped GeO2-PbO waveguide amplifier exhibiting a propagation loss of ~ 1.75 dB/cm; and in14 the authors reported an internal gain of ~ 4.6 dB/cm for double line WGs in an Er3+/Yb3+ co-doped germanate waveguide amplifier (10 kHz PRF). Overall, it is apparent that to achieve efficient germanate laser operation reductions in WG propagation losses in germanate glasses are required.The reported studies of FSL written WGs in germanate cover the low and high pulse repetition frequencies (PRF) ranging from 1 kHz to 1 MHz. Early work reported a 1 kHz FSL (80 fs pulses) single line WG inscribed in Er2). The high linear refractive index of the germanate glass (1.8–2) is not matched with standard (n ~ 1.5) index matching liquid (e.g. oil) and oil immersive objective lenses which are designed to compensate spherical aberrations in the focal region. Furthermore, irradiating FSL pulses can exceed the self-focusing critical power limit of the glass which can occur even at low average laser power in conjunction with other FSL parameters15. This combination typically results in non-spherical and elongated structures. Currently there are no commercially available high-index matching oils and objective lenses; thus limiting the parameter space that can be explored.One of the challenges of FSL writing in germanate glass is its high 1st order (n) and 2nd order refractive index  and verify their near-IR lasing capability. The selection of this composition of GPGN glass is based on the fact that its low loss fabrication has extensively been researched for mid-IR fibers17. Employing GPGN glass for this study extends its utilization to inscribe low loss waveguides for laser applications.In this paper, we investigate a range of FSL inscribed low propagation loss single line and double line induced WGs in a rare earth doped lead-germanate glass GeO3+ is doped in the glass to evaluate the lasing performance of the fabricated WGs. We believe that this study will provide a road map towards the fabrication of low loss WGs for efficient lasing in the mid-IR. The study is conducted for both the athermal (low PRF) and thermal (high PRF) regimes. The effect of nonlinear interaction of FSL with GPGN glass and the resulting modified index region in athermal (100 kHz) and thermal (1 MHz and 5 MHz) regimes of FSL are discussed, and the parameters to achieve lowest propagation loss in GPGN glass are presented. A high slope efficiency of ~ 28% is achieved for 1 MHz double line inscribed WG which is ~ 5 times higher than achieved with FSL parameters presented in our previous study16.The aim of the study is to optimize the FSL parameters in such a way that low propagation loss guiding structures are induced in the GPGN glass. Yb3+ doped GPGN glass ingots are fabricated by a conventional melt quench technique18. The ion density of Yb3+ selected for small cavity WG laser operation is ~ 7 × 1020 ions/cm3. The 15 × 12 × 4 mm3 glass sample is cut to have parallel surfaces. The sample is polished to optical grade for WG writing, characterization, and laser demonstration.Yb19. These parameters need to be balanced with glass properties  to produce modifications that are unique to each glass composition. For writing WGs in GPGN glass a ~ 250 fs pulse width Yb3+ fiber laser (IMRA FCPA-µJewel) operating at 1047 nm (Pavg. = 2.4 W) and externally frequency doubled to λ = 524 nm is used. The WG writing setup is demonstrated in Fig. The parameters to achieve controlled and symmetric refractive index modifications in glasses depends on the careful tuning of the FSL parameters such as pulse energy, translation speed, PRF, and the laser spot size22. Moving the sample transversely with respect to the focus of the FSL produces tracks  of respective index modifications at the focus point all the way through the sample Selected writing depth into this n = 1.82 glass is a balance between shallow writing depths resulting in thermally induced cracking, while deeper focusing increases the spherical aberration. An externally set writing depth of 150 μm is chosen which results in an estimated writing depth of ~ 234 μm  structural features of the modified region; and measurement of (ii) coupling loss (CL) into the WG; (iii) propagation loss (PL) through the WG determined from the measured transmission loss (TL); and (iv) numerical aperture (NA) of the WG. For the structural characterization, the cross-sectional details of the modified region are collected via rear illuminated brightfield microscopy. To characterize the WG losses  a flexible probe beam technique is developed that can simultaneously measure CL and TL. The alignment error is estimated by taking 4 measurements of each WG. Figure 3+ doped glass does not have ground state absorption at this wavelength. A polarized beam splitter cube (PBS-104) followed by a quarter waveplate (λ/4) is used to separate the probe beam going-to and traveling-back from the glass sample Optical loss (OL) through the PBS and the λ/4 waveplate.(ii)Mirror M2 loss (ML) while retroreflecting the probe beam to ‘b’.(iii)Coupling loss (CL) which arises due to the mode mismatch between the probe beam and the WG dimensions and WG NA. A reasonable mode-match is selected based on a f = 10 mm focusing lens.(iv)Fresnel reflection loss (FL) due to the surface reflections at the air/glass/air interfaces of the glass sample.(v)Transmission loss (TL) due to the scattering, impurities, and defects in the WG.With reference to Fig. b’ from ‘e’ to allow determination of TL, independent of CL . The probe beam is retroreflected at ‘e’ via butting a R = 97% mirror to the end of the glass sample. The 3% leakage from the butted mirror has the advantage of allowing the launched mode to be monitored at ‘e’. A phosphor-coated CCD camera beam profiler  is used to measure the beam intensity of the WG fundamental mode by selecting it using the built-in aperture of the beam profiler software . CL and TL are calculated based on measured beam intensities at selected locations in the experimental layout using the derived loss equations given in detail in the supplement 1. The PL of a WG is then extracted from the measured TL using PL = TL/d, where d is the length of a WG in cm. The final set of equations through which TL and CL are calculated are given below1 = OL + ML, L2 = FL, L3 = CL + TL + FL, and L4 = 2TL + 2FL + CL + OL + ML.The probe beam is then retroreflected back to ‘16.The NA of a WG is estimated by the 1550 nm probe beam divergence angle (θ) out of the WG. Using this estimated NA of the WG, and the refractive index of GPGN glass, the refractive index modification is estimated using Figure a’ to ‘e’, then back to ‘b’, is measured under conditions where the focusing lens and glass sample is removed. OL is measured to be ~ 0.43 dB which includes 0.13 dB of M2 loss.With reference to Fig. 2 + 1)/2n). The value of n is taken from4.The GPGN glass sample is then inserted into the setup (without the focusing lens) to determine FL. The probe beam is passed through the GPGN bulk glass (i.e. the portion of the sample without WGs). The FL due to the air/glass/air interfaces of the sample is measured to be 0.73 ± 0.05 dB. The experimentally determined FL is in close approximation to the theoretical value of FL = 0.74 dB which is calculated from the refractive index of GPGN glass (n = 1.82) using FL = 10 × log  in the focal volume of a glass sample that eventually leads to permanent local modification of the refractive index23. Due to these peak powers at the focal volume, a non-linear multiphoton absorption takes place resulting in the release of an electron. Once a bounded electron is free through the initial ionization process, it then undergoes avalanche ionization where the free electron linearly absorbs the remaining laser pulse, gets excited and transfers its energy to the neighboring atom which then releases another electron once it has attained sufficient energy. The avalanche ionization process, therefore, results in exponential growth of free electrons, which generates a plasma and melts the localized volume of the sample. Due to the very short lifetime of free electrons (of the order of ps23), the electrons decay to the valence band almost immediately. This rapid cool down of the melted glass (fast quenching) thus freezes the structural modification and can lead to a change in the refractive index of the focal volume.FSL-based WG writing involves tightly focused FSL pulses that produce extremely high peak powers Pcritical) is reached, where it counters the diffraction and produces an elongated structure as the beam continues propagating deeper into the glass24. The critical self-focusing power is given by Eq.  of a related lead-germanate glass (GPLN) from25 is considered. Applying n2 = 56 × 10–20 m2 W−1, n = 1.82, and λ = 524 nm in Eq.  = κ/(ρ Cp) where κ is the thermal conductivity, ρ is the glass density and Cp is the specific heat capacity. To evaluate α for GPGN glass we consider κ = 0.7 Wm−1 K−1 (for a lead-germanate glass in27), ρGPGN = 5.61 g/cm316 and Cp = 500 J/(kg K) (for another germanate glass in28). The focusing conditions for the FSL used in this study (focusing radius ~ 0.2 μm for λ = 524 nm and NA = 1.25) leads to τ ~ 0.8 μs for GPGN glass which is comparable to the thermal diffusion time of glass mentioned in29 . As τahermal > τ, the time between the pulses is sufficiently long to carry the heat away from the focal volume before the arrival of the next pulse.For WG writing in the athermal regime, the PRF of the FSL is set to 100 kHz. The time between the successive 100 kHz laser pulses τ(i)1 in Fig. An elongated non-guiding dark structure at the focal volume, labelled as A(ii)2 above A1. A2 is referred to as WG due to its strong ability to guide light, andA bright guiding structure A(iii)3 below A1.A weak guiding structure AFor initial index modifications within the bulk GPGN glass, intermediate pulse energies of 80 nJ and 100 nJ are selected. 100 nJ appeared to be the approximate threshold pulse energy for producing observable modified regions in the glass. To write more prominent regions in GPGN glass the writing pulse energy is increased to 200 nJ, i.e. twice the writing threshold. Figure 2, WG3, and WG4, respectively.Figure 1 in Fig. 30.As discussed earlier in results “FSL  for lead-germanate, the low self-focusing threshold is most likely the cause of elongation of A1  are also reported earlier in other germanate glasses and crystals such as in31.As the Presults “” at 100 itical ~ 0 kW for 1  written with the 100 kHz FSL (pulse energy = 200 nJ) is measured to be 0.028 with an estimated refractive index change (∆n) of ∆n ~ 1 × 10–4. For 100 kHz PRF, the maximum FSL pulse energy of 200 nJ is not sufficient to induce heat accumulation in GPGN glass. However, as observed for silicate glasses in32, increasing the pulse energy above 200 nJ can result in sufficient thermal diffusion to initiate weak to modest heat accumulation even at low PRF of 100 kHz (depending on the focusing conditions). This in turn can also lead to larger ∆n and WG dimensions and needs to be explored further for GPGN glass.WG2-WG4 compared to WG1 but is also accompanied by higher PL values .In the thermal writing regime, the time between the FSL pulses reaching the focal volume is small compared to the diffusion time of the glass  is 1 μs and is almost the same as the diffusion time of the GPGN glass (τ ~ 0.8 μs). Depending on the FSL pulse energy 1 MHz is an intermediate regime between thermal and athermal. In other words, the modified region at low pulse energies in 1 MHz PRF appears similar as in athermal regime. As the pulse energy increases the athermally modified region gradually transforms to the heat accumulation/ thermal region with varying index layers.For 1 MHz PRF, the time between the incoming FSL pulses, τthermal (5 MHz) = 200 ns) is much less than the diffusion time of the glass (~ 0.8 μs). More heat starts accumulating in the focal volume and eventually produces larger volume of melted glass (compared to 1 MHz) which upon quenching results in larger heat accumulated regions.In contrast to 1 MHz, the modified region with 5 MHz is predominantly a heat accumulated region with complex structural features of bright and dark layers (discussed in next section). This is because the time between the 5 MHz PRF  of ~ 5 × 10–4 for thermally induced WGs.Maximum NA of the largest WG in the thermal writing regime increases to 0.057 for WG5-WG10 for pulse energies varying from 50 to 200 nJ. Figure (i)1 in Fig. FSL ~ 200 kW) exceeds the self-focusing threshold limit and results in the elongation of B1 (described in “An elongated dark structure (lower refractive index) at the focal volume in the middle of the modified region, labeled as B(ii)2 below B1.A bright guiding structure (WG) labeled as B(iii)3 with varying index layers above B1. The heat accumulation in B1 results in the formation of B3.An elliptical-shaped annulus structure B1 MHz single and double line FSL modified regions are inscribed in GPGN glass with pulse energies varying from 50 to 200 nJ with 50 nJ interval  below B1. As the pulse energy further increases  giving way to the formation of heat accumulated region which becomes more prominent (with no defined guiding structure) as the pulse energy increases . A low PL of ~ 0.2 dB/cm is measured for WG10 which is the lowest PL observed in GPGN glass to date.WG5, WG6, WG9 and WG10 with ~ 7%, 10%, 19% and 28% slope efficiencies, respectively 1 in Fig. An elongated dark structure at the focal volume in the center of the modified region labeled as C(ii)2.A bright circular-shaped guiding structure C(iii)3 around C2.An elliptical shaped annulus structure C(iv)4 around C3.An outer annulus structure C5 MHz PRF single line FSL modified regions are inscribed in GPGN glass with pulse energies varying from 60 to 120 nJ with 20 nJ intervals Fig. a–e. The 1. The additional heat accumulation in C1 compared to B1 [as τthermal(5 MHz) (200 ns) ≪ τthermal (1 MHz) (1 μs)] results in much higher plasma density which upon immediate quenching of the glass melt densifies around C1 and results in C3. The uniform spherical cooling of the focal volume leads to the formation of non-uniform bright and dark index layers surrounding C1  originates from the center of the modified region . With further increase in FSL pulse energy to 120 nJ, C2 represented as WG14 again turns multimode and thus PL increases again . The highest laser slope efficiency of ~ 13% is achieved for WG–4 while for thermal waveguides ∆n increases to ~ 5 × 10–4. It is concluded that the heat accumulation in thermal regime is responsible for the generation of smooth guiding structures with large ∆n. This in turn suppresses the impact of aberrations and self-focusing on the guiding region thus reducing the PL in thermally inscribed WGs in the GPGN glass. The lowest loss of ~ 0.2 dB/cm at 1550 nm is measured in a 1 MHz written double-line WG. To the best of our knowledge, the PL of 0.2 dB/cm is the lowest loss reported in a FSL inscribed WG in germanate glass and is equivalent to the PL observed in a silica glass33. Near-IR waveguide laser operation (~ 1060 nm) in the lowest loss 1 MHz written WG results in a best achieved laser slope efficiency of ~ 28% which is the highest yet reported in GPGN glass. Further improvement in slope efficiency of the Yb3+ doped GPGN glass waveguide laser can be realized by lowering the propagation losses by balancing the FSL parameters in such a way that self-focusing is minimized and symmetrical waveguide structures are inscribed.We fabricated a range of waveguides in a RE-doped lead-germanate GPGN glass using FSL in three different PRF i.e. 100 kHz , 1 MHz and 5 MHz , and verified their lasing capability in a near-IR region. The aim of the study is to optimize the FSL parameters in such a way that low propagation loss guiding structures are induced in GPGN glass for their efficient utilization in near to mid-IR laser applications. Irradiating GPGN glass with FSL pulses results in positive refractive index change in the GPGN glass in the vicinity of focal volume . The refractive index contrast (∆n) for athermal waveguides is ~ 1 × 10There is still room for further improvement in the WG laser slope efficiencies by reducing the PL in WGs with careful tuning of variable FSL parameters in such a way that induced WGs are still single mode. Moreover, the self-focusing, which arises due to the interaction of high power FSL pulses with GPGN glass, needs to be further explored  to minimize its impact on the inscribed WGs in GPGN glass .Supplementary Information.