New opportunities in transmission electron microscopy of polymers

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New opportunities in transmission electron microscopy of polymers Brooke Kueia, Melissa P. Aplanb, Joshua H. Litofskyb, Enrique D. Gomeza,b,c, a b c

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Materials Science and Engineering, The Pennsylvania State University, University Park, PA 16802, USA Chemical Engineering, The Pennsylvania State University, University Park, PA 16802, USA Materials Research Institute, The Pennsylvania State University, University Park, PA 16802, USA

A R T I C LE I N FO

A B S T R A C T

Keywords: Soft materials TEM STEM Radiation damage In-situ High resolution imaging

Recent advances in instrumentation for transmission electron microscopy have pushed the resolution limit, leading to remarkable instruments capable of imaging at 0.5 Å. But, when imaging soft materials, the resolution is often limited by the amount of dose the material can handle rather than the instrumental resolution. Despite the strong constraints placed by radiation sensitivity, recent developments in electron microscopes have the potential to advance polymer electron microscopy. For example, the focused ion beam creates opportunities for site-specific imaging, recently developed sample holders enable liquid TEM, monochromated sources lead to spectroscopy and imaging based on the valence electronic structure, and direct electron detectors minimize the required dose for imaging. Transmission electron microscopy has transformed the field of polymer science, and it is poised to do so again in the near future.

1. Introduction Transmission electron microscopy (TEM) has been an invaluable characterization tool throughout the latter history of polymer science. As early as 1959, electron microscopy elucidated the nucleation mechanism for the formation of Nylon-6 spherulites [1]. In 1985, electron diffraction coupled with high resolution electron microscopy surpassed the accuracy of X-ray diffraction in the structure determination of Cyanine Green and demonstrated the feasibility of direct imaging of organic molecules in crystals [2]. Towards the turn of the twenty-first century, cryogenic TEM revealed a new class of tough, thin-shelled block copolymer capsules [3] and toroid formation from triblock copolymers [4]. Electron microscopy has also evolved from two-dimensional (2D) projections to three-dimensional (3D) volume renderings. For example, in 2003, electron tomography revealed the coexistence of three kinds of cylindrical structures within a star ABC terpolymer in 3D [5]. More recently, cryogenic electron tomography demonstrated the self-assembly of crystalline nanotubes from amphiphilic diblock copolymers [6]. Imaging and diffraction at the nanoscale have transformed our understanding of polymer structure and often function. Analytical techniques within TEM, such as electron energy-loss spectroscopy (EELS), have also been instrumental in polymer science. EELS has been used to map the spatial distribution of water in a frozenhydrated polymer despite the complication of distinguishing small variations in water content from noise, a task that highlighted the

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necessary optimization of dose to achieve good spatial resolution [7]. Mapping of the lithium distribution within nanostructured block copolymers [8] and sulfur in mixtures of conjugated polymers [9–12] has also been demonstrated. Despite the tremendous progress that has been achieved in polymer science with electron microscopy, the resolution limit for imaging of soft materials is hindered by their inherently low contrast and their beam sensitivity. In 1971, Glaeser expressed Rose’s analysis of the limits of the image formation process in the context of resolution in the TEM:

Cd ≥

SN fDC

(1)

where C is the contrast, d is the smallest resolvable feature size, SN is the minimum acceptable signal-to-noise ratio, DC is the critical electron dose, and f is the fraction of electrons that actually enter the lens aperture and contribute to the image [13]. From this equation, it is immediately apparent that higher resolution requires an increase in electron dose. Herein lies the problem: for beam sensitive materials like polymers, beam damage can destroy the specimen such that an image with high contrast may no longer contain accurate information at the desired resolution. Nevertheless, recent developments in instrumentation, often not designed for soft materials, provide opportunities to overcome many of the challenges of polymer microscopy. The purpose of this Review is to

Corresponding author at: Chemical Engineering, The Pennsylvania State University, University Park, PA 16802, USA. E-mail address: [email protected] (E.D. Gomez).

https://doi.org/10.1016/j.mser.2019.100516 Received 19 August 2018; Accepted 23 July 2019 0927-796X/ © 2019 Published by Elsevier B.V.

Please cite this article as: Brooke Kuei, et al., Materials Science & Engineering R, https://doi.org/10.1016/j.mser.2019.100516

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sensitive materials, making spherical aberration correctors important for extending the point resolution limit in soft materials. In addition to spherical aberrations, chromatic aberrations arise from the inability of lenses to focus electrons of different energies to the same point; this is particularly problematic for thicker samples. As a consequence, chromatic aberrations can be limiting for polymeric materials when thicker samples are desired or required, such as in imaging of multiscale structures or for samples that are challenging to thin below 100 nm. Furthermore, energy-filtered TEM can be a powerful tool to image polymeric systems with heterogeneity in elemental composition, such as in salt-containing block copolymers [8] or in polymer-fullerene blends used in organic photovoltaics [12,17–20]. But, chromatic aberrations may limit the resolution in energy-filtered imaging. Chromatic aberration correctors have been demonstrated to push elemental mapping to atomic resolution in hard materials [21–23], to improve resolution at low acceleration voltages (˜80 kV), and for thick samples [24,25], but has not been sufficiently explored for soft materials. We discuss spherical and chromatic aberration correctors in the next section, and highlight a few opportunities for soft material electron microscopy in Sections 5.2 and 7.2. Nevertheless, the main limitation in resolution for soft materials comes from their sensitivity to the beam: contrast and sensitivity to beam damage are interrelated, as the latter limits the electron dose that can be used for imaging. The contrast C between domains must be larger than the minimum acceptable signal-to-noise ratio SN. Taking DC to be the maximum dose a sample can handle and d as the smallest resolvable feature size, the noise in an image is 1/ d 2DC . Thus, the product Cd must be greater than SN / DC , as shown in Eq. (1) [13,26].

summarize the accomplishments of polymer microscopy thus far while highlighting the evolution towards new opportunities made possible by technological advancements such as monochromators, aberration correctors, phase plates, direct electron detectors, specimen preparation tools, and in situ sample holders. Altogether, the field of polymer electron microscopy is poised to make significant advances in the near future. 2. Pushing the resolution limit with new instrumentation Even with a perfect lens, the resolution of any imaging system must be defined in terms of the Rayleigh criterion because diffraction of rays at the outermost collection angles of a finite sized lens results in a point being imaged as an Airy disk. Practically, in addition to this diffractionlimited resolution, resolution in the TEM is governed by inherent aberrations that come from imperfect lenses. Section 2.1 will describe spherical and chromatic aberrations in the context of imaging of soft materials, as well as the added constraint of dose-limited resolution in soft materials. Section 2.2 will describe how new advancements in instrumentation have the potential to push the resolution of soft material imaging to new limits. 2.1. Theoretical limits of the TEM The resolution of a TEM is defined as the minimum resolvable distance within the specimen and arises as a consequence of aberrations in the microscope. In TEM, the resolution is governed by the quality of the objective lens, whereas in STEM resolution is limited by the size of the probe. In both cases, this resolution is ultimately limited by a combination of spherical and chromatic aberrations in the microscope. For soft materials, beam sensitivity adds the constraint of dose-limited resolution (Eq. (1)). Spherical aberrations arise because electrons that are different distances from the optical axis are focused to different points, causing smearing in the image [14–16]. In particular, the rays scattered to high angles, which carry information about the smaller spacings in the object, are incorrectly focused because they are closer to the electromagnetic lenses and deflected more than they should be. If we assume that there is no chromatic aberration, the resolution is limited by a combination of the aforementioned Rayleigh criterion and spherical aberration which together results in a point resolution limit approxi-

2.2. Development of new instrumentation Although often not designed with polymers in mind, new instrumentation, such as monochromators, aberration correctors, phase plates, and direct electron detectors, can offer significant advantages for imaging of polymeric materials. In this section, we briefly describe these instrumentation advances in the context of hard materials and polymer microscopy, and discuss opportunities in greater detail in later sections. A monochromator is an optical component that can transmit a chosen narrow band of wavelengths of radiation; in a TEM, its purpose is to select electrons of a certain energy emitted from the source. The four types of monochromators include the single Wien filter (FEI), the double Wien filter (JEOL), the omega-shaped electrostatic monochromator (CEOS), and alpha-type magnetic monochromator (NION). At 100 kV, the energy spread of Tungsten, LaB6, Schottky FEG, and cold FEG sources are 3 eV, 1.5 eV, 0.7 eV, and 0.3 eV, respectively [27]. Monochromators can thus reduce the energy spread of the beam; for

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mated by (Cs λ3) 4 , where CS is the coefficient of spherical aberration and λ is the wavelength of the incident electron. Although we can improve this resolution by decreasing λ with higher accelerating voltages, it can be impractical, it decreases the scattering cross-section (potentially requiring thicker samples), and it is potentially harmful for beam

Fig. 1. (A) Zero loss peaks acquired in vacuum using conventional FEG (Schottky source) and with a monochromator. Reprinted from [28], with permission from Elsevier. Images and corresponding FFTs of gold particles (B) without a monochromator with an energy resolution of 0.93 eV and (C) with a monochromator with an energy resolution of 0.10 eV. The distance from the center of an FFT represents the spatial frequencies present in the image; use of a monochomator increases the achievable resolution. Reprinted from [30], with permission from Elsevier. 2

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Fig. 2. (A) Schematic of lens assembly of a hexapole aberration corrected TEM showing the path of intermediate images (axial ray, red) and diffraction images (field ray, green). The net effect of this lens assembly is the correction of spherical aberrations up to third order. Reprinted from [31], with permission from Elsevier. Comparison of holograms of a GaAs/AlAs-multilayer recorded by (B) an uncorrected CM30 Special TEM and (C) an aberration corrected Tecnai F20 TEM, showing the improved contrast achieved through aberration correction. Reprinted with permission from [32]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

generated from electrons hitting the detector are integrated over a fixed frame rate and then summed for the total exposure time in a single image. In contrast, the high-speed readout capability of direct detectors allows for counting mode, where individual electron events are detected such that noise arising from signal readout and scattered electron signal are rejected. The final image is then recorded as a stack of high signal-to-noise frames. Examples of the application of direct electron detectors in imaging techniques for soft materials such as HRTEM, 3D reconstruction, and 4D STEM are discussed in later sections.

example, reducing the spread of a 0.7 eV FEG beam to 0.18 eV (Fig. 1A) [28]. Because a large energy spread is the main contributor to chromatic aberration, monochromators can lower the information limit from 0.7 Å to 0.5 Å [29]. In addition, monochromators allow for about 0.1 eV energy resolution, which has practical applications in low-loss electron energy loss spectroscopy (Section 4) and spectrum imaging (Section 7) of polymers. Use of monochromators has vastly improved the information limit of imaging in hard materials (Fig. 1B, C) [30], and has the potential to do so in soft materials as well, both for spectroscopy and imaging. Aberration correctors have also advanced resolution limits. Multipole (non-round) lenses generate a negative spherical aberration CS to cancel the positive CS of objective (round) lenses, thus creating a net zero CS (Fig. 2). In this way, rays scattered at high angles with respect to the optical axis (these are the rays carrying information about smaller spacings in the object) are brought to the correct focus, thereby allowing access to higher frequencies and improving contrast (Fig. 2B, C). There are two types of correctors, octupole/quadrupole correctors and hexapole correctors. Section 5, which discusses high-resolution TEM, includes examples of aberration correctors in practice. Another advancement that has applications in soft materials imaging is the advent of the phase plate, a device inserted in the diffraction plane (Fig. 3A) that modulates the contrast transfer function (CTF) (Fig. 3B). Traditionally, defocus phase contrast is generated by inherent (spherical aberrations) or induced (defocus) aberrations (Fig. 3C). Nevertheless, this method results in a loss of information, particularly in the low frequency region. Alternatively, contrast can be improved through the use of a phase plate (Fig. 3D). The primary purpose of a phase plate is to change the sine-type CTF of a conventional TEM to a cosine-type function, thus improving contrast at low frequencies. The two main types of phase plates are Zernike (generates a circularly symmetric modulation pattern) and Hilbert (generates an asymmetric modulation pattern). Section 8 includes several examples of phase plates applied to soft materials. More recently, direct electron detectors have created new opportunities for the imaging of beam sensitive materials. In a traditional charge-coupled device (CCD), a scintillator converts primary electrons into photons before they hit the detector; this inefficiency limits the resolution and signal-to-noise ratio. On the other hand, a direct electron detector bypasses the scintillation step (Fig. 4). Improving the signal-tonoise ratio is especially useful for beam sensitive materials where, as discussed in Section 2.1, there is a limit to how much dose the sample can withstand. In addition, direct electron detectors have the ability to operate in two distinct modes. In linear mode, the accumulated charge

3. Advanced sample preparation techniques for new microscopy opportunities A key limitation in TEM of many polymeric materials lies in sample preparation. Thus, before discussing the applications of new microscope instrumentation to polymers, we first discuss the advanced sample preparation techniques that make polymer microscopy feasible. With a high energy (80–300 keV) electron beam, samples must maintain electron transparency and be thin enough to prevent multiple scattering in order to accurately examine composition, microstructure, and phase separation (this is a thickness on the order of 10 s to 100 s of nanometers) [36–39]. Coupled with the need for minimally deformed samples devoid of artefacts, sample preparation techniques must be able to generate a large amount of thin, uniform, and high-quality samples in a repeatable fashion [40]. 3.1. Focused ion beam One commonly used technique for creating samples suitable for the TEM is the focused ion beam (FIB). The FIB is able to prepare films in multiple ways, or it can thin existing films, removing possible artefacts and achieving the necessary sample thickness. Films produced by the FIB are electron-transparent and often show little or no effect from sample preparation [41–46]. The FIB instrument operates similarly to a scanning electron microscope, where the focused beam is capable of etching away most samples. Often, FIB instruments are equipped with a second electron beam, to enable site-selective milling with nanometer precision. For FIB usage in sample preparation, a gallium source is contacted with a nanometer-sized tungsten needle, and when coupled with a heavy electric field, ions are emitted through the tungsten tip [41,42]. Gallium is generally used as the liquid-metal ion source because of its low vapor pressure, volatility, and melting temperature [47]. Alternatively, noble gas ion sources have recently become available. As these ions travel down the ion column and raster over the 3

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Fig. 3. (A) Schematic of TEM with phase plate. (B) Moduli of contrast transfer function (CTF) with and without phase plate. Reprinted from [33], with permission from Elsevier. Lyophilized GroEL chaperonin protein imaged with (C) conventional TEM (underfocus 1960 nm) and (D) with a Zernike phase plate. Reprinted from [34], with permission from Elsevier.

Fig. 4. Schematics of a direct electron detector vs a traditional CCD (top). Comparison of sulfur elemental maps of the polymer/fullerene blend poly(3-hexylthiophene-2,5-diyl)/ [6,6]phenyl-C61-butyric acid methyl ester (P3HT/ PCBM) taken with a direct electron detector (K2) and traditional CCD (UltraScan) at identical imaging conditions (bottom). Reprinted with permission from [35].

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Fig. 5. H-bar sample preparation for TEM. Schematic diagram (a) and image (b) detailing traditional H-bar sample preparation. After the sample is mechanically cut from the bulk, it is placed on a TEM grid and further polished and thinned to reduce time under the ion beam. Then, the region of interest is milled by the FIB to an electron-transparent thickness. Reprinted from [51], with permission from Elsevier.

normal tilt at 0° and then thinned further using progressively smaller ion beams and completely cut free from the bulk specimen. After putting the lifted-out sample under an optical microscope, a micromanipulator tip can place the fully milled and electron transparent sample onto a TEM grid. The INLO technique allows for thinner milling than EXLO or conventional FIB milling, because sections produced by EXLO cannot be milled further once they are placed on the grid [50]. Nevertheless, INLO samples may be slightly less robust than EXLO, requiring great care to ensure the specimen is not damaged during handling or transfer to the TEM. Fig. 6e and f show examples of the in-situ lift-out technique used on block copolymer thin films. Although damage was apparent at the top film surface, the lamellar structure of poly(styrene-block-2-vinylpyridine) and the spherical morphology of poly(2-vinylpyridine-blockstyrene-block-2-vinylpyridine) are clearly apparent [55]. As this approach can significantly damage soft materials due to the direct impact of the ion beam on the sample, we hypothesize that staining the sample with iodine and the high glass transition temperature (near 100 °C for both blocks) were crucial to stabilize the microstructure during FIB milling. Thus, to allow for general applicability to beam-sensitive soft materials, minimization of damage due to the ion beam is warranted, as discussed below. New advances in the field of FIB milling will enable sample preparation of polymeric materials that are sensitive to the ion beam. Traditional FIB milling allows for the ion beam to directly impact the sample and can lead to thermally-induced damage or damage from impregnation of ions (e.g. Ga ions). One approach that aims to overcome these limitations is shadow-FIB milling (Fig. 7), a relatively new technique to prepare freestanding and electron-transparent samples [62–64]. This technique relies on milling the sample from the back, in order to protect the sample of interest. A shard or thin region of a film or bulk sample (on a grid or electron-transparent silicon nitride window) can be mounted onto a support and loaded into a FIB chamber upside down. The sample is then milled to the desired thickness using relatively low FIB currents from the back of the sample. Using a micromanipulator tip, the milled sample, including the grid or silicon nitride support structure, is moved to the testing chamber. To remove the silicon support and create the electron-transparent sample, electronbeam-assisted etching is used in a xenon difluoride precursor gas to prevent deposition of silicon onto the sample. This entire procedure is done in the shadow geometry, such that the ion beam never directly impacts the sample, thereby lessening damage and minimizing artefacts due to ion implantation [63,64]. This approach has been demonstrated to lead to minimal damage of block copolymers, and consequently has been used to image block copolymer films [46,49,64–66]; an example of a hexagonally packed cylindrical morphology of poly(styrene-block-

sample, they are able to precisely remove sample material by sputtering it away, allowing for nanoscale milling [48]. As stated above, the key advantages of FIB sputtering and milling are site-selectivity at the nanoscale coupled with low amounts of induced damage on the sample or specimen. Due to its nanometer-sized tungsten tip and ion column, the ability to control the milling thickness and location is high [43]. Additionally, because the incident ion beam column mills the sample at an angle that is not perpendicular to the sample, homogeneous thinning can be achieved even within rough samples [49]. This allows for the preparation of samples with specific geometries from film or bulk samples, enabling significant flexibility in sample preparation [50]. Traditional FIB preparation, sometimes referred to as the H-bar technique, was originally described by Stevie et al. in 1995 and is shown in Fig. 5 [52–56]. Although not specific to polymers, this technique is nonetheless important as it demonstrates the application of a focused ion beam. Thinning of the sample is performed before FIB milling to minimize the time under the ion beam [41]. Tungsten, or a similar metal, is deposited around the region of interest to mark and protect the sample. The FIB then thins the sample to a uniform, sub-micron thickness, alternating on either side of the sample to reduce redepositing material onto the specimen. This technique cuts a “trench” in the specimen about 5 to 20 microns in size. Because of the high control of sample location within the specimen bulk and low damage, multiple TEM samples can be prepared from the same piece of material [57]. Other FIB preparation techniques rely on extracting the region of interest from the sample, such as in the in-situ lift-out (INLO) and the exsitu lift-out (EXLO) approaches. The INLO technique is used to create a freestanding film (Fig. 6), which can then be positioned onto a copper TEM grid for sample analysis and experimentation [44,51,55,58]. This process begins by depositing a single-micron scale thick metal line to mark the region of interest for milling and to protect the sample from ion beam damage and artefact introduction during the sample preparation process. Following this demarcation, the sample is FIB milled to 10–20 microns in thickness and then extracted from the bulk sample and mounted on a TEM-compatible sample grid [59]. From the grid, the sample can then be further FIB milled to electron transparency or the desired thickness. The EXLO technique begins the same as INLO, depositing a metal protection or guide. Rather than milling the sides of the samples, the front and the rear of the sample are milled away, generally using a “stair step” milling procedure [41,42,58,60,61]. The sample is then progressively milled to single-micron thickness, and then tilted to at least 45° to cut away almost all remaining attached material holding the sample area to the rest of the bulk. The sample is then returned to the

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Fig. 6. In-situ lift out for TEM sample preparation. Images taken during the FIB milling process; the region of interest is initially selected in (a), as outlined by the white rectangle (scale bar, 5 μm). Using a focused ion beam (FIB) trenches are cut around the previously determined region of interest (b) (scale bar, 10 μm). The wedged TEM sample is extracted from the specimen using a micromanipulator tip (c) (scale bar, 10 μm). The sample is then attached to a copper mesh TEM grid (d) (scale bar, 5 μm). Reprinted by permission from [54]. TEM images of poly(styrene-b-2-vinylpyridine) are shown in (e) and (f); P2VP is shown in black in the micrographs. Reprinted from [55], with permission from Elsevier.

isoprene) copolymer film is shown in Fig. 7e. Given that thermal damage is one of the main challenges in applying the FIB to polymeric or organic materials, the use of a cryogenic stage is transformative for sample preparation (Fig. 8); the cryo-FIB minimizes damage while maintaining the precise control that comes with FIB milling [43,45,67,68]. As such, the shadow-FIB, EXLO and INLO FIB milling, and H-bar milling can be combined with cryogenic temperatures to examine a wide array of polymer and biological samples for TEM, such as cells and cellular structures, polymer films, both highly crystalline and amorphous, organic semiconductors, and organometallic compounds [49,69–73]. This is discussed in more detail in Section 3.3. One advantage of milling material away with nanoscale precision is the possibility of creating 3D reconstructions by imaging during the milling process [44]. Using a dual-beam instrument, slight milling and polishing by the ion beam alternates with imaging by the electron beam; this technique can be referred to as serial sectioning [44,61,74,75]. Once the sample is aligned at the eucentric point, such that the ion beam and the electron beam interact with the sample at the same point, the ion beam is tilted to offset it from the electron beam. Images are then gathered using the electron beam system, which can

include backscattering detectors or energy dispersive X-ray spectroscopy detectors, alternating with FIB milling in between every image or elemental map. These milled layers are on the nanometer-scale thickness, and over the course of hours, around 1000 μm3 can be reconstructed [61]. Care must be taken as to not redeposit eroded material on the sample surface, and drift correction routines must be run to minimize the consequences of ion beam drift.

3.2. Oscillating diamond knives to improve ultramicrotomy A long-standing approach to prepare polymer samples for the TEM is the ultramicrotome, a tool that can cut very thin slices of a material. Ultramicrotomes can cut electron-transparent sections, which are on the order of tens of nanometers to a micron [74,76,77]. Sharp knives, often made from diamond, tungsten carbide, or glass, are crucial to achieve such thin sections. The ultramicrotome often relies on using these sharp knives at an angle of approximately 35–45 degrees to create a small crack in the sample, and then propagating the crack through the material to create a thin section [78]. As discussed in previous reviews, ultramicrotomy can introduce surface roughness or lead to compression of the sample due to the sectioning process [79,80]. Often, optimization

Fig. 7. Shadow-FIB milling. SEM images of a block copolymer (PS-b-PI) (a–c) during shadow FIB preparation and (d) after sample is prepared for TEM. (e) TEM micrograph of the section. Reprinted from [46], with permission from Elsevier. 6

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Fig. 8. TEM micrographs of PS-b-PI block copolymer prepared in a FIB microscope at cryogenic conditions (−177 °C) and cryotransferred to the TEM. (a) High magnification TEM image with overlaid computer simulated gyroidal structure. (b) TEM image of two different regions with different thicknesses with their computer-generated projections as insets. Reprinted with permission from [73]. Copyright 2005 American Chemical Society.

Samples sectioned via ultramicrotomy are less uniform in thickness than those milled with the FIB [43]; often, knife defects are visible in sections. Furthermore, it can be easier to control the thickness of sections with FIB milling. Delamination of layers or particles within the samples from the cutting process and applied mechanical stress can introduce artefacts [77,79,87]. Also, ultramicrotomy leads to approximately millimeter site-selectivity, as opposed to near nanometer precision of FIB sample preparation [43]. A comparison of sections produced by the FIB and ultramicrotomy can be seen in Fig. 10. Nevertheless, there are also multiple advantages of ultramicrotomy. Ultramicrotomy typically requires less time to produce sections when compared to the FIB; production of samples for the TEM can be tens of minutes with the oscillating diamond knife [43,51,79]. Furthermore, the lack of exposure to an ion or electron beam during the sectioning process removes this often limiting cause of sample damage in the FIB. At the current state-of-the-art for both techniques, thinner samples are achievable using ultramicrotomy when compared to FIB milling [79,89].

of the cutting speed and cutting temperature, in addition to ensuring knife blades are sharp and free of defects, is critical to produce thin sections suitable for TEM experiments. Cryogenic ultramicrotomy, for example, is often needed to minimize deformation and enable sectioning of soft samples where the glass transition temperature is near or below room temperature, even if only one component in a multicomponent system is soft, rubbery, or liquid-like [79,81–83]. For instance, cryogenic ultramicrotomy followed by cryo-transfer to the TEM enabled the imaging of poly(methyl acrylate) (PMMA)-grafted-silica nanoparticles in poly(ethylene oxide) (PEO) [84]. One approach to minimize sample compression is to oscillate the diamond knife (Fig. 9). A diamond knife is attached to a low-voltage piezoelectric translator operating at frequencies of 20–25 kHz [77,79]. Samples can be frozen prior to sectioning if needed [77,85,86]. Cutting speeds are generally on the order of 0.1–1 mm/s [43]. After cutting, the thin sections can either be floated on water and then transferred to a TEM grid, or transferred to the grid directly. With this approach, electron-transparent samples with less compression or added surface roughness are produced, where the improvement in sample quality has been attributed to reducing the cutting angle of the knife with respect to the sample to approximately 35 ° [43,76,77,79,87,88]. The oscillating diamond knife has been demonstrated on a variety of soft materials, such as organic semiconductors for electronics, polymer membranes, nanopores, and living cells. Compared to the FIB, ultramicrotomy has various disadvantages.

3.3. Vitrification of liquid samples at cryogenic temperatures Despite the advantages of the above techniques, freezing many polymer liquids or solutions prior to sectioning through the FIB or ultramicrotomy leads to damage due to crystallization of the solvent or material itself [38]. Alternatively, liquids can be coated on a grid as a Fig. 9. Comparison of biological samples (HM20-embedded dinoflagellates) using (a) conventional sectioning and (b) oscillating diamond knife sectioning. Compression of the sample using conventional sectioning is 22.5%, while compression using the oscillating diamond knife is 7.5%. Scale bar: 100 μm. Reprinted, by permission, from [79]. Copyright 2000 John Wiley & Sons, Inc.

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Fig. 10. A comparison of TEM and STEM images with different sample preparation techniques: (a, c) FIB and (b, d) ultramicrotomy. Although sections produced by the FIB are more uniform, those made by ultramicrotomy can be slightly thinner. Additionally, no evidence of mechanical stress is seen in (a) and (c). Reprinted from [43], with permission from Elsevier.

used primarily for biological tissue fixation, but can also be used for polymers and other soft materials [90,93]. In comparison to highpressure freezing, slam-freezing better preserves the integrity and structure of layers in samples, is simpler in terms of equipment, and can make larger area samples. Nevertheless, forcefully slamming samples against frozen copper blocks can introduce some strain (flattening) and distortion during the sample preparation process. Plunge freezing instead relies on smaller amounts of sample to create vitrified films of liquids on TEM grids [110,111]. This procedure relies on a liquid solution being first dropped onto a copper grid and then plunged into a cryogen, usually liquid ethane or propane to ensure fast heat transfer (ethane and propane have higher heat capacities than nitrogen or helium, and are less able to form an insulating vapor layer) [112]. The simplicity of this approach, coupled with automatic tools to ensure the proper amount of liquid and electron-transparent films, have led to cryo-TEM imaging of a variety of samples, such as small macromolecules, proteins, and whole cells [113], including 3D reconstructions [114,115]. Cryo-TEM has also been successfully applied to a range of polymer assemblies, such as toroidal micelles, cylindrical micelles, and vesicles (Fig. 11). Disadvantages of this technique are associated with creating thin-films of liquid suitable for TEM experiments, which can induce artefacts due to the water-air interface including unwanted alignment of particles or molecules, as well as from the moderate level of control of the freezing process that can lead to the formation of ice or sample heterogeneity. A newer cryo-TEM technique is self-pressurized rapid freezing (SPRF), first published by Leunissen in 2009 [117]. Building upon the theory behind high-pressure freezing, SPRF takes advantage of the adaptability of high-pressure freezing but adds the advantages of ease of sample preparation [70]. Liquid samples are loaded into a copper capillary tube and the tubes are then sealed at both ends and submerged into liquid nitrogen at −210 °C. As the system is kept at constant volume, the increased volume of the solvent freezing (when water is the solvent) will drive an increase of pressure within the capillary, thereby achieving high-pressure during the freezing process. This sample can

thin film and the sample can be quickly vitrified to prevent crystallization, after which the frozen section (and grid) can be transferred, without allowing the sample to warm up, into a TEM equipped with a cryogenic stage. This approach is often termed cryo-TEM. Operating at cryogenic temperatures can also reduce the loss of volatile components from the sample and minimize damage of the sample from the electron beam [90–93]. Although many informative reviews of cryo-TEM can be found in the literature [94–104], we briefly highlight five vitrification techniques, high-pressure freezing, slam-freezing, plunge freezing, selfpressurized rapid freezing, and cryogenic-FIB. Perhaps the oldest of these cryo-TEM techniques is high-pressure freezing as introduced by Moor and Riehle in 1968 [105,106]. Highpressure freezing generally occurs on the order of a few thousand bar to suppress crystallization of solvents such as water [107,108]. The process begins with the creation of a sample solution that is dropped onto an aluminum husk immersed in n-hexadecane; the hexadecane acts as a pressure transfer medium [108]. Once this frozen drop of sample is within its aluminum half-shell, another half-shell of the same size is placed over the first, completely encasing the sample drop. This fullshell is then transferred into liquid nitrogen at ambient pressure for storage. The use of high pressure within the aluminum full-shell greatly reduces the nucleation rate of ice and its subsequent crystal growth, allowing for samples up to hundreds of microns to be frozen and studied [70,92]. This technique can preserve samples in native-like states; due to the high pressures, however, there is ample opportunity to introduce artefacts into the sample [90,108]. Another commonly used technique for making cryo-TEM samples is slam-freezing, such as described by Escaig in 1982 [109]. This technique is simpler compared to high pressure freezing, but operates using similar conditions; first, a drop of solution is deposited on a copper substrate and allowed to stabilize at room temperature but under humid conditions to prevent solvent evaporation [91,107]. These discs are attached to a Teflon substrate and then projected at a liquid-helium cooled 10 K copper block, thus completing the “slam-freezing” process. The frozen sample is then stored in liquid nitrogen. Slam-freezing is 8

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Fig. 11. Cryo-TEM of different types of vitrified polymer assemblies. (A) Vitrified fresh solution of poly(acrylic acid-b-methyl acrylateb-styrene) triblock copolymer micelles in 1:2 by volume tetrahydrofuran (THF) to water showing mixture of cylindrical and spherical micelles and a small amount of toroids. (B) Same solution as (A) after evaporation of THF showing predominantly toroidal micelles. Scale bar: 100 nm. From [4]. Reprinted with permission from AAAS. (C) Poly(acrylic acid)b-polystyrene cylindrical micelles. Scale bar: 100 nm. Reprinted with permission from [116]. (D) Polyethyleneoxide-polyethylethylene amphiphilic diblock copolymer vesicles. Scale bar: 20 nm. From [3]. Reprinted with permission from AAAS.

damage include knock-on damage, radiolysis, and local heating [27]. Nevertheless, previous studies on the beam damage of organic materials have different theories regarding how damage occurs in these beam sensitive materials and suggest that damage depends on many different factors besides the total electron dose, such as dose rate (flux of electron beam) [124,125], probe size [126], temperature [127], and accelerating voltage [128]. Damage has been examined extensively by measuring changes in diffraction intensities of crystalline materials [125,129–131], but changes in thickness (mass loss) and image contrast have also been quantified [124,127,131,132]. More recently, spectrometers provide an opportunity to track damage to chemical bonding, either by identifying changes in core-loss spectra associated with specific bonding orbitals or in the low-loss spectra that is associated with the valence electronic structure. Coupled with the development of monochromators, the combination of low-loss EELS and diffraction experiments will likely add new insights by revealing how chemistry and structure change with radiation damage.

then be stored in liquid nitrogen for further sectioning and imaging [117,118]. As mentioned in Section 3.1, operating the FIB with a cryogenic stage allows for milling of soft, hydrated or solvated materials. This technique was first demonstrated on biological samples in 2003 and polymers in 2005, and has since been greatly expanded for both [73,119]. The sample is created in a way similar to traditional FIB milling, taking place in a vacuum, but with a cryogenic temperature stage that is cooled once the sample is placed in the chamber, or a cryogenic-transfer stage that can insert cold samples through a loadlock stage (similarly to cryo-TEM stages). Samples are prepared through the same techniques as described in Section 3.1, and once they are thinned to electron transparency they can be either warmed to room temperature or transferred cold to the TEM. Cryo-FIB milling reduces thermal damage, local melting, or devitrification that can hinder traditional FIB sample preparation of soft (or liquid-like) samples [67,73,90,120]. As a consequence, tomography of cryogenic samples in the FIB through serial sectioning can be performed as well [121,122]. For example, cryo-FIB cross sections of vitrified liquid electrolyte were recently used to reconstruct the 3D structure of dendrites [123]. Similar to how damage from an ion beam is a limitation during FIB sample preparation, electron beam damage of polymers in the microscope also limits resolution and is discussed in the next section.

4.1. Mechanisms (knock-on, radiolysis, local heating) This review will not discuss origins in detail, but will provide a brief explanation of the aforementioned primary beam damage processes that can affect polymer samples (knock-on damage, radiolysis, and local heating). We refer the reader to other excellent resources for more indepth discussions [131,133–139]. Knock-on damage is the displacement of atoms from the crystal lattice and occurs when the energy of an incident electron is above some threshold value that scales with atomic number (about 25–80 keV

4. New insights on beam damage from low-loss EELS Despite the clear impact of beam damage on soft materials in the TEM, it is currently not fully understood. In general, mechanisms for 9

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for carbon, depends on the bonding type) [135,140]. This can generate point defects within the lattice or cause sputtering, the ejection of atoms from the surface. Low density (soft) materials, primarily made up of low atomic number elements, are particularly susceptible to this damage mechanism. Radiolysis occurs when inelastic scattering deposits enough energy to break chemical bonds. Polymers are especially susceptible to radiolysis; the high inelastic scattering cross section can lead to chemical changes. Furthermore, because most polymers are electrical insulators, accumulated charge cannot dissipate. Radiolysis damage can be localized in either the polymer backbone, leading to mass reduction, or side groups, often leading to charged species, free radicals, and crosslinking. Another common source of radiation damage in polymers is local heating caused by phonons (lattice vibrations). Polymers are thermally insulating, with thermal conductivities of about 0.1-0.5 W/mK, and will experience a more significant temperature rise than ceramics (1–200 W/mK) or metals (10–400 W/mK) [141]. In addition to thermal degradation, if heated above the glass transition temperature (Tg), chains will start to flow and the structure of the sample will change. Fig. 12 illustrates the relative time scales of the different beam damage mechanisms. Electronic processes occur on the order of femtoseconds and chemical reactions occur on the order of nanoseconds. In insulating materials, radiolysis is typically 103 – 106 times faster than knock-on damage [138]. Molecular vibrations, which lead to local heating, occur on the femtosecond timescale [142]. Assuming a diffusion coefficient of 106 cm2/s (similar to that of organic molecules in butyl rubber), molecules take about 10 ns to diffuse 1 nm (about the lattice spacing of polymer crystals). Thus, although knock-on damage can disrupt the lattice immediately, damage from radiolysis often relies on solid-state diffusion prior to apparent disruption of the structure.

Table 1 Summary of reported critical doses for polymers. Polymer

Dc (e−/nm2)

PE

2031 (2000 kV) 268 (100 kV) 749 (125 kV) 936 (500 kV) 1498 (1000 kV) 4484 (2000 kV,100 K) 100 (200 kV) 624 (100 kV) 1248 (100 kV) 2809 (100 kV) 1450 (100 kV, 100 K) 6200 (200 kV) 620 (120 kV, 77 K) 3745 (80 kV) 31,208 (80 kV) 62,400 (120 kV, 127 K) 1123 (120 kV, TEPD) b 8000 (200 kV) 936 (100 kV) 599 (100 kV) 31,208 (80 kV) 125 (80 kV) 374 (80 kV) 3745 (80 kV) 1872 (80 kV) 3200 (80 kV) 33,000 (80 kV) 10,800 (300 kV,100 K) 2800 (80 kV) 4000 (80 kV)

PEO POM Nylon-6 PEEK film PET PMMA

PS iPS PVP PBT PCTFE PC Collodion

Formvar RR P3HT

RRa P3HT PGeBTBT

a

Signal monitored

Reference

Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Diffraction pattern Low-loss EELS (5-10 eV) EELS (C K-edge) EELS (O K-edge) EELS (C K-edge) Low-loss EELS (7 eV) Diffraction pattern Low-loss EELS (˜10-40 eV) Diffraction pattern EELS (Cl edge) EELS (O edge) EELS (N edge) EELS (O edge) EELS (C edge) EELS (O edge) Low-loss EELS (˜ 2-4 eV) Diffraction pattern Diffraction pattern Low-loss EELS (˜ 2-4 eV) Low-loss EELS (˜ 1-3 eV)

[145] [130] [130] [130] [130] [145] [130] [130] [130] [130] [146] [147] [148] [149] [149] [150] [151] [152] [153] [154] [149] [149] [149] [149] [149] [18] [18] [155] [18] [18]

PE: polyethylene, PEO: poly(ethylene oxide), POM: polyoxymethylene, PEEK: polyether ether ketone, PET: polyethylene terephthalate, PMMA: poly(methyl methacrylate), PS: polystyrene, iPS: isotactic polystyrene, PVP: poly(vinyl pyrrolidone), PBT: poly(butylene terephthalate), PCTFE: poly(chlorotrifluoroethylene), PC: polycarbonate, RR P3HT: regioregular poly(3-hexylthiophene-2,5-diyl), RRa P3HT: regiorandom poly(3-hexylthiophene-2,5diyl), PGeBTBT: poly[(4,4′-bis(2- ethylhexyl)dithieno[3,2- b :2′,3′- d]germole)2,6-diyl-alt-(2,1,3- benzothiadiazole)-4,7-diyl]. a Operating voltage listed in parentheses. All measurements performed at room temperature unless otherwise stated. b TEPD = total end point dose.

4.2. Summary of reported critical doses for polymers Accurate experiments require the effects of beam damage to be carefully monitored. This is necessary to ensure the data accurately describe the properties of the sample and are not a result of the imaging procedure. Measuring and reporting a critical dose for damage is necessary for accurate imaging of polymer morphology. To ensure the image is not of a damaged sample, it is imperative the critical dose not be exceeded during imaging. Historically, beam damage in polymers was monitored mostly by the loss of a diffraction peak, such that beam damage has been mostly quantified for semicrystalline polymers. For uncorrelated damage events, the signal corresponding to the presence of pristine material (e.g., diffraction peak intensities) decays exponentially with accumulated dose. The critical dose (Dc), typically reported as electrons/area, is then the inverse of the exponential decay constant and is taken as the radiation dose above which the material will be significantly changed (intensity of measured signal, such as diffraction peak intensity, drops to 1/e of its initial value), as shown in Eq. (2).

D I = Aexp ⎛− ⎞ + Ib ⎝ Dc ⎠ ⎜

⎟

(2)

A is an exponential prefactor and Ib is the background intensity [143]. Thus, the critical dose can be calculated by monitoring the peak intensity in the diffraction pattern (I) as a function of electron dose (D) [144]. Dc is material-specific and we propose that it can even vary within the same material depending on the type of damage measured (damage to crystal vs damage to the chemical structure). Low-loss EELS provides an alternative method to quantify beam damage. The critical dose can be determined by monitoring the decay of low-loss EELS peaks as a function of dose. EELS spectra will quantify damage of the electronic structure whereas the diffraction pattern will quantify loss of crystallinity due to beam damage. Importantly, with EELS, beam damage of amorphous polymers can be monitored. A summary of published critical doses (Dc) is reported in Table 1. The values in Table 1 were obtained by various groups over several decades. One must exercise caution when interpreting the presented values, as Dc can vary significantly with experimental conditions including accelerating voltage, temperature, and dose rate (the latter is often not reported). Furthermore, the electron dose required to completely destroy the signal from a diffraction pattern or EELS peak, or total end point dose (TEPD), may be used to quantify beam damage instead of Dc. Nevertheless, one trend evident in the table is that polymers with increased unsaturation appear to be more resistant to

Fig. 12. Different beam damage mechanisms occur on different timescales. In insulators, radiolysis is typically the most significant. 10

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decreased, another peak at ˜5 eV appeared (Fig. 13b). This peak was assigned to degraded phenyl rings. When the carbon K-edge of the EELS spectrum is measured beyond the critical dose, a broadening of the C=C 1s to π* transition is observed while the peak corresponding to the CeH 1s to σ* transition does not change. Based on the damage being manifested in signals corresponding to the phenyl rings in both the EELS and low-loss EELS spectra, a crosslinking mechanism is suspected. Siangchaew et al. performed similar low-loss EELS beam damage experiments in STEM mode to examine the 7 eV peak in PS films [126]. Unexpectedly, they noticed that as the size of the probe decreased and dose rate increased, the low-loss peak was significantly less susceptible to beam damage. They proposed that fast secondary electrons emitted roughly perpendicular to the incident beam travel laterally within the sample. The seemingly high resistance to beam damage was because the fast secondary electrons would travel outside the probe area and thus, not be monitored by the technique. To support this idea, they performed scans at systematic distances away from the initial probe site. They found that at 80 nm away from the site, each pixel scanned produced the same peak intensity, suggesting the probed material had not been previously damaged. When the distance from the initial probe site is decreased to 5 nm, the intensity of the 7 eV peak was significantly reduced, suggesting that the probed material had been previously damaged. Thus, when the probe is small, damage is delocalized beyond the sample region being probed by the beam, leading to lower apparent radiation damage. In a more recent study, Egerton et al. supported their theory using Monte-Carlo simulations of fast secondary electron trajectories, although the theoretical calculations also indicate that fast secondary electrons cannot be the only explanation for the high resistance to damage that is apparent with small probe sizes [156]. To achieve the observed beam damage resistance, fast secondary electrons would need to travel at least 500 nm away from the probe. Fast secondary electrons have a mean transport range on the order of 10 nm in organic solids. Thus, fast secondary electrons cannot be the only factor contributing to the high resistance of PS. It is suggested that delocalization of inelastically scattered electrons may also contribute. Thus, the mechanism for the large Dc with small probes and radiation damage away from the probe site remains unresolved; we speculate that diffusion of reacting species away from the probe may also contribute to the large Dc value observed. Conjugated polymers, often composed of conjugated backbones and solubilizing alkyl side chains, have also been examined using low-loss EELS. These materials make up an important class of polymers that can be used as the active layer in many organic electronic devices. Delocalized π-electron densities impart semiconducting properties, as well as absorption features in the low-loss region (1–10 eV). Thus, lowloss EELS imaging has the potential to significantly enhance imaging of

damage from the electron beam. The Dc values measured at ambient temperature with a 100 kV accelerating voltage for PE, POM, PBT, Nylon-6, and PEEK are 268, 624, 936, 1248, and 2809 e/nm2, respectively. PE and POM are both completely saturated polymers and are reported to have lower Dc values (on the order of 100 s of e/nm2). Nylon-6 is a polyamide which incorporates a carbonyl group within its molecular structure, and PBT and PEEK both incorporate phenyl rings within their molecular structures. Nylon-6, PBT, and PEEK have relatively higher Dc values than PE and POM. In general, polymers with more degrees of unsaturation, and especially polymers with aromatic moieties, may be more radiation resistant. Due to enhanced delocalization, these polymers may be able to distribute energy deposited by the electron beam more efficiently. As highlighted above, tabulated values can still be useful in extracting general trends, even though precise characterization of beam damage is sensitive to experimental conditions.

4.3. Monochromated low-loss EELS beam damage experiments EELS spectral imaging has significant potential to investigate systems not easily measured by conventional TEM, particularly polymers. Because polymer degradation is the main limitation of EELS, understanding radiation damage is very important [18,126,135,147,148,150,156–159]. Unlike when beam damage is monitored using the diffraction pattern, low-loss EELS data can reveal useful information regarding the breakdown of the chemical and electronic structure. Beam damage in the low-loss region was first measured by Siangchaew and Libera using low-loss EELS to examine the decay of the π-π* transition in polystyrene using a field emission source with a full width at half max for the zero-loss peak of 1 eV (Fig. 13) [160,161]. Degradation of the π-π* transition is attributed to chemical effects caused by differences in bonding and conjugation as the material is damaged. The authors conclude that radiation damage studies provide important information that defines appropriate conditions for imaging polymeric samples. Furthermore, the authors propose that the scientific community would benefit if more EELS studies of polymers included quantitative assessment of radiation stability specifically through the creation of a database focused on spectral fingerprints and radiation stability of polymers. Since these reports, beam damage in polystyrene has been studied extensively using low-loss EELS; it is convenient to study the peak at an energy-loss of 7 eV, corresponding to the π- π* transition of the phenyl ring. Varlot et al. used EELS to assign the electronic transition energies at the carbon K-edge for the C]C 1s to π* and CeH 1s to σ* transitions [150]. Using the peak in the low-loss region (7 eV), corresponding to the π- π* transition of the phenyl ring, the authors calculated a critical dose for polystyrene of about 62,000 e/nm2. As the π- π* peak intensity

Fig. 13. Examining beam damage in PS using low-loss EELS as a function of electron dose. (a) The exponential decay of the π- π* transition in polystyrene as a function of electron dose. Reprinted with permission from [162]. (b) The low-loss spectrum of polystyrene for different electron doses. As the dose increases from a) 300 e/nm2 to b) 300,000 e/nm2 to c) 1,000,000 e/nm2 the peak at 7 eV energy-loss decreases and a new peak at 5 eV energy loss increases in intensity. Reprinted with permission from [150].

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Fig. 14. Quantifying beam damage in rr-P3HT using monochromated low-loss EELS. (a) rrP3HT low-loss EELS spectra at various different electron doses. (b) Integrated low-loss EELS intensities as a function of electron dose. As electron dose increases, the intensity of the peak at 2.6 eV gradually decreases and shifts to higher energy, plateauing at 3.2 eV. Reprinted with permission from [18].

5.1. Beam damage minimization approaches

semiconducting polymers, but understanding beam damage will be essential. For materials with excitations near the visible region, the use of a monochromated electron source is crucial to minimize the background from the zero-loss peak. In a recent study, the critical dose of regioregular poly(3-hexylthiophene-2,5-diyl) (P3HT) was quantified from both π-π* transition peak intensities and diffraction intensities using monochromated 80 kV electrons with an energy spread of 0.15 eV [18]. Because of the minimal energy dispersion of the probe, the authors were able to extract the electron absorption spectra near the bandgap at 2 eV. Degradation of the electronic structure of P3HT as a decrease in peak intensity was measured, and a gradual shift to higher energy of the peak absorption, from about 2.6 eV to 3.2 eV, was observed (Fig. 14). It was found that the low-loss EELS critical dose was 3200 e/nm2 compared to the critical dose from the diffraction pattern of 33,000 e/nm2. Based on the large difference in magnitude of critical dose measurements, the authors conclude that the valence electronic structure is damaged before the crystal structure. This is useful information for measuring the morphology of semiconducting polymers; the optimum acquisition scheme can be determined depending on what characteristics are being measured. Examples of taking advantage of the low-loss spectra for imaging are highlighted in Section 7.3. In a similar study, the beam sensitivity of organic materials was measured with STEM-EELS and compared to data from variable angle spectroscopic ellipsometry (VASE). Because STEM-EELS has high spatial resolution but causes damage while VASE has low resolution but does not induce damage, comparison of the complex dielectric function obtained from the two experiments allows for optimization of acquisition parameters for STEM-EELS [163]. Dose-dependent studies have revealed much about the mechanisms behind radiation damage in the TEM, but a complete picture remains elusive. Often, lower accelerating voltages are useful to reduce knockon damage, although it remains unclear whether greater ionization damage will be a detriment to low-loss EELS measurements. Reducing the probe size appears to significantly reduce the sensitivity to radiation dose by allowing damage to diffuse away from the probe, suggesting interesting opportunities with focused-probe experiments (e.g., analytical techniques such as STEM-EELS). Overall, measuring the limits of dose tolerance and developing measurement approaches to minimize damage is crucial to push the resolution limit, as discussed in the next section.

Starting with the relationship between DC, SNR, C, and d from Eq. (1) and taking into account the detective quantum efficiency DQE (a measure of the signal and noise performance of a detector), the smallest resolvable feature size can be expressed as [156] −1

−1

d = (DQE )

2

SNR DC ⎛ ⎞ C ⎝ e ⎠

2

(3)

Assuming DQE = 0.2 (typical for a CCD detector at the Nyquist spatial-frequency limit), SNR = 5 (from the Rose criterion), and C = 0.1 (typical for unstained polymers) as was done by Egerton et al. [156], we can estimate the dose required for several desired resolutions. These values can be found in Table 2 below. We observe that in order to achieve the resolution required to resolve chains within crystals (d ˜ 0.4 nm) a dose of nearly 80,000 e/nm2 is needed, and to resolve atomic positions (˜ 0.1 nm resolution) the required dose is 1,250,000 e/ nm2. Clearly, there is a need for minimizing beam damage in order to achieve these high doses. As discussed in Section 4, beam damage studies of organic materials to date have differing theories regarding how damage occurs in these beam sensitive materials. For example, EELS of the 7 eV π-π* peak of polystyrene reveals that π bonding is more stable at higher dose rates [164], but EELS experiments on thin films of collodion show first an increase then decrease in stability with increasing dose rate [124]. Diffraction experiments on P3HT also show an increase then decrease in critical dose with increasing dose rate, which is explained by radiolysis followed by the slow diffusion of a reacting species [125]. Conversely, another diffraction study of P3HT/PCBM suggests that there is no dose rate dependence on beam damage [165]. Studies on the relationship between beam sensitivity and voltage do not follow a clear trend either: beam sensitivity increases as voltage decreases near 100 kV, but the trend is reversed at very low voltages [128]. The effect of other variables such as temperature and the presence of water or oxygen during sample preparation have also been shown to affect stability [165]. Further beam damage studies would be valuable in working towards a unified theory for the damage mechanisms in polymers. Despite an incomplete understanding of how beam damage occurs, some strategies for minimizing damage are known. It is helpful to know how knock-on damage, radiolysis, and local heating affect the polymer Table 2 Required dose in electrons per squared nm to achieve desired resolution calculated from Eq. (3).

5. Making HRTEM possible for polymers The resolution limit of organic materials is typically set by the dose the sample can tolerate in the electron microscope. As such, details matter, in the sense that any mitigation of radiation damage can result in pushing the resolution limits beyond what is currently achievable. Here we discuss how beam damage can be minimized, the evolution of HRTEM of polymers thus far, and how aberration correctors and direct electron detectors offer new opportunities for the future. 12

d (nm)

Required dose (e/nm2)

1000 100 10 1 0.4 0.1

0.0125 1.25 125 12,500 78,100 1,250,000

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Fig. 15. (a) Bright field TEM image of regioregular P3HT grown by directional epitaxial solidification. (b) Low dose HRTEM image. (c) Schematic of P3HT chain packing along the b and c axis. (d) Schematic of flat-on and edge-on lamellae corresponding to HRTEM image. Reprinted with permission from [168]. Copyright 2009 American Chemical Society.

rise to an exit wave ψx. The Fourier transform of ψx is ψk, where k is the scattering vector. The role of the objective lens is to multiply ψk by a contrast transfer function T = sin χ(k), and then to inverse Fourier transform ψkT to give the image wave ψi. Thus, sin χ(k) gives the phase changes of diffracted beams with respect to the direct beam as

being imaged, as it can help in choosing experimental parameters to minimize damage. For example, because the scattering cross section can Ze be given by σ = πr 2 for r = Vθ where Z is the atomic number, e is the charge of an electron, and V is the accelerating voltage, a sample that is most susceptible to knock-on damage should be imaged with a lower accelerating voltage whereas a sample that is most susceptible to radiolysis should be imaged with a higher accelerating voltage. In addition to choice of accelerating voltage, the common practice of using cryogenic conditions for biological samples is also transferrable to polymeric materials; a recent study demonstrated that cryogenic conditions increased the beam stability of P3HT/PCBM [165]. Cryogenic conditions could also minimize the effects of local heating from the electron beam. As mentioned previously, beam damage studies as a function of dose rate reveal that higher dose rates (smaller beam size) could also minimize damage. While this might be an artefact from under-sampling of damaged material, dose rate and beam size optimization could be applicable in STEM mode where rastering the probe quickly enough may be able to outrun damage. Additionally, sparse sampling STEM techniques coupled with reconstruction algorithms could be a useful approach for beam sensitive materials [166]. Although some combination of the aforementioned approaches may enable imaging experiments at about 105 e/nm2, thereby resolving the separation of polymer backbones in many materials, achieving doses of 106 e/nm2 and atomic resolution seems out of reach. As discussed in the next section, perhaps the application of aberration-corrected microscopes and direct electron detectors, in combination with exceptionally radiation-hardy soft materials, will enable near atomic resolution imaging.

χ (k ) = πλΔfk 2 +

1 πCS λ3k 4 2

(4)

where λ is the wavelength of the incident electron, k is the spatial frequency, Δf is the defocus, and CS is the coefficient of spherical aberration, which describes the quality of the objective lens. As a consequence, negative defocus values can be used to offset the positive spherical aberrations inherent to the lens. Nevertheless, sin χ(k) is positive only for intermediate values of k, meaning that in this region all information is transferred with positive phase contrast and can be easily interpreted, but once the function crosses the k axis (this is the point resolution) and begins to oscillate strongly, interpretation becomes more convoluted. The implication here is that for an uncorrected microscope, the resolution limit follows the aforementioned point re1

solution (Cs λ3) 4 (Section 2.1). To overcome this resolution limit, non-rotationally symmetric lenses can be used to produce negative aberrations that cancel out the positive aberrations. Eq. (4) is a simplified description of the contrast transfer function; in reality, it should contain higher order terms (C5, C7, etc.), with each order corresponding to electrons scattered at higher frequencies. Thus, whereas an uncorrected microscope is C3-limited, a C3corrected microscope is now C5-limited, thereby allowing for higher frequency information to be obtained [167]. Aberration correction also reduces delocalization in the image, which minimizes the spread of contrast from nonperiodic features in the specimen [31]. At the highest frequencies, chromatic aberrations, rather than spherical aberrations, become the limiting factor. The large lattice constant of polymer crystals facilitates imaging without aberration correctors. The first published polymer TEM image with lattice resolution was in 1969, when Bassett and Keller imaged the 1.8 nm (100) spacings in beta-PPX using two-beam dark-field (DF) imaging [169]. Since then, many polymers have been imaged with HRTEM, as listed in detail in previous reviews [170]. In 2009, the

5.2. Aberration corrected microscopes and direct electron detector for low dose HRTEM We present a brief overview of the physics of image formation in the TEM to motivate the advantage offered by aberration correctors, and to highlight the potential opportunities for imaging of polymers. Consider the incident electron beam ψo and the electron potential function of the sample φ, such that as the beam interacts with the sample, it will give 13

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Fig. 16. (a) HRTEM image of rr-P3HT:PCBM showing the (100) fringes of rr-P3HT crystals and the rr-P3HT/PCBM phase boundaries (red lines). (b) HRTEM of pure rr-P3HT showing (100) fringes of lamellar crystals. Images were taken on a Cs-corrected FEI Titan operating at 300 kV with a low-dose system. Reprinted with permission from [19]. Copyright 2011 American Chemical Society. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 17, comparison of expected imaging results and experiments suggest that atomic-scale information is revealed. Thus, although earlier work with a conventional electron microscope relied upon electron diffraction to reveal molecular information about polyvinylidene fluoride (PVDF) electrospun nanofibers [172], imaging with the CScorrected TEAM 0.5 revealed features that are interpreted to be rows of CF2 groups [173]. The HRTEM image shown in Fig. 17 was acquired with 2000 e/nm2 after exposing the sample to a dose of 170,000 e/nm2. Previous work with electron diffraction of chain folded crystals of PVDF had demonstrated a loss of diffraction spots at about 200 e/nm2 (originally reported as 38 C/m2) [174]. The authors of the study that produced the images in Fig. 17 suggest that their observed larger radiation tolerance is due to the presence of heavy metals (Pt, used to enhance contrast) in earlier work, because these heavy atoms could cause secondary radiation events that promote damage. Nevertheless, achieving near angstrom resolution with 2000 e/nm2 for a single image suggests that aberration-corrected microscopes are moving beyond the limitations outlined in Table 2, likely due to a higher contrast and lower acceptable signal-to-noise ratio than our assumptions used in Eq. (3). The combination of direct electron detectors with aberration correctors has the potential to push resolution limits for polymers even further. The advantage of direct electron detectors is that they have an improved modulation transfer function (MTF), which describes the effect of an electron being detected as signal in multiple pixels, as well as better detective quantum efficiency (DQE), which describes how the detector affects the signal-to-noise ratio in an image [175]. Thus, the application of direct electron detectors will reduce the dose required to image at a given resolution when compared to traditional CCD detectors. In addition to better MTF and DQE, it has also been demonstrated that direct detectors operating in counting mode, as opposed to integrating mode, offer improved resolution as well. For example, Stach et al. took advantage of the Gatan K2-IS direct electron detector operating in counting mode to achieve aberration corrected HRTEM images of P3HT by taking multiple images that can be drift corrected prior to summing (Fig. 18) [176]. Another technique with growing potential due to direct electron detectors is known as scanning nanodiffraction, also called 4D STEM. When the electron beam is focused into a probe, a convergent beam electron diffraction (CBED) pattern is formed in the far field. This pattern is rich in crystallographic and scattering data as well as information regarding thermal vibrations and energy losses and is particularly amenable to soft materials because of the low dose required for a large amount of diffraction information. Traditional STEM imaging, however, simply uses a monolithic detector such as an annular dark field (ADF) or bright field (BF) detector, throwing away most of the diffracted signal information. By placing a high-speed pixelated direct electron detector in the far field, every pixel of the CBED pattern can be recorded at each probe location with millisecond dwell times, resulting in a 2D CBED pattern (reciprocal space) at each point of a 2D STEM image (real space) — hence the name 4D STEM [177]. The recent

semicrystalline structure of regioregular poly(3-hexylthiophene-2,5diyl) (rr-P3HT) thin films grown by directional epitaxial solidification were imaged with HRTEM by Brinkmann (Fig. 15) [168]. Crystalline order was aligned through directional epitaxy, and images were enhanced using a Fourier filter that emphasized frequencies corresponding to the lattice spacing. While HRTEM has been accomplished on a number of polymer systems before the advent of aberration correctors, aberration correctors have improved resolution and opened up opportunities for imaging of polymers. For example, Drummy et al. combined EFTEM with low dose HRTEM on a CS-corrected FEI Titan to image rr-P3HT crystals and their orientation within rr-P3HT-rich domains of rr-P3HT/PCBM blends (Fig. 16) [171]. Distinct regions corresponding to rr-P3HT and PCBM crystallites are apparent without any image processing, in contrast to previous work on HRTEM of rr-P3HT [168]. Although the improvement may be due to a better detector, we speculate that the image quality is also enhanced by extending the frequency at which the CTF crosses zero using Cs correctors, given that finite defocus is needed to enhance phase contrast. This could effectively increase the signal-to-noise ratio. Recent work on a CS-corrected microscope has also imaged polyvinylidene fluoride (PVDF) fibers at high resolution. As shown in

Fig. 17. (A) HRTEM image of a segment of a thin nanofiber of PVDF compared with (B) a simulation of PVDF. Images are overlaid by a stick model in which carbon-carbon bonds are green, CH2 groups are white, and CF2 groups are red. Reprinted with permission from [173]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) 14

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crucial parameters such as the contrast and signal-to-noise ratio and assumes that a traditional CCD is used. We propose that revisiting the resolution limits for soft materials in modern electron microscopes is needed, and that new instrumentation likely enables near atomic resolution for many polymers. 5.3. Improved modeling software and computational tools Image interpretation can often be a challenge for materials with significant disorder, as is often the case for polymers. Fig. 17 highlights one example of how computational tools can aid in analysis of HRTEM images. Another example is a recently published software called GRATE (GRaph based Analysis of TEM images) which converts HRTEM images of polymer films into easily interpreted line drawings from which microstructural information can be extracted [182,183]. Software has also been helpful for analyzing 3D reconstructions, where a discrete algebraic reconstruction technique (DART) [184–186] has been demonstrated to be a useful tool in the interpretation of STEM-tomography tilt series of P3HT/PCBM (Fig. 20) [187]. In a typical tomography reconstruction procedure, many gray-scale levels are allowed for every voxel, such that eventual segmentation (the assignment of dark or bright domains to different materials) is done manually. DART, on the other hand, automatically segments reconstructions through the use of constant gray-level assignations to each material. The addition of this constraint adds information to the reconstruction, and thereby improves reconstructions for a given exposure or number of images in a tilt series. Besides software that is used for analysis post-data acquisition, recent work on automating sample preparation, image acquisition, and image analysis all at once has enabled high-throughput TEM capable of generating phase diagrams of block copolymer amphiphiles [188]. In addition to advancements in software, developments in instrumentation and techniques are clearly warranted to minimize the effects of radiation damage in the TEM. The next section overviews some approaches to address this central challenge in polymer electron microscopy.

Fig. 18. HRTEM images of P3HT domains. Aligned summation of 20 images, each acquired at 12,500 e/nm2. Images were taken on an aberration corrected microscope using a Gatan K2-IS direct electron detector operating in counting mode.

development of an electron microscope pixel array detector (EMPAD) has allowed for ptychographic reconstructions capable of revealing deep sub-angstrom resolution (0.39 Å) in 2D and beam-sensitive materials [178]. Although various different types of 4D STEM experiments can be done, including position-averaged convergent beam electron diffraction (PACBED), virtual dark field imaging, fluctuation electron microscopy, strain measurements, and phase contrast imaging methods such as ptychography and MIDI STEM (see discussion in Section 7.2), 4D STEM has only been demonstrated on polymers very recently. In earlier studies, 4D STEM was used to map islands of P3HT in a PS matrix [179], whereas more recent studies have used 4D STEM to create orientation maps of π-π stacking in a small molecule, 7,7′-(4,4bis(2-ethylhexyl)-4H-silolo[3,2-b:4,5-b′] dithiophene-2,6-diyl)bis(6fluoro-4-(5′-hexyl[2,2′-bithiophen]- 5-yl)benzo[c][1,2,5]-thiadiazole), and a conjugated polymer, poly[2,5-bis(3-tetradecylthiophen-2-yl) thieno[3,2-b]thiophene] (PBTTT) (Fig. 19) [180]. Pair distribution function analysis of 4D STEM data has also been proposed as a method for characterizing short and medium range order in aperiodically packed organic molecules [181]. Aberration-correctors and direct electron detectors are pushing the resolution in images of polymers, as shown by the examples highlighted in this Section. Our simple calculations of the dose required for imaging at specific length scales shown in Table 2 relies on rough estimates of

6. Low-dose techniques for polymer microscopy Although the beam sensitivity of soft materials appears prohibitive for high resolution imaging as exemplified in Table 2, advances in instruments and techniques aim to push the limits established by Glaeser. This Section highlights a few examples of instrument and technique development that attempt to move beyond Eq. (1); another example, single-particle reconstructions, is described in Section 8.1. 6.1. Monochromators for low-dose imaging Most modern microscopes are designed to maximize flux. As a consequence, controlling the dose rate to achieve low-dose imaging for soft materials can be challenging. Monochromators significantly reduce

Fig. 19. π-π stacking orientation maps obtained from 4D STEM. (a) Orientation map of as-cast PBTTT thin film showing no mesoscale order. (b) Orientation map of annealed PBTTT thin film showing enhanced ordering with domains several nanometers in size. Reprinted with permission from [180]. 15

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Fig. 20. Volume reconstructions of as-cast and thermally-annealed blends of P3HT and the endohedral fullerene Lu3 N@C80-PCBM. Reproduced with permission from [187]. Fig. 21. (A) High-dose (2000 e/Å2s) TEM image of carbon nanothreads with FFT (inset) showing amorphous structure despite (C) predicted hexagonal packing from X-ray diffraction data. (B) Low-dose (50 e/Å2s) TEM image and (D) line profile showing thread-like features that correlate with the predicted model. Reprinted with permission from [189].

flux by filtering electrons from the gun based on energy, providing an opportunity to control flux by increasing the energy resolution or even by misaligning the monochromator with respect to the emission maximum. For example, Juhl et al. tuned the monochromator to lower the dose rate and thereby achieved low-dose imaging conditions, which was crucial to detect the hexagonally packed carbon nanothread structure that agreed with density functional theory predictions (Fig. 21) [189,190]. A similar strategy enabled HRTEM of highly-oriented polyacetylene [191]. Simply by affording the operator more time, we propose that controlling the electron flux is useful for imaging of radiation-sensitive samples. Furthermore, dose rate may affect radiation damage [192]. While low dose rate techniques may minimize damage in the case of reversible reactions induced by the electron beam, other studies that demonstrate an advantage of high dose rate

[125,126] may be more applicable when slow, irreversible reactions dominate. 6.2. Single shot dynamic TEM In a traditional TEM, only a single electron is traveling down the microscope column at any given time (to put this into perspective, a 200 keV electron travels at approximately 2/3 the speed of light). On the other hand, in dynamic TEM (DTEM), an ultrafast laser pulse illuminates a photocathode source that can photoemit a billion electrons in a single packet. For example, at the Pegasus facility at UCLA, long electron pulses produced by an electron source will enter a linear accelerator whose electromagnetic fields compress the pulse several meters downstream into 10 fs pulses (Fig. 22) [193]. 16

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Fig. 22. At the Pegasus facility at UCLA, an electron pulse (green) passes through a linear accelerator that compresses the pulse to below 10 fs. This duration time is so short that the pulse can “outrun” all atomic motion in molecules, thereby enabling an image to be captured before damage occurs. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

the particles, the low-dose rate method avoids sample degradation. While the single recorded image with a dose of 10 e/Å2 has poor contrast, the accumulated dose of the focus series comes out to 1360 e/Å2. The images can then be aligned to obtain the complex exit wave function, creating an in-line hologram with high contrast and no sample degradation. This method has been used on many beam sensitive systems, such as graphene [196], halide perovskites [197], Ziegler-Natta catalysts [198], and gallium nitride [199]. The innovation of a resonator capable of pulsing an electron beam with picosecond resolution has enabled HRTEM of the beam sensitive Ziegler-Natta catalyst magnesium chloride; this study also suggests that phonons play a role in beam damage, such that by pulsing the electron beam at the time scale of phonon vibrations, the sample can “heal” between pulses [200]. By systematically increasing dose rates, in-line holography can also be used to investigate dynamics. Whether through instrumental optimization or computational analyses, from Eq. (3) it is apparent that enhancing contrast in the TEM is crucial to maximize resolution. Polymeric systems often contain elements other than carbon and hydrogen, creating opportunities through analytical TEM. Furthermore, the inner potential of domains within different polymers often varies, such that phase interference can be used to enhance imaging contrast. These strategies are discussed in the next two Sections.

At time scales this short, all atomic motion is essentially frozen. One of the appeals of ultrafast DTEM is thus the ability to make “molecular movies” [194]. From a polymer microscopy standpoint, another advance of DTEM is the ability to capture atomic images beyond the sample’s damage threshold. In other words, we can “diffract and destroy” or “outrun damage” – electrons will interact with and pass through the sample before damage propagates and affects the structure.

6.3. Low dose rate in-line holography While the “diffract and destroy” method is one solution to the radiation damage problem, another strategy is to “divide and conquer”. In other words, we can deliver electrons in such small quantities at once that the distortion of structure is negligible as long as the excitation is reversible and decays before the next probing electron hits the molecule. This can be repeated over and over again, resulting in a large image series that when averaged, reveal a high resolution image that remains faithful to the undamaged structure [195]. Such an approach uses a Nelsonian illumination scheme that bypasses the conventional collimator setup and is able to deliver electrons in a controlled manner exclusively to the imaged area. This technique has been demonstrated for imaging gold particles with minimal beam-induced damage (Fig. 23) [195]. Whereas the traditional method of imaging with high dose rates shows alterations in

Fig. 23. A comparison of high dose rate images of gold particles to the averaging of low dose rate acquisitions. The low dose rate approach simultaneously avoids beam damage and boosts resolution. Reprinted with permission from [195]. 17

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7. From elemental mapping to mapping local electronic structure

Recent work has highlighted the unique set of challenges for EFTEM when applied to polymeric materials, which are largely due to the inherently low contrast and similar electron densities between domains. An optimized procedure for elemental mapping of polymeric materials was demonstrated with the polymer/fullerene mixture poly(3-hexylthiophene-2,5-diyl)/ [6,6]-phenyl-C61-butyric acid methyl ester (P3HT/PCBM) and emphasizes three factors: focusing at zero-loss with the aid of Fast Fourier Transforms (FFTs), using an objective aperture, and ensuring sufficient signal-to-noise and counts. Additionally, generating a set of images that include a bright field image, a thickness map, and an elemental map at the same region is useful for minimizing misinterpretation of elemental maps [35].

Morphologies of multicomponent polymer systems are inherently challenging to examine using electron microscopy. Soft materials are often amorphous and have similar densities and elastic scattering crosssections. Because of this, conventional TEM, where imaging is based on elastic scattering, is typically limited to samples that are chemically or physically modified with heavy stains or to systems with high contrast (such as organic-inorganic nanocomposites). Spectroscopic imaging exploits contrast mechanisms that conventional TEM cannot. Inelastic scattering of electrons by soft materials generates contrast based on chemical and electronic structure. The development of monochromators that can reduce the energy spread of the electron beam to 0.1 eV (see Section 2.2) as well as spectrometers with detection resolution as low as 0.04 eV have facilitated the development of electron energy loss spectroscopy (EELS) and low-loss EELS. As the electron passes through a sample, it can interact with core or valence shell electrons, causing it to lose energy. Electrons detected by the spectrometer can be recorded as a function of energy-loss; interactions with core electrons result in high energy loss (100 s of eV) and interactions with valence electrons result in low energy loss (1–10 s of eV). Thus, at high energies, the EELS spectrum can be used to quantify elemental composition and at low energies, low-loss EELS can be used to characterize the electronic structure of the sample. Analogously, energy-filtered TEM (EFTEM) can be used for elemental or valence mapping. This can be done in TEM mode, where images are collected at specific energy-losses, or in STEM mode, where an EELS spectrum is collected at each pixel.

7.2. Recent elemental mapping and spectrum imaging examples Elemental mapping by energy-filtered TEM is a convenient technique to determine the structure of polymer systems with significant compositional contrast [5,9,12,203–212]. This technique has been applied to image ions within polymer electrolytes [8,213], to generate elemental maps of solid-electrolyte interfaces [123], and to characterize nanoscale phase separation in mixtures of organic semiconductors [9,11,12,17,19,20,214]. Energy-filtered electron tomography has also revealed the 3D structure of conjugated polymer/fullerene mixtures [215]. Most examples of elemental mapping require very high doses, such that this approach relies on minimal atomic diffusion. As a consequence, EFTEM and elemental mapping have been demonstrated to reveal the mesoscale structure of materials with domains larger than about 5 nm. The distribution of ions within a polymer matrix affects the performance of electrolytes used in lithium batteries [216,217]. EFTEM was used to image bis(trifluoromethane)sulfonimide lithium salt (LiTFSI) within a poly(styrene-block-ethylene oxide) (PS-b-PEO) matrix. LiTFSI ions are selectively imaged using F and Li elemental maps and the PEO block is imaged using the O map (Fig. 24). While LiTFSI does contain O, the low ion concentration ensured the signal was dominated by PEO; using oxygen maps to characterize the PEO regions agreed with both bright-field TEM and small angle X-ray scattering (SAXS) data. EFTEM results demonstrate that the lithium salt preferentially aggregates in the center of the PEO lamellae as the PEO chain length increases, and this localization has been attributed to increases in ionic conductivity [8]. Although elemental maps are quantitative in terms of the total amount of a specific element, variations in thickness can confound compositional information. A simple approach to account for variations in film thickness was demonstrated using a mixture of a conjugated polymer and PCBM. By dividing elemental maps by thickness maps that measure the thickness in units of the mean free path, the local composition of polymer can be extracted. Mean free paths for the individual components were predicted from elemental compositions. As shown in Fig. 25, this approach leads to maps of the volume fraction of polymer by essentially accounting for inelastic scattering at a given edge and the total amount of inelastic scattering at every pixel [9]. EFTEM can also generate images based on differences in plasmon resonances. An advantage lies in that the zero-loss and plasmon peaks typically dominate the scattering cross-section of a material, although generating contrast requires distinct plasmon peaks. Unfortunately, plasmon resonances in polymers are generally broad and overlapping [218]. Nevertheless, due to their high intensity and slight shifts in peak position, plasmon resonances have been used to image various polymer systems [19,20,206,212–214,218–222]. Recent work demonstrates the value of using direct electron detectors for EELS and spectrum imaging. The improved point spread function and pixel density of the direct electron detector provides both improved energy resolution and energy field of view, suggesting that direct electron detectors could facilitate elemental mapping [223]. Indeed, direct detection has allowed for faster acquisition of energy-

7.1. Strategies for optimized EFTEM of polymers EFTEM for elemental mapping may be performed using the two window method or three window method [139,201]. At a minimum, two images must be taken to allow for background subtraction. The two window method involves taking an image at an energy before (preedge) and after (post-edge) the ionization edge. The pre-edge image is subtracted from the post-edge image. This method is insufficient for quantitative analysis [202]. The three window method enables quantitative elemental mapping; two pre-edge images are taken to generate an averaged image for background subtraction from the post-edge image. The resulting map is directly proportional to the total amount of the specific element, such that variations in composition and thickness are convoluted. While the three window method will allow for images with intensities that are proportional to elemental content, taking three separate exposures to create one elemental map can be costly in terms of beam damage to radiation-sensitive polymers. Nevertheless, taking multiple images can be useful to properly account for the background or overlapping resonances [139]. Moreover, the inelastic scattering cross section of electrons used for elemental mapping, core-shell electrons, is relatively low; this necessitates high exposure times for sufficient intensity. The inelastic scattering cross section decreases with increasing energy-loss to the fourth power, such that the inelastic scattering cross section of low-loss energies resonant with valence band transitions is greater than the ionization K-edge by several orders of magnitude [136,137]. Thus, one strategy to obtain high-contrast energy-filtered images for quantitative analysis while minimizing beam damage is to use low-loss energies for imaging. Contrast is generated by differences in valence band electronic structure. Spectroscopic imaging has benefitted from the development of monochromators (Section 2.2). The energy resolution associated with spectral imaging by EELS is limited by the energy spread of the electron source, quantified by the FWHM of the zero-loss peak. Newer monochromators can achieve energy resolutions small enough to resolve the valence band structure of polymeric samples, as low as about 0.1 eV [136]. 18

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Fig. 24. Using elemental maps to quantify domain spacing of a lithium salt distributed in a PS-b-PEO matrix. (a) Chemical structures of block copolymer and lithium salt. (b) Elemental maps of F, Li, and O. Lithium ions localize near the center of PEO lamellae. (c) Lamellae thickness of lithium and fluorine using EFTEM elemental maps. Lamellae thickness of PEO quantified using O EFTEM elemental map (O), bright-field TEM (BF), and SAXS (d x fEO). Scale bar: 50 nm. Reprinted with permission from [8]. Copyright 2009 American Chemical Society.

filtered images of labeled biological specimens, thereby overcoming the problem of drift [224].

at higher energies was collected at higher doses. The resulting composite maps of P3HT and PCBM were generated using a principle component analysis algorithm, and are representative of local variations in the electronic structure of the film (Fig. 27). Further work in STEM mode (e.g., STEM-EELS) could achieve higher energy resolution to resolve subtle features associated with vibronic states or intermolecular coupling. Overall, this work exemplifies the potential of lowloss EELS spectrum imaging to take advantage of contrast arising from the characteristic valence band structure of semiconducting polymers and how advances in instrumentation can drive the development of new techniques. Despite the tremendous potential for analytical TEM, most energyfiltered experiments require very high doses. Low-loss spectrum imaging and EELS does have the potential to image at low doses because of the higher inelastic scattering cross-section at low energy losses, as exemplified in Fig. 27. Nevertheless, developing alternative approaches to enhance contrast is crucial for imaging polymers, as discussed in the next section.

7.3. Monochromated low-loss EELS SI to highlight differences in valence electronic structure Monochromated low-loss EELS spectrum imaging (SI) uses electronic transitions as an additional source of contrast for samples with similar mass densities and elemental composition. In 2000, Varlot et al. used low-loss EELS SI to characterize a triphase polymer composite of PS, PMA, and PB [209]. The PS phase was mapped using the low-loss ππ* transition of the aromatic ring. A slit width of 5 eV, corresponding to the limit of the instrument, was used. Using such a small slit width made obtaining sufficient signal-to-noise without severely damaging the sample challenging. Nevertheless, contrast was significantly enhanced relative to the zero-loss image as seen in Fig. 26. In 2015, Guo et al. used monochromated low-loss EELS to characterize P3HT blended with PCBM [18]. These materials exhibit absorption features at low loss (2–8 eV) that are a result of the conjugated electronic structure. Advances in instrumentation enabled a slit width of 1.5 eV for spectrum imaging, resulting in acquisition of images with good signal-to-noise ratio at 1 eV increments. Experiments were designed to collect data at energies corresponding to P3HT prior to reaching the critical dose for this material, although the “background”

8. Enhancing contrast in polymers with phase plates In the past, one method of increasing contrast in polymeric systems has been to use heavy element stains such as osmium tetroxide, which reacts with unsaturated but non-aromatic carbon-carbon bonds, or Fig. 25. Sulfur map and corresponding composition map of 1:3 poly[(4,4′-bis(2-ethylhexyl)dithieno[3,2-b:2′,3′-d]germole)-2,6diyl-alt-(2,1,3-benzothiadiazole)-4,7-diyl] (PGeBTBT) mixed with [6,6]-phenyl-C 71 -butyric acid methyl ester (PC71BM) annealed at 160 °C for 20 min. The color legend represents the scale for the volume fraction of PGeBTBT. Reprinted with permission from [9].

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Fig. 26. Low-loss EELS SI of a triphase polymer composite incorporating PS, PMA, and PB. (a) Zero-loss image of elastically scattered electrons. (b) PS is mapped using the low-loss image at an energy-loss of 7 eV, corresponding to the π to π* transition of the aromatic ring. Scale bar in both images is 30 nm. Reprinted from [209], with permission from Elsevier.

Fig. 27. Monochromated low-loss EELS spectrum imaging using the valence electronic structure of P3HT/PCBM blends. (a–d) image slices taken at 3, 4, 5, and 6 eV, respectively. (e) P3HT and (f) PCBM phase maps generated using principal component analysis to deconvolute micrographs shown in (a–d). In image (e), P3HT fibers are the bright areas and PCBM-rich domains are bright in image (f). (g) Quantitative composite image generated from phase maps in images (e) and (f). P3HT-rich domains are red and PCBM-rich domains are green. Reprinted with permission from [18]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

8.1. Phase plate contrast enhancement in TEM mode

ruthenium tetroxide, which reacts with most carbon-carbon double bonds, including aromatic rings. Several shortcomings exist in the practice of polymer staining, such as the fact that contrast is qualitative (the way in which the stain interacts with the polymer is often not completely characterized) and the fact that stains can introduce artefacts such as inorganic nanostructures [136]. A recently developed approach, in which the contrast in weak phase objects can be enhanced, relies on phase plates. While the use of phase plates has been useful in a number of organic and biological materials, such as in the aid of 3D reconstructions of proteins [34], the use of phase plates to image polymers remains underexplored.

In organic materials, the phase contrast ratio is often approximately 5%, making them weak phase objects. As described in Section 5.2, interactions between the incident beam and the sample produce a phase shift in the exit wave. Consequently, if the amplitude of the scattered wave is small compared to the unscattered wave, the image contrast will be small. While phase contrast can be enhanced via spherical aberrations under defocused conditions, this leads to oscillating contrast at high spatial frequencies and difficulty in image interpretation. This method also results in a loss of phase information, particularly at 20

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Fig. 28. TEM images of NR-CB with conventional TEM at focus (a), with conventional TEM at underfocus (b), and with a Zernike phase plate (c). Reprinted with permission from [226]. Copyright 2005 American Chemical Society.

high-angle annular dark field detector (collects mostly incoherently scattered electrons and leads to Z-contrast). In any of these cases, the STEM probe prevents interference between scattered and forwardscattered electrons, thereby eliminating the limitations in resolution due to the contrast transfer function of the objective lens. But, this also eliminates contrast enhancement due to defocus and aberrations of the objective lens. In theory, interference between scattered and forwardscattered electrons to create phase contrast is possible in STEM mode. In order to achieve this, however, a very small detector is needed to collect a perfect parallel beam, leading to low signal-to-noise. The combination of spherical aberration correctors and phase plates holds promise in enabling phase-contrast STEM. Ophus et al. demonstrated a new kind of phase-contrast electron microscopy called matched illumination and detector interferometry (MIDI)-STEM. In a MIDI-STEM set up, a ring pattern phase plate with alternating thicknesses of Si3N4 placed at the probe-forming aperture generates a probe with a built-in reference wave; this probe can be rastered across the sample as in traditional STEM imaging. The electrons scattered to high angles can then be collected by a standard ADF detector while a pixelated direct electron detector records an image of the forward-scattered beam at each scanned position. The images generated from the forward-scattered beam are then processed by fitting a virtual detector that matches the geometry of the phase plate. Aberration correction is crucial for MIDI-STEM (and any other phase contrast method) because the q4 dependence of third order spherical aberrations causes phase shifts past π/4, where contrast gets worse or flips sign, in the large scattering vector (q = 1 to 2 1/Å) range. Thus, while the MIDI-STEM phase plates fill in most of the low spatial frequencies, C3 spherical aberration correction is necessary to image with the high spatial frequencies. MIDI-STEM, which combines phase plates, aberration correction, direct electron detectors, and phase reconstruction with an interference pattern, produces almost ideal linear phasecontrast images over a wide range of spatial frequencies, making it an attractive method for imaging soft materials [228].

smaller frequencies. Nevertheless, defocusing to enhance contrast based on differences in the inner potential of two components of a block copolymer has been demonstrated to reveal the microstructure [225]. Here we highlight a few successful examples of polymer microscopy using phase plates. For example, a Zernike phase plate was used to enhance contrast in natural rubber filled with carbon black (NR-CB) (Fig. 28) [226]. The phase plate consists of a thin carbon film with a small hole in the center placed at the back focal plane of the objective lens, such that the phase of the scattered beams is shifted by the carbon film while the unscattered beams pass through the center hole without being phase shifted. The result is that at the image plane, interference between the scattered and unscattered beam enhance phase contrast. Another type of phase plate, the semicircular Hilbert phase plate, also shows promise for revealing structure in polymeric materials. Whereas the Zernike phase plate must be precisely aligned in order for its center hole to be on the optical axis, the semicircular Hilbert phase plate covers one half of the back focal plane such that the central beam passes through an open area close to the edge of the phase plate and only electrons passing through the plate experience an additional phase shift. The lamellar structure of polystyrene-block-polyisoprene (PS-b-PI) was clearly seen in an image taken with a Hilbert phase plate (Fig. 29A) without the use of a stain [226]. More recently, PS-b-PI was also imaged successfully with another type of hole-free phase plate that uses a uniform thin film in the back focal plane (Fig. 29B) [227]. Here, the difference in phase shift arises from the fact that charging occurs only at the beam crossover. Because the primary beam will induce the emission of secondary electrons from the film, the consequential local bias in the film creates a phase shift in the diffracted beams with respect to the primary beam. 8.2. Spherical aberration corrected STEM with phase plates Image acquisition in STEM mode can occur via a bright field detector (collects forward-scattered electrons), an annular dark field detector (collects scattered electrons similarly to dark-field imaging), or a

Fig. 29. (A) Phase contrast image of an unstained lamellar polystyrene-polyisoprene diblock copolymer (PS-b-PI) using a 300 kV microscope equipped with a Hilbert phase plate. Reprinted with permission from [226]. Copyright 2005 American Chemical Society. (B) TEM image of PS-b-PI taken with a hole-free phase plate (inset was taken with a conventional TEM at defocus). Reprinted with permission from [227].

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Fig. 30. (a) 2D image of free standing film of poly(styrenesulfonate-b-methylbutylene) after exposure to air with RH = 98% for 24 h. (b) 3D electron microtomography image obtained from 53 tilt series images of sample in (a). Reprinted with permission from [66].

Fig. 31. (a) Bright-field cryo TEM (2D projection) of a frozen-hydrated, as-cast 100 nm Nafion membrane. (b) Cryo TEM 3D reconstruction showing two perpendicular slices through the tomogram. Isosurface rendering is used to mark the spatial distribution of the central region of the dark hydrophilic phase in yellow. Reprinted with permission from [229]. Copyright 2014 American Chemical Society. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

9. Three-dimensional imaging opportunities using direct electron detectors Understanding the three-dimensional (3D) structure of a system can offer invaluable insights on the morphology that is otherwise difficult to interpret through two-dimensional (2D) projections. Electron tomography is often performed by taking many images at different sample tilt angles to elucidate the 3D morphology of various polymer systems. For example, the orientation of domains in polymer electrolyte membranes made from poly(styrenesulfonate-b-methylbutylene) copolymers with different mol% of sulfonated polystyrene moieties in contact with humid air was imaged via electron tomography (Fig. 30) [66]. Fiftythree TEM images were obtained using tilt angles ranging from -52° to +52° and then aligned with the aid of gold nanoparticles that served as fiducial markers. While the conventional 2D image can only suggest the presence of perpendicular cylinders, the 3D structure demonstrates that the vertically oriented cylinders span the entire thickness of the film. Electron tomography has also been used to study the hydrated form of Nafion membranes, revealing an interconnected channel-type network with a domain spacing of about 5 nm (Fig. 31) [229]. Recent work has demonstrated the need for 3D tomographic reconstructions to properly characterize closed voids and surface area (surface roughness) in water filtration membranes [230]. Cryo-tomography of P3HT assemblies in vitrified organic solvents has enabled even higher resolution, revealing a 3D lamellar structure of the 1.7 nm stacking of conjugated backbones showing increased order in the bulk of nanowires [231]. Like in the previous examples, the 3D reconstruction was obtained with a TEM tilt series. This approach has been successfully applied to many polymeric systems, and progress, applications, challenges

and opportunities have been summarized in various reviews [232–236]. We thus focus our discussion on techniques for 3D reconstructions beyond the acquisition of tilt series. 9.1. 3D structure from single particle reconstructions While electron tomography of polymeric materials has been achieved through tilt series of TEM images, the repetitive imaging at one location deposits high doses to the imaged area. As such, singleparticle cryo-electron microscopy, which does not require tilting of the sample and instead uses 2D images of individual particles, is a promising method for beam-sensitive materials. Single-particle cryo-electron microscopy has proved to be useful in determining the macromolecular structures of many biological materials, such as proteins and viruses. The workflow for this process generally involves flash-freezing the biological sample, collecting 2D projections via cryo-electron microscopy, aligning and averaging the images, and then finally constructing the 3D model (Fig. 32) [237]. Unlike in tomography, the collection of 2D images does not involve a tilt series at a single location on the sample. Instead, 2D images of individual particles dispersed throughout the sample grid are taken. Although each 2D projection alone is too noisy to discern atomic detail, the signal can be improved by averaging over thousands of particles using stages that can hold multiple samples and acquire images 22

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automatically. With the development of direct electron detectors, 2D projections can be collected with much greater sensitivity. Moreover, because particles are frozen in random orientations throughout the sample grid, the 2D projections of such a large number of individual particles provides the distribution of 2D views necessary to reconstruct the 3D structure. Once the 2D images are aligned and averaged, software can be used to construct a 3D map that is iteratively refined. While this method has been successful for biological samples where each particle is identical, the conformational freedom of most polymers and polymer nanostructures makes averaging over multiple regions and samples essentially impossible. Nevertheless, a combination of crystallographic and single-particle methods has been successful on a highly ordered, synthetic polymer: a peptoid [238,239]. As seen in Fig. 33, these images revealed the V-shaped motif of a peptoid nanosheet with atomic resolution. We predict that single particle reconstructions may reveal 3D reconstruction of other highly ordered polymeric materials that are challenging to achieve through other means, such as for the unit cell of nanocrystalline polymers. 9.2. Exit wave reconstructions Exit wave reconstruction is a well-established technique for HRTEM in which a series of images taken at different defocus can be used to improve image interpretation and increase resolution. Phase information is lost when recording an image or scattering intensities, although in principle the information is present due to interference; exit wave reconstructions take advantage of phase interference from different values of defocus to recover phase information of the exit wave. More recently, exit wave reconstructions have been used to extract 3D information at the atomic level. This is possible because although images acquired in HRTEM are 2D projections, the reconstructed exit wave contains 3D information due to the sample thickness, relative position of atomic layers, and multiple scattering events that occur as the electrons pass through the sample. In fact, simulations have demonstrated that the use of direct detectors allow for accurate exit wave reconstructions that can lead to extraction of 3D information [240]. This method is especially attractive for polymers because it is compatible with low-dose imaging, such that beam damage is minimized without compromising atomic resolution. 9.3. MicroED tomography Electron crystallography has been a powerful tool for studying atomic structure for many years, but it has traditionally been restricted to 2D patterns from crystals and has only modest resolution (4–10 Å). In 2013, electron diffraction was used to determine the structure of a protein in 3D using a method called MicroED [241]. The overall idea of this technique is that if many diffraction patterns are taken from a single crystal, a large enough region of reciprocal space would be represented such that the crystal data set could be properly indexed. This is analogous to X-ray crystallography, but has two important distinctions. First, electrons interact more strongly with matter despite less energy onto the sample, thereby allowing electron diffraction to extract high-resolution data from very small crystals. Second, electrons have a much smaller wavelength and are therefore able to probe smaller length scales. The first iteration of MicroED involved a series of still diffraction patterns as the crystal was rotated in discrete angles between exposures, with the entire process conducted under cryogenic conditions and a dose rate on the order of 0.01 e/Å2s. This resulted in a 3D lysozyme structure with 2.9 Å resolution. Following this proof of concept, a more sophisticated version of MicroED was developed. This involved continuous-rotation of the crystal and was made possible by the high frame rates achievable by direct electron detectors [242]. Continuous-rotation MicroED not only resulted in improved data quality, but allowed data processing to be done with standard X-ray crystallographic software.

Fig. 32. Workflow of single-particle cryo-EM of a protein sample. Reprinted with permission from [237].

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Fig. 33. A combination of crystallographic and single-particle methods, originally developed for cryo-electron microscopy of biological macromolecules, was successfully implemented on self-assembled nanosheets of a peptoid polymer. The atomic length scales achieved with this technique represents a new level of resolution for polymer microscopy. Reprinted with permission from [238]. Copyright 2019 American Chemical Society.

Fig. 34. Workflow for microED (top) and example structures determined by this method (a–d). Structures are shown with their resolution in parentheses, the full model on the left, and the representative region of the model and density map on the right. Reprinted from [243], with permission from Elsevier.

damage under the electron beam during imaging, it is difficult to distinguish between morphological changes caused by the beam versus changes caused by the applied external stimuli. A better understanding of beam damage combined with direct electron detectors that will allow for low-dose imaging makes it an ideal time to revisit in situ TEM experiments for polymers.

The workflow for MicroED involves identification of a suitable crystal, continuous collection of diffraction data as the sample is rotated in the beam, and then indexing by standard X-ray crystallography programs (Fig. 34) [243]. So far, MicroED has solved the crystal structure for a variety of biological samples, such as lysozyme, trypsin, and thermolysin, among others. Nevertheless, to our knowledge, it has not yet been applied to synthetic polymers. Most of the samples discussed in this Section have been captured in their frozen state. Nevertheless, in order to study solutions in their true, liquid form or to study solid samples under different external stimuli, in situ TEM methods, which are discussed in the next Section, are needed.

10.1. Heating/cooling stages Perhaps the most common of in situ TEM studies involves temperature changes during experimentation. The earliest use of temperature in situ TEM comes from Easterling in 1970 to study iron and iron-nickel composites at the atomic level [245]. The ability to look at changes in the microscopy images and their Fourier transforms during real-time heating events allows for monitoring growth processes at the atomic scale to understand growth and growth mechanisms [246–249]. Recent developments in precise measurement control of TEM holders have further pushed the boundaries of temperature-dependent morphological resolution [248,250]; sample holders with very small thermal mass allow for increased temperature control. A schematic of in situ TEM heating can be seen in Fig. 35, along with electron diffraction patterns of the thermolysis of ammonium tetrathiomolybdate [246]. As the sample is heated from room temperature

10. In situ TEM using new sample holders Demonstrations of in situ electron microscopy go back to work by Pashley in 1956, who examined material epitaxy in thin films [244]. Combining TEM and external pressures, such as temperature changes, mechanical stresses, electrical bias, and chemical exposure, allows microscopists to study instantaneous material changes at the atomic level. Using various techniques to study samples while they undergo these phenomena is critical to understanding certain dynamics of these materials. One of the major challenges of in situ TEM of polymers is again electron beam damage. Because polymers are susceptible to 24

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Fig. 35. An in-situ heating TEM experiment is shown. (a) A schematic of the heating stage is presented, with the red “solid precursor” being (NH4)2MoS4. (b–e) Selected-area electron diffraction of the thermolysis of (NH4)2MoS4 is shown at various temperatures (25 °C, 400 °C, 780 °C and 900 °C), demonstrating the crystallization of MoS2 up to 800 °C followed by decomposition of MoS2 into metallic Mo at about 800 °C. This work is licensed under the Creative Commons Attribution 4.0 International License. To view a copy of this license, visit http://creativecommons.org/licenses/by/4.0/ or send a letter to Creative Commons, PO Box 1866, Mountain View, CA 94042, USA. [246]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

to 900 °C, crystallization is observed, followed by metallic decomposition. By observing these processes in situ, not only can the existence of crystallization at a specific temperature be confirmed, but the grain size of these crystals can also be determined. Phase changes can be directly observed in real-time as well, further increasing the flexibility of this technique. In situ heating in the TEM has been applied to many inorganic systems, such as for observing the growth mechanisms of 2D molybdenum disulfide flakes [246], the rotation of grains in nanocrystalline platinum [247], the martensitic transformation of nickel titanium shape memory alloys [249], and grain growth in nanocrystalline copper thin films [251]. To our knowledge, its application to organic materials has not been widely published. Nevertheless, in situ heating in the TEM has offered valuable contributions to organic-inorganic systems, such as for monitoring the degradation of organometallic halide perovskite solar cells under heating [248] and observing the loss of a capping polymer in the heating of polymer-capped platinum nanocrystals [252]. Even minimal damage below the critical dose as measured via electron diffraction could cause chemical changes that alter the dynamics and phase behavior of organic materials, thereby limiting in situ heating TEM studies. As a consequence, a series of in situ heating images that are taken with a total dose below which damage to the chemical structure is significant is needed; we propose that low-loss EELS may be useful to define this acceptable dose by measuring damage to the valence electronic structure. On the other hand, materials can also be studied in the TEM while undergoing cooling (in situ cooling). Using TEM cooling stages, various phase and magnetic transformations can be accessed that cannot generally be observed with isothermal TEM stages [253,254]. While the effects of heating are often detrimental to polymeric materials, cooling slows down transient processes and has enabled imaging of dynamic processes, for example, in the cooling of nanolaminates [255]. We predict that the slowing of reactions due to cooling could be applicable to polymer electron microscopy as well. Unlike the aforementioned cryo-TEM, where motion is essentially frozen, variable cooling could enable imaging of dynamic processes at select temperatures, such as polymer crystallization upon quenching. Such a study could also be useful for visualizing processes associated with Tg for polymers with Tg below room temperature. Care must be taken to use low dose rates such that local heating from the electron beam does not alter the desired imaging condition temperature.

developments in TEM sample holders have focused on the miniaturization of mechanical testing instrumentation that is compatible with sample stage space limitations while maintaining the ability to accurately control and measure applied stress, such as in situ nanoindenters [256,257] and on-chip microelectromechanical systems (MEMS) [258,259]. Since its inception a few decades ago, in situ TEM mechanical testing has made important advances in materials science, by revealing the nanoscale origins for deformation [260], tearing [261], and compression [262] of molybdenum disulfide, the compression of metal-organic framework microcrystals [263], tensile deformation of carbon nanofibers [264], and the nucleation and propagation of dislocations in aluminum films [265]. In situ mechanical studies of polymers would be valuable because there is a wide range of micromechanical processes that occurs in polymers under load, such as nanometer scale changes to individual macromolecular segments and micrometer scale plastic yielding (e.g., crazing and shear bands). Unfortunately, the main limitation, in addition to aforementioned interpretation difficulties that arise as a consequence of beam damage, is the challenge of conducting mechanical tests on polymer films that are thin enough to be electron transparent. One solution has been high-voltage electron microscopy (HVTEM), which allows for thicker specimens. For example, in situ straining of poly(styrene-co-acrylonitrile) (PSAN) blended with poly (methyl methacrylate) (PMMA)/acrylate-rubber/PMMA core-shell particles in a HVTEM (1000 keV) at room temperature revealed simultaneous crazing and shear yielding of the PSAN and fibrillation of the rubber particles with drawing of material from the PMMA cores (Fig. 36) [266]. For this study, microtomed polymer sections were glued between two metallic films to create a “sample/metallic-film sandwich” that could be mounted into a tensile holder. Continued advances of in situ mechanical testing systems will likely be necessary in order to overcome current limitations in the field. For example, current MEMS chips are not amenable for soft materials, as specimens are often attached to MEMS devices inside a FIB. The advent of the cryo-FIB, discussed in sections 3.1 and 3.3, could help mitigate this issue. Furthermore, as increasingly creative methods for measuring the mechanical properties of polymer thin-films are explored, such as tensile testing thin-films on a liquid surface [267] or an elastomer [268], in situ mechanical testing of thin films in the TEM could be more widely adopted and used to explore phenomena such as shear-induced crystallization [269,270].

10.2. Mechanical deformations

10.3. Liquid cell TEM

Although mechanical properties of materials are directly linked to their structure, ex situ mechanical testing complemented by imaging of stress-induced microstructural evolution struggles to connect nanoscale phenomena to macroscopic responses to applied stress. Recent

Recent development in sample holders that can encase a droplet of liquid have enabled liquid cell TEM [271–274]. The use of 100 nm Si3N4 films to create liquid sample holders was demonstrated for the TEM in 2003 by Williamson while at the same time Thiberge developed 25

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Fig. 36. (A) TEM image of stained core-shell particles in PSAN matrix. (B) Deformation of core-shell particles at room temperature. Direction of deformation indicated by arrow. Reprinted with permission from [266].

a similar cell for the scanning electron microscope (SEM) [273,274]. Because the window material adds background scattering, Si3N4 films have been further thinned to 25 nm [275]. Liquid cells have enabled studies of nanoparticle and nanocrystal synthesis growth, as well as growth of lithium dendrites at the interface between metal electrodes and organic electrolytes [276–281]. Controlling the sample and window thickness to minimize background scattering, as well as improvements in detectors, continue to push the limits of liquid cell TEM [282]. Two other limitations of liquid TEM, electrical charging and radiolysis of the liquid under the electron beam, have also recently been addressed by the development of encapsulating graphene sheets that alleviate charging and using deuterated water to prolong the sample lifetime [283]. Liquid cell TEM has been demonstrated on polymer solutions. Fig. 37A–C shows a small amount of liquid in between two graphene sheets where solubilized polymer chains are apparent [284]. The authors could discern changes in polymer conformations as a function of time at the timescale of seconds. Given that polymer segmental dynamics are expected to be much faster, the authors speculated that polystyrene sulfonate chains were absorbed onto graphene, thereby slowing down motion. Liquid cell TEM has also been conducted on a solution of amphiphilic block copolymer micelles, revealing the evolution of micelle fusion and growth in real time with nanometer spatial resolution (Fig. 37D) [285]. These studies clearly demonstrate the possibility of examining dynamical processes of polymers in solutions. 10.4. Electrical bias Fig. 37. Liquid cell TEM. (a) Schematic of liquid cell TEM, showing a polymer solution between two electron transparent substrates (graphene sheets). (b) Low magnification TEM image that shows the polymer solution channel in between graphene sheets. (c) TEM image of polymer molecules in solution. Reprinted, by permission, from [284]. Copyright 2017 by John Wiley & Sons, Inc. (D) Direct observation of growth and evolution processes in block copolymer micelles. Reprinted with permission from [285]. Copyright 2017 American Chemical Society.

The application of electric fields is possible either in the solid state or in solution through microfabricated chips and sample holders. Electrical bias experiments in the TEM have elucidated field-induced phenomena such as oxygen vacancy migration in cerium oxides [286], reversible formation of lamellar domains in a piezoelectric [287], structural damage of multiwalled nanotubes [288], and nanostructure changes in perovskite solar cells as a function of a current-voltage stimulus [289]. In the case of polymers, resolving phase transformations, structural rearrangement, or time-resolved processes due to an applied field requires enough radiation hardness to acquire multiple images. As a consequence, no direct application to polymers has been demonstrated. The development of liquid cell TEM allowed for the development of electrochemical sample holders by adding electrodes to Si3N4 windows [258,274,275,277,280,290–293]. In situ TEM electrochemistry allowed for studies of dendrite growth, intercalation and interfacial changes in electrochemical cells such as batteries [277,280,291,293]. The electrochemical polymerization of poly(3,4-ethylenedioxythiophene) (PEDOT) has also been visualized, as shown in Fig. 38 [294]. TEM

studies reveal that particulates are formed in solution as the deposition proceeds, followed by deposition of these polymer particles. This refines previous descriptions of electrochemical polymerization, which assumed that only monomers approach the surface prior to chain growth. 11. Summary TEM has transformed our understanding of the structure and morphology of polymeric materials. Throughout the past few decades, 26

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Fig. 38. Direct imaging of the in situ electrochemical deposition of PEDOT from an aqueous solution of EDOT monomer using TEM with an electrochemical liquid flow cell. Captured frames from live video of growth of PEDOT clusters during electrochemical deposition (a–f) and projected areas of individual PEDOT clusters (g).Reprinted with permission from [294]. Copyright 2015 American Chemical Society.

Acknowledgements

polymer microscopy has evolved to include high-resolution imaging, spectrum imaging, tomography, and in-situ characterization. Nevertheless, growth in the field has always been hampered by the inherent lack of contrast and beam sensitivity of polymers. The accomplishments of polymer microscopy despite these constraints have been impressive, although new advancements in instrumentation such as specimen preparation tools, monochromators, aberration correctors, phase plates, direct electron detectors, and in situ sample holders can push imaging limits even further. For example, the development of sample preparation techniques based on the FIB, new microtome technology, and cryo-preparation methods will enable new imaging experiments. The application of monochromators to polymer microscopy provide a means to control beam flux, and a way to probe beam damage through monochromated EELS, which in turn paves the way for optimized EFTEM and spectrum imaging using differences in valence electronic structure. Various experiments have been made possible by new sample holders, such as liquid TEM and in situ electrochemistry. By highlighting the limitations and potential of HRTEM of polymers, we establish that a combination of protocols to minimize the effects of beam damage with aberration correctors and direct electron detectors will push the resolution limits of polymer microscopy despite beam sensitivity. We propose that electron microscopy is poised to leap forward and continue to transform our understanding of structure and structure-property relationships in polymers. Indeed, various techniques highlighted here are currently underexplored in polymer microscopy, such as 3D reconstructions beyond tilt series and 4D STEM.

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Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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