grain boundary engineering thesis

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Grain boundary engineering in Ni-based alloys: processing, mechanisms and effect on mechanical properties

The property of polycrstalline nickel alloys are crucially dependent upon grain boundaries, as they are the strength limiting sites at high temperatures. Engineering of grain boundary structure and phases can be effective means of pursuing optimum properties of a component. Grain boundary morphologies are known to have critical impact on time dependent deformation, such as creep and fatigue. However, the reason for such an improvement is only understood qualitatively, the mechanisms of its...

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• Published: 02 January 2020

Engineering grain boundaries at the 2D limit for the hydrogen evolution reaction

• Yongmin He 1 , 2   na1 ,
• Pengyi Tang   ORCID: orcid.org/0000-0002-2306-095X 3 , 4   na1 ,
• Zhili Hu 5 , 6   na1 ,
• Qiyuan He   ORCID: orcid.org/0000-0001-9083-8285 1   na1 ,
• Chao Zhu 1 ,
• Luqing Wang 6 ,
• Qingsheng Zeng 1 ,
• Prafful Golani 1 ,
• Guanhui Gao 7 ,
• Zhiqi Huang 1 ,
• Caitian Gao 8 ,
• Juan Xia 9 ,
• Xingli Wang 10 ,
• Xuewen Wang 11 ,
• Chao Zhu   ORCID: orcid.org/0000-0002-1589-855X 1 ,
• Quentin M. Ramasse 12 , 13 ,
• Ao Zhang   ORCID: orcid.org/0000-0002-9427-9641 1 ,
• Boxing An 14 ,
• Yongzhe Zhang 14 ,
• Sara Martí-Sánchez 3 ,
• Joan Ramon Morante   ORCID: orcid.org/0000-0002-4981-4633 4 ,
• Liang Wang   ORCID: orcid.org/0000-0002-3771-4627 15 ,
• Beng Kang Tay   ORCID: orcid.org/0000-0002-3776-3648 10 ,
• Boris I. Yakobson 6 ,
• Achim Trampert 7 ,
• Hua Zhang   ORCID: orcid.org/0000-0001-9518-740X 1 , 16 ,
• Minghong Wu   ORCID: orcid.org/0000-0002-9776-671X 15 ,
• Qi Jie Wang   ORCID: orcid.org/0000-0002-9910-1455 2 , 17 ,
• Jordi Arbiol   ORCID: orcid.org/0000-0002-0695-1726 3 , 18 &
• Zheng Liu   ORCID: orcid.org/0000-0002-8825-7198 1 , 8 , 17 , 19  

Nature Communications volume  11 , Article number:  57 ( 2020 ) Cite this article

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• Catalyst synthesis
• Electrocatalysis
• Two-dimensional materials

Atom-thin transition metal dichalcogenides (TMDs) have emerged as fascinating materials and key structures for electrocatalysis. So far, their edges, dopant heteroatoms and defects have been intensively explored as active sites for the hydrogen evolution reaction (HER) to split water. However, grain boundaries (GBs), a key type of defects in TMDs, have been overlooked due to their low density and large structural variations. Here, we demonstrate the synthesis of wafer-size atom-thin TMD films with an ultra-high-density of GBs, up to ~10 12  cm −2 . We propose a climb and drive 0D/2D interaction to explain the underlying growth mechanism. The electrocatalytic activity of the nanograin film is comprehensively examined by micro-electrochemical measurements, showing an excellent hydrogen-evolution performance (onset potential: −25 mV and Tafel slope: 54 mV dec −1 ), thus indicating an intrinsically high activation of the TMD GBs.

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Introduction

Grain boundaries (GBs) are commonly found in atom-thin or so called two-dimensional (2D) polycrystalline materials 1 , 2 , 3 , where they can be described as line defects. They play a key role in shaping the properties and performance of 2D materials in applications as varied as mechanical strengthening 4 , photovoltaics 5 , 6 , electronics 7 , 8 , 9 , and catalysis 10 , 11 . Engineering the structure and/or the density of GBs in 2D materials could thus become a promising way to tailor their performance. One particular class of 2D materials, transitional metal dichalcogenides (TMDs) have attracted a great deal of attention for its possible uses in electrocatalytic reactions 12 , 13 , including the hydrogen evolution reaction (HER) 14 , 15 . Due to the low cost, earth abundance and good stability of a wide range of TMDs, substantial work has thus been undertaken to improve their electrocatalytic activity 12 , 13 by, e.g., exposing their edges 16 , 17 , 18 , 19 , doping with heteroatoms 20 , and/or creating and straining structural defects 11 , 15 , 19 , 21 , 22 . In contrast, less attention has been paid to the role of GBs due to their typically low number density and large structural variations, even though GBs have been predicted to be highly electrocatalytically active 23 . The poor control over the density and structure of GBs stems from the fast gaseous kinetic processes and the multiplicity of chemical phases involved during the growth of TMDs 2 , 24 , 25 . To date, the most common methods employed to synthesize atomically thin polycrystalline TMDs include chemical vapor deposition (CVD) 2 , 3 , 26 , 27 , 28 , 29 , 30 , 31 , 32 , 33 , 34 , 35 , 36 , physical vapor deposition (PVD) 37 , 38 , 39 , and metal organic chemical vapor deposition (MOCVD) 40 , 41 . Films grown using these techniques usually exhibit grain sizes ranging from hundreds of nanometers to several millimeters, resulting in a low GB density, as summarized in Fig.  1a (See the statistical method in Supplementary Fig.  1 ).

a Overview of the grain size and density in TMD materials obtained by various fabrication methods, such as CVD 2 , 3 , 26 , 27 , 28 , 29 , 30 , 31 , 32 , 33 , 34 , 35 , 36 , MOCVD 40 , 41 , PVD 37 , 38 , 39 , top–down syntheses of sulfurization/selenization of metal and metal oxide thin films 73 , 74 , 75 , 76 , 77 , and thermal decomposition of thiosalts thin films 78 . Although small grains (~20 nm in diameter) were observed in TMDs synthesized by the top–down methods of sulfurizing Mo-based thin films 73 , 74 , or thermal-decomposition of Mo-thiosalts thin films 78 , the grain density in the film is not very high owing to their interlayer diffusion synthesis mechanism. b Schematic of the wafer-scale growth of TMD nanograin films. Ultra-high-density Au quantum dots (QDs) were used to grow the MoS 2 (as well as WS 2 atom-thin films: see  Supplementary Information ). c TMD nanograin film from wafer scale to nanoscale, including photograph (i), optical image (ii), SEM image (iii), and HRTEM image (iv). d False-colored SEM image of the Au QDs on a SiO 2 /Si substrate, showing an average diameter of 4.8 nm. e Statistical distributions of Au QDs obtained from Au films with different deposition times. The evaporation rate is ~0.1 Å s −1 in our experiments. f Raman spectra acquired from the MoS 2 film. The difference ( ∆ ~20 cm −1 ) between the out-of-plane (A 1g ) and in-plane (E 1 2g ) Raman modes indicates that the MoS 2 film is 1–3 layers (1–3 L) 2 , 44 . Compared to CVD-grown and exfoliated samples, the reduced intensity of the in-plane mode indicates a highly polycrystalline structure for this MoS 2 film.

Here, we fabricate wafer-size atomically thin TMD films with sub-10 nm grains by means of Au-quantum-dots (QDs)-assisted vapor-phase growth, demonstrating an ultra-high-density of GBs (up to ~10 12  cm −2 ). The quality of the films was examined by high-resolution transmission electron microscopy (HRTEM), aberration-corrected scanning transmission electron microscopy (STEM), scanning electron microscopy (SEM), and Raman spectroscopy. Experimental evidence as well as phase-field simulations demonstrate that the Au QDs regulate the formation of the TMD grains. We then investigated the catalytic activity of these nanograin films by a four-electrode micro-electrochemical cell for hydrogen evolution. An excellent performance (−25 mV and 54 mV dec −1 for the onset potential and the Tafel slope) was obtained for our MoS 2 nanograin films, indicating a good intrinsic activation of the GB-rich 2D basal plane.

Controlled growth of the TMD nanograin film

The main obstacles that hinder the vapor growth of TMDs at the 2D limit to grain sizes typically <10 nm consist of the precise control of the nucleation sites and the growing rate of the grains. These can be addressed by using a high-density of Au QD seeds and a low mass flow rate of the vapor sources. Figure  1b illustrates the wafer-scale growth process (see also a full description of the growth setup in Supplementary Fig.  2 ). A wafer-scale Au QD layer was fabricated on 2-inch sapphire or SiO 2 /Si substrates and used to subsequently grow the atom-thin MoS 2 films. Figure  1c shows the geometries of the as-grown wafer-size film from the centimeter to the nanometer scale. The SEM image (Fig.  1d ) shows the as-prepared Au QDs with an ultra-high density up to ~2 × 10 12  cm −2 and average diameter down to ~4.8 nm, which was achieved by heating a thin Au film on sapphire or SiO 2 /Si substrates at a high temperature. We attribute the formation process of these Au QDs to a solid-state dewetting behavior at the interface between the SiO 2 (or Al 2 O 3 ) and Au 42 , 43 (Supplementary Fig.  3 ). The initial deposition time of the Au films prior to heating determines the density and the size of the final Au QD structures: a shorter deposition time results generally in a smaller size and a higher density of Au QDs, as shown in Fig.  1e and Supplementary Fig.  4 . The MoS 2 film is then grown using a vapor-phase growth technique (see further experimental details in Method section), after which Au QDs can be removed from as-grown films, by using KI/I 2 etchant at room temperature. The successful removal of the Au QDs after the etching treatment was assessed by X-ray photoelectron spectroscopy and STEM imaging in Supplementary Figs.  5 – 7 , showing only minimal amounts of residual Au present. The resulting MoS 2 films were examined by Raman spectroscopy, as shown in Fig.  1f . The Raman spectroscopy measurements indicate that the MoS 2 film is 1–3 layers thick (1–3 L). A weak in-plane mode (E 2g 1 at 385 cm −1 ) was observed in our 1–3 L film, with an intensity ratio with the out-of-plane mode A 1g (at 405 cm −1 ), \(I_{E_{2g}^1}/I_{A_{1g}}\) (~0.28), significantly smaller than the ratios typically measured on the CVD-grown (~1.89) and exfoliated (~1.57) MoS 2 films in our experiment, suggesting a highly polycrystalline structure in this film (see the details in Supplementary Table  1 and Supplementary Note  1 ).

Au QDs regulate the formation of TMD grains in two ways. The first way is that Au QDs can facilitate the TMD nucleation at the initial stage of the TMD growth. The melting temperature of Au will dramatically decrease as its size decreases to several nanometers, as shown in Supplementary Fig.  8 . As a result, at the growth temperature of TMDs (650–800 °C), Au QDs tend to be in the liquid phase due to their small size. Such liquid Au droplets can facilitate the formation of the TMD nucleation sites during the growth process. The second way is that the Au QDs can confine the size of the TMD domains below 10 nm. Our deposition method can produce well-dispersed Au QDs with a very high number density, with average spacing between QDs down to a few nanometers. As a result, the formation of the TMD grains will be constrained at this length scale.

Atomic structure of the TMD nanograin films

We investigated the atomic structure of the TMD nanograin films using transmission electron microscopy (TEM). Figure  2a shows an overview TEM image of a uniform and continuous MoS 2 film suspended on a Cu-supported lacey carbon TEM grid. This region is representative of the tens of samples examined, and the MoS 2 film comprises polycrystalline patches with regions of 1–3 L MoS 2 (consistent with the measured Raman frequencies 44 ). To evaluate the grain distribution, we randomly chose six regions of the film, labeled 1–6 on Fig.  2a , for further HRTEM imaging. The original HRTEM images from these six regions, corresponding Fourier transforms (FFTs) and frequency filtered images (IFFTs) 45 are shown in the top (a1–a6), middle (b1–b6), and bottom panels (c1–c6) of Fig.  2a , respectively. To aid the visualization of the different grains, black-dashed lines are added along the edges of the grains in the bottom panel (c1–c6). From these images, we can identify 8–10 distinct MoS 2 grains in a ~700 nm 2 region, suggesting an ultra-high grain density (up to ~10 12  cm −2 in this representative patch, consistent with the initial estimate derived from the Au QD number density) in our sample. Accordingly, the average diameter of the grains is <10 nm, with some observed grains smaller than 5 nm in diameter. To our knowledge, this is the smallest grain size obtained to date in materials at the 2D limit 17 , 46 , 47 .

a TEM image of the 1–3 L MoS 2 nanograin film, showing a uniform and continuous atom-thin layer with Au nanoparticles (dark spots). Top panel (a1–a6): HRTEM images from regions 1–6. Middle panel (b1–b6): Fourier transforms (FFTs) of regions 1–6 from the corresponding TEM image. Bottom panel (c1–c6): false-colored frequency filtered (IFFTs) images of the same regions. Grains with different orientations give rise to distinct sets of rotated spot patterns in the FFTs, as indicated by colored circles in the middle panel (b1–b6); in turn, the localization of the grains can be visualized in real space by color coding the corresponding contribution to the IFFT of a given set of diffraction spots, and overlaying the IFFTs into a composite colored image (in bilayer regions, the colors are combined). In order to better visualize the different grains, the grains are highlighted in the IFFT images with black-dashed lines. b HAADF STEM investigation of MoS 2 grains. Left: HAADF STEM image showing the GBs between three MoS 2 grains. Inset on the left figure: Fourier transform of the image showing three distinct sets of MoS 2 diffraction patterns with rotation angles of 11.6° and 25.6°. Middle: composite color-coded IFFT image. Right: dilatation map of b after applying GPA routine to the monolayer MoS 2 . c Higher magnification atomic resolution HAADF STEM image of the brown-dashed line marked region in b , showing the detailed atomic structure of the GB, which is composed of 5|7 and 4|8 rings. Mo atoms are marked with indigo dotted circles and S atoms marked with yellow dotted circles. d Grain size dependence on Au nanoparticle size (average diameter) in the MoS 2 nanograin film. The error bars of Au nanoparticles are extracted from the full-width at half maximum in Fig.  1e .

We then employed atomic-resolution high-angle annular dark-field (HAADF) imaging in the STEM to examine the atomic structure of the GBs on the as-grown nanograin films, representative examples of which are shown in Fig.  2b, c and Supplementary Fig.  9 . The inspection of dozens of boundary locations systematically found that nanograins in our films are stitched by GBs. The atomic structure of the GBs varied significantly (great care was taken to exclude any structure modification arising from beam damage, see Experimental section for details). Structures comprising a combination of 5|7 and 6|8 rings were however dominant in our nanograin films and are illustrated in Fig.  2c . A small number of 8|4|4 structures were found, but in contrast, 12|4 rings were never observed in “pristine” GBs (i.e., before any substantial exposure to the electron beam). Their atomic structures in MoS 2 are schematically illustrated in Supplementary Fig.  10 .

Although the nature of atomic-resolution microscopy precludes large-scale statistical studies, we investigated as many different regions in a MoS 2 nanograin film as practically possible. Six further regions over a large range were carefully examined with atomically resolved HAADF–STEM imaging (Supplementary Fig.  11 ), and 3 to 5 GBs can be observed in a 400 nm 2 area, confirming the ultra-high-density of GBs observed in STEM (consistent with the ~10 12  cm −2 suggested by the Au QD density). Moreover, tens of samples with different sizes of Au QDs were grown: the grain size measured in the resulting MoS 2 films exhibits a clear linear relationship with the QD size, as shown in Fig.  2d . This provides a strong proof that the TMD grain size can be carefully controlled by the Au QD substrate. Remarkably, our method can be also extended to other TMD materials such as WS 2 (Supplementary Fig.  12 summarizes similar results to those described above, using WS 2 instead of MoS 2 ), making it as a universal approach for the wafer-size synthesis of atom-thin films with sub-10 nm grains.

Growth mechanism of the TMD nanograin film

A climb and drive zero-dimensional (0D)/2D interaction is proposed to explain the Au QDs-assisted-growth mechanism of the TMD nanograins at the atomic limit, which is supported by real-time phase-field simulations and experimental evidence. The phase-field simulation of the climb process shows that, once the Au QDs encounter a MoS 2 edge at the growth front, they migrate swiftly from the SiO 2 surface onto the MoS 2 surface (Fig.  3a ). This behavior is mainly attributed to a smaller wetting angle of the Au QDs droplet (Au is in the liquid phase at the growth temperature) on MoS 2 than that on the SiO 2 substrate 48 (Supplementary Fig.  13 ). This phenomenon is also experimentally confirmed by cross-sectional TEM, as shown in Supplementary Fig.  14 , where nearly all the Au QDs sit on the surface of the MoS 2 . Subsequently, the phase-field simulation suggests the formation of a second MoS 2 layer will drive the Au QD droplets along its growth direction (Fig.  3b ). A number of the Au QDs may thus coalesce and form bigger droplets during the driving process. This is supported by our experimental evidence: the size of the Au QDs on the MoS 2 is usually larger than those on SiO 2 , and some of them sit at the edge of the MoS 2 sheet, as illustrated in Fig.  3c and Supplementary Fig.  15a–b . This drive process will be immobilized when the Au QDs reach a critical size (as a result of the Au returning to the solid phase). This would be consistent with experimental observations that larger Au particles seen in the minority multilayer regions of the nanograin film appear to be encapsulated by the MoS 2 film, suggesting they were immobilized after reaching a critical size, leading the MoS 2 to grow over them rather than driving them forward (see Supplementary Figs.  15c, d and 16 ).

a Schematic of the climb stage: the Au QDs droplets tend to climb onto the MoS 2 monolayer from the SiO 2 substrate due to the surface tension difference. b Schematic of the drive stage: the growth of a second MoS 2 layer tends to push the Au QDs along the growth direction. Most of the Au QDs will be stabilized at the grain boundaries in order to minimize their surface energy. c SEM image showing the distribution of the Au QDs on the edges of growing MoS 2 layers (marked by the red-dashed line). Au QDs are usually larger on the MoS 2 layers (indicated by blue arrows) than those observed on the SiO 2 substrate. d Real-time phase-field simulation showing the formation of the MoS 2 nanograin film. The Au QDs numbered 1–4 show that the growth of the MoS 2 layer will push Au QDs toward the grain boundaries. Au QDs numbered 5–6 show a zipper effect to suture the neighboring grains together. The movie can be found in Supplementary Movie  1 (Real-time phase-field simulation of the growth process).

Figure  3d shows successive snapshots that illustrate the growth of TMD nanograins films via a phase-field simulation (also see Supplementary Movie  1 ). The Au QDs serve as nucleation sites to form the first MoS 2 layer (Fig.  3d–i ). The growth of a second MoS 2 layer will then drive the Au QDs. Some of the Au QDs will be immobilized at the GBs (Fig.  3d ii and iii , number 1–4). More interestingly, we found that neighboring grains could be sutured together through a zipper effect (Fig.  3d iii-v , number 5 and 6, see Method and Supplementary Fig.  17 ). It is important to emphasize that the growth model for the TMD nanograin films we describe here is driven by a 0D/2D interaction at the atomic limit, which is distinct from the conventional metal-film-assisted growth of 2D materials (2D/2D) 29 , 35 , or from the metal-nanoparticle-assisted growth of nanowires (0D/1D) 49 .

Hydrogen production of the TMD nanograin films

We first performed first-principle calculations 50 to examine all of the HER catalytic active sites in model MoS 2 geometries (Fig.  4a ), including the basal plane sites, edges with different passivation (Mo or S) 51 , and GBs with different atomic structural configurations (5|7, 6|8, 4|6, 12|4, and 8|4|4 rings). The calculated hydrogen adsorption free energies ( ∆G H ) 52 of various atomic structures (Supplementary Fig.  18 ) are shown in Fig.  4b and Supplementary Table  2 . The MoS 2 basal plane has a ∆G H as high as 1.79 eV 15 , indicating a HER-inert surface. For edges and GBs, we compare the activities of most of the experimentally observed structures, e.g., 5|7 (0.132 eV), 6|8 (−0.237 eV), and 8|4|4 (0.52 eV, −0.044 for defect) ring combinations, 2 , 53 50% S passivated Mo edge (0.561 eV), and 50% passivated S edge (0.446 eV) 53 , 54 . It can be seen that GBs indeed show comparable or even better activity than the edges in general, suggesting GBs are promising candidates as high-efficient catalyst sites. In addition, we note that some edges are not further considered here owing to their instability in air; for instance, the otherwise energetically favorable Mo edge without passivation of S atoms (−0.446 eV).

a Schematic of MoS 2 with an illustration of the main types of catalytically active sites for the hydrogen evolution reaction (HER) including: basal plane atomic sites, edges, and GBs. b corresponding ∆G H of various types of catalytically active sites in MoS 2 catalysts. Some types of GBs (e.g., 844, 6–8, and 5–7) show better electrocatalytic activity than edges.

We then developed a micro-electrochemical cell to investigate the HER activity of our sub-10 nm nanograin films. Furthermore, the HER activity of model single-GBs (mirror GB) 3 , 11 , single-edge (Supplementary Fig.  19 ) 53 , and basal plane structures in MoS 2 samples from CVD method 1 were also examined for comparison. Figure  5a illustrates the micro-cell’s structure. Figure  5b shows a picture of the micro-cell and electrodes configuration where a graphite counter electrode and a micro-reference electrode were used. In our experiment, a vertical MoS 2 /graphene heterostructure was designed, in which graphene played two important roles: one is to efficiently inject electrons to MoS 2 . Such a strategy has already been widely adopted in TMD-based semiconductor devices 55 , 56 , while the other is to provide a fair comparison by eliminating the contact resistance variations of the individual MoS 2 electrodes onto the substrate. Previous work 56 , 57 suggested the formation of an inconsistent contact barrier between the MoS 2 and the Au electrode, owing to their complex metal–semiconductor interface leading to effects such as Fermi pinning, creation of alloy structures, etc. The formation of such a contact barrier was also studied by phase engineering 21 and field-effect gating 58 in the HER process. Considering the important role of this graphene supporting layer, we introduced one more working electrode in the micro-electrochemical cell (four-electrode micro-electrochemical cell 59 , 60 , see circuit diagram in Supplementary Fig.  20 ) to monitor in situ its conductance. The device fabrication of the MoS 2 /graphene heterostructure is detailed in Supplementary Figs.  21 and 22 , and Supplementary Notes  2 and 3 . The heterostructure was also characterized by Raman spectroscopy (Supplementary Fig.  23 ). Figure  5c–f shows the optical images obtained from the typical devices fabricated with a MoS 2 nanograin film, a single-GB model structure (the presence of a GB was confirmed by Raman mapping, as illustrated on Fig.  5d , right panel), a single edge, and a basal plane, respectively. In these optical images, only the exposed MoS 2 in the reaction window contributes to the electrocatalytic reaction, as indicated by the white arrows. The rest of the areas are covered by poly(methylmethacrylate) (PMMA), and the graphene supporting layer (Supplementary Fig.  24 ) are also electrochemically inert.

a Schematic of the micro-electrochemical cell for HER measurements, where graphene serves as a vertical electron injector. b Photograph of the micro-electrochemical cell. c Optical images of the MoS 2 nanograin device, consisting of a PMMA reaction window, a MoS 2 nanograin film, a graphene supporting layer, and a SiO 2 /Si substrate from top to bottom. d – f Optical images of the MoS 2 devices with a single GB ( d ), a single edge ( e ), and basal plane ( f ), respectively. In these devices, the HER process occurs at the exposed widows on the PMMA passivation film as indicated by the white arrow. g , h Polarization curves of the current density ( g ) and the corresponding Tafel plots ( h ) of the MoS 2 devices. The window size is ~80 μm 2 for each device.

Figure  5g, h shows the polarization curves and the corresponding Tafel slopes in a 0.5 M H 2 SO 4 solution and provide several interesting observations. First, the resistivity of the graphene supporting layer is found to be as low as 10 −4  Ω mm during the HER (see the inset in Fig.  5g ), indicating the excellent current injection performance of the graphene layer. Secondly, the single-GB device delivers a better activity than the single-edge device (see also the measurement of the active length per unit area of the GB in Supplementary Fig.  25 ), and both are superior to the basal plane device, echoing our theoretical calculation results shown above. More importantly, our MoS 2 nanograin film shows a remarkable HER performance: −25 mV and 54 mV dec −1 for the onset potential and the Tafel slope, respectively.

In order to evaluate the HER performance of the TMD nanograin film accurately, we have fabricated hundreds of devices and tested them in a size-controlled micro-electrochemical cell. The HER data, including current density, Tafel slope, and onset potential of the nanograin film and other types of electrostatically active sites are presented in Fig.  6a–c and Supplementary Fig.  26 . The window size is controlled from 25 to 150 μm 2 . Our results demonstrate the excellent HER performance of the nanograin film. The HER current density of the nanograin film is up to ~1000 mA cm − 2 while the Tafel slope and onset potential are down to ~50 mV dec −1 and ~−25 mV, thus exhibiting a performance superior to devices using the basal plane of the CVD film, the single edge and the single GB. This suggests that sites on the otherwise HER-inert basal plane of the MoS 2 have been activated by the presence of GBs. The obtained HER performance is also comparable to the performance of strained sulfur vacancies 15 or the metastable 1T-phase 61 , which are arguably more challenging to realize experimentally.

a – c Current density ( a ), Tafel plots ( b ), and onset potentials ( c ) for the MoS 2 devices with serial window sizes from 25–150 μm 2 . The nanograin films show higher current density, lower Tafel slope, and onset potential compared to the basal plane, single edge, and single GB.

Finally, as a comparison, we have conducted conventional macro-cell measurements on a MoS 2 nanograin film on a glassy carbon electrode, as shown in Supplementary Figs.  27 – 32 and Supplementary Notes  4 and 5 . A remarkable HER performance (−44 mV and 55 mV dec −1 for the onset potential and the Tafel slope, respectively) of the MoS 2 nanograin film is observed (Supplementary Fig.  28 ), which is consistent with the results obtained in the micro-cell. It also shows an excellent long-term stability for HER, as shown in Supplementary Fig.  29 . It is worth mentioning that Au single-atom exists in our MoS 2 nanograin film and may contribute to the overall HER performance. However, based on the structure (Supplementary Fig.  6 ) and extremely low content (Supplementary Fig.  7 ) of Au single atoms as well as their low HER-activity (Supplementary Figs.  33 and 34 , and Supplementary Note  6 ), we conclude that the residual Au is not the main contributor of the good HER activity in MoS 2 nanograin film. The measured performance should be mainly from GB due to its superior activity and ultra-high density. As a proof-of-concept practical application, hydrogen production on a wafer-size (2 inch) MoS 2 nanograin film was also demonstrated (see Supplementary Movie  2 ). A large amount of H 2 bubbles could be produced during the reaction, demonstrating the tantalizing potential for practical application of GBs-engineered catalysts for hydrogen production.

In summary, we have engineered wafer-size ultra-high-density GBs at the 2D limit in TMD films. Phase-field simulations reveal a climb and drive 0D/2D-interaction-based growth mechanism. As a proof-of-concept application in electrocatalysis, devices based on our TMD nanograin films deliver a superior hydrogen evolution performance (onset potential: −25 mV and Tafel slope: 54 mV dec −1 ), indicating an intrinsic activation of the GB-rich 2D plane. Beyond electrocatalysis, the nanograin films may provide further diverse potential applications, such as in resistance-memory devices, flexible devices, or for use as mechanical films and molecular sieving films.

Growth of water-size MoS 2 and WS 2 nanograin films

The first step consists in depositing Au layers of various thicknesses on clean 2-inch sapphire or SiO 2 /Si wafers by e-beam evaporation at the rate of 0.1 Å s −1 ; these are used as the growth substrate for the nanograin films. Subsequently, Au-coated wafers are introduced into a CVD apparatus, which is depicted in detail in Supplementary Fig.  2 . MoO 3 or WO 3 powders were loaded in an aluminum oxide boat and used as sources. Sulfur powder is placed in a second aluminum oxide boat upstream of the MoO 3 or WO 3 in the outer tube. In a third step, Ar (flow 100 sccm) and Ar/water vapor (flow 50 sccm) are introduced as carrier gases for the growth in the inner and outer tubes, respectively. The growth temperature was kept at 780 °C in low-vacuum conditions of 1–10 kPa; the S powder is separately kept at ~160 °C. Adjusting the growth time is used to control the thickness of films. For example, Condition (I): 3–5 min of the growth time for 1–3 layers (1–3L); Condition (II): 5–10 min of the growth time for 4–5 layers (4-5 L). The growth parameters for WS 2 nanograin films are similar to those used for MoS 2 , except for a higher growth temperature (900 °C). Finally, Au nanoparticles can be completely removed from as-grown nanograin MoS 2 or WS 2 films by using a KI/I 2 etcher for 40–60 min at room temperature (see Supplementary Figs.  5 and 6 ).

Special care needs to be taken on two important points: (i) the Au QDs will form once the temperature is above 300 °C, due to the solid-state dewetting behavior 42 , 43 , 62 at the interface of SiO 2 (Al 2 O 3 ) and Au (see Supplementary Fig.  3 ). Prior to the TMD growth, the substrate therefore stabilizes at this temperature to ensure the presence of Au QDs. (ii) Small amounts of water vapor are used to prevent the Mo or W sources from complete sulfurization before they reach the growth wafer, thus enabling a steady supply of sources during growth.

Theoretical calculations of ∆G H

The structural optimizations were carried out by adopting the generalized gradient approximation (GGA) with the Perdew–Burke–Ernzerhof (PBE) exchange-correlation functional, along with projector-augmented wave (PAW) potentials. The electronic wave functions were expanded in a plane wave basis set with a kinetic energy cutoff of 300 eV. For the Brillouin zone integration, 1 × 5 × 1 Monkhorst–Pack k -point meshes were used. The energy convergence criterion for the electronic wavefunction was set at 10 −5  eV. A vacuum distance of about 10 Å, both laterally between MoS 2 ribbons and vertically between layers, was chosen when constructing the supercell to minimize possible spurious interaction between ribbons due to the periodic boundary conditions.

The hydrogen adsorption energy is defined as:

where \(E_{({\mathrm{MoS}_{2}} + {\mathrm{H}})}\) is the energy of the MoS 2 system with a H atom adsorbed, \(E_{({\mathrm{MoS}_{2}})}\) is that of the MoS 2 system without adsorbed H, and \(E_{({\mathrm{H}_{2}})}\) is that of gas phase H 2 molecule.

The hydrogen adsorption free energy \(\Delta G_{\mathrm{H}}\) was defined as:

where Δ E H is the hydrogen adsorption energy, Δ E ZPE is the zero-point energy difference, T is the temperature, and Δ S H is the entropy difference for hydrogen between the adsorbed state and the gas phase. The entropy of hydrogen adsorption is calculated as \(\Delta S_{\mathrm{H}} = 1/2S_{{\mathrm{H}_{2}}}\) , where \(S_{{\mathrm{H}_{2}}}\) is the entropy of hydrogen molecule in the gas phase at standard conditions. We have calculated the vibrational frequencies of H adsorbed on different MoS 2 systems, using finite differences to determine the Hessian matrix. According to the Sabatier principle, a good catalyst should bind a key intermediate in a manner that it is strong enough to allow the reagent H atoms to stay on the MoS 2 system for reaction but weak enough to enable the release of any produced H 2 molecules.

Phase-field simulations: simulation of the drive mechanism: MoS 2 pushing Au QDs long the growth direction

The equilibrium status of the Au QDs at atomic steps dictates they can be pushed by the MoS 2 growth front or not, which is crucial to dynamic evolution process. We undertook a case study of the equilibrium status of a Au QD at the edge of a MoS 2 grain.

We adopt a phase-field model 63 to simulate how the wetting process influences the equilibrium shape profile. The free energy function is given as,

where ρ is the phase field for the density of the liquid, V is the space volume, C  = 10 and κ  = 1 are constants of bulk energy and surface tension, respectively. ρ  = 1 denotes the liquid phase and ρ  = 0 denotes the vapor phase. The term \(\lambda \xi (\frac{{\rho ^3}}{3} - \frac{{\rho ^2}}{2})\) provides the driving force to the Au particle evolution with a coupling constant λ  = 180, the dimensionless supersaturation ξ is set as ( v 0 – v t )/ v vapor , where v 0 is the initial volume of the particle, v t is the real time volume of the particle, and v vapor is the volume of vapor. The evolution equation in the Allen–Cahn form is,

This evolution equation differs from that of Borcia et al.’s work 63 , which considers full fluid dynamics by including extra terms from the Navier–Stokes equation. In our study, we focus on the equilibrium and do not need to consider the exact evolution kinetics, thus leading to a simpler form of the evolution equation.

We control the contact angle between the Au QD and the substrate by fixing ρ  =  ρ S at the substrate. Then, the contact angle can be analytically obtained via the equation \(\cos \theta = - 1 + 6\rho _{\mathrm{S}}^2 - 4\rho _{\mathrm{S}}^3\) . We set ρ S  = 0.5 for the MoS 2 substrate as corresponding to a contact angle between the Au and the substrate of 90°. Likewise, we set ρ S  = 0.21 for the MoS 2 substrate as corresponding to a contact angle between the Au and the substrate of 138°.

We solve the evolution equation in Matlab by discretization with a time step of 0.01 and let the simulation run for a long enough time to ensure the system reaches equilibrium. The particle has a diameter of ~11 nm if resting on the MoS 2 sheet.

The result shows that this Au QD will stay on the lower part of the MoS 2 atomic step, but keeps in contact with the step edge. This suggests that if in the liquid phase, the Au QD would be pushed forward by the growth front of the second MoS 2 layer during the growth.

Phase-field simulations: simulation of the zipper effect suturing the neighbor nanograins

If the Au QDs can be pushed by the growing MoS 2 grains, they are likely to coalesce when two growth fronts meet, resulting in a complicated evolution. To obtain insights into this process, we perform another phase-field simulation. The phase-field model for the grain evolution and the set of parameters is almost identical to that reported in the literature for a similar system 64 . We ignore the evolution of the supersaturation and keep it at a constant of 0.1 everywhere. A detailed formulation is omitted here for brevity and can be found in other work 65 , 66 . In addition to the grains, we pay attention to the evolutions of the Au QDs. These QDs are modeled as hemispherical balls sliding on the substrate. For simplicity, we ignore the effect of the Au QDs on the growth of MoS 2 grains. Then these QDs move on the substrate at a speed of v  = 50 tan −1 ( t )/ π , where t is the dimensionless time of evolution. When two QDs meet, they merge and form a larger QD. The volume of the new QD is the sum of the volumes of the previous two QDs and its center is at the midpoint between the centers of the previous two QDs. We then solve the evolution equation in Matlab by discretization with a time step of 0.01. Initially Au QDs are randomly placed on the substrate and there is in the vicinity of these QDs, a circular nucleus of MoS 2 with a radius of r 0 and a randomly assigned grain orientation.

To make the simulation consistent with the experiment, we extract some parameters from the experimental data. The experiment shows that the QDs are typically 4–5 nm. It also shows that when observed on top of the MoS 2 sheet, the Au QDs are larger and less dense by around 2.5 times. This suggests by pushing Au QDs on top of MoS 2 sheet, the final number of QDs will be 2/5 of the initial QDs. A subtle issue involved in the simulation is the ratio between n Au (initial number of Au QDs) and n n (the number of nuclei). To address this issue, we take a simulation case study to find a reasonable ratio. By fixing n n  = 5, we obtain that when n Au  = 9.2, the final number of QDs is 2/5 of the initial QD number, consistent with the experiments. We therefore choose n Au  = 10 in further simulations.

Device fabrication procedure

First, large scale and high-quality single-layer graphene films were grown on Cu foils by CVD, and then transferred onto the prepatterned chips through a standard PMMA-assisted transfer method. Second, electron beam lithography (EBL) and O 2 plasma were employed to pattern the graphene film into isolated strips with desired size, shape, and location. Third, as-grown MoS 2 films were transferred on the patterned graphene strips, also using a PMMA-assisted transfer method. For the nanograin films, an additional step carefully positioning an EBL-defined PMMA template was needed. Fourth, an annealing process at 200 °C under high vacuum condition (1 × 10 −5 Torr) was conducted to remove trapped residual molecules between the graphene and MoS 2 films to optimize their interfaces. Fifth, EBL followed by thermal evaporation was employed to fabricate the electrodes (Cr/Au, 2 nm/60 nm) on graphene to connect the device with the Au contact pads. Finally, the entire device as fabricated on the chip was passivated by a 500 nm PMMA film, followed by EBL to remove the PMMA above the region of interest in the MoS 2 film, and expose this section and this section only of the active catalyst to the electrolyte in HER test. Any electrochemical activity only occurs within the exposed window on the nanosheet, while the rest of regions, including the electrodes and the contact pads, are fully passivated by the PMMA to ensure a full electrochemical inertness.

A mirror GB in CVD-grown MoS 2 is chosen as representative of single GBs, due to the ease with which it can be distinguished in optical images. It also consists mainly of 8|4|4 or defect 8|4|4 rings 3 , 11 and has therefore a similar atomic structure to many of the complex boundaries observed in the nanograin films. The length–width ratio of the exposed window for the single-GB device or for the edge device (exposing the edge of an otherwise single-crystalline patch of CVD-grown MoS 2 ) is fixed at a ratio of about 2:1, and the length of the window is comparable to the length of the edge or GB in experiments.

Four-electrode micro-electrochemical measurements

A micro-electrochemical cell with four electrodes (the circuit diagram for this cell is introduced in Supplementary Fig.  20 ) was developed in our experiment. Among the four electrodes, two serve as counter and reference electrode, using pencil graphite and a Ag/AgCl micro reference electrode (Harvard Apparatus), respectively. The remaining two electrodes were connected to the graphene supporting layer to monitor the conductance signal of graphene and the electrocatalytic signal of MoS 2 during HER. The device fabrication procedure is shown in Supplementary Figs.  21 and 22 , and Supplementary Notes  2 and 3 . In all experiments, only the exposed region of the MoS 2 nanosheets contributes to the recorded HER performance.

The typical four-electrode micro-electrochemical measurements were conducted in a 0.5 M H 2 SO 4 electrolyte solution. The scan rate is 5 mV step −1 . Both electrocatalytic current ( I c ) and conductance current ( I ds ) are simultaneously detected. The leakage current in I c is about 10 −10  A (the electrochemical signal without any exposed window, i.e., from a region with PMMA passivation). The electrochemical current density is calculated by normalizing the current to the area of the exposed window on the MoS 2 . In 0.5 M H 2 SO 4 solution, we measured:

Standard macro-electrochemical measurements

The macro-electrochemical measurements were conducted on a glassy carbon electrode (3 mm in diameter). Using a similar PMMA-assisted transfer method as described above, we transferred CVD-grown graphene and MoS 2 nanograin films layer-by-layer on a glassy carbon electrode, and the sample area exceeding to glassy carbon electrode was scraped off before measuments. A standard three electrode system was used, and a Pt plate and Ag/AgCl rod served as counter and reference electrode, respectively. The measurement was performed on a biological electrochemical station in H 2 -saturated 0.5 M H 2 SO 4 solution. Linear sweep voltammetry (LSV) was conducted at a scan rate of 5 mV s −1 . The onset potential is defined as the beginning potential of Tafel linear region. The stability test was carried out by taking continuous potential cycling in the potential window of −0.181 to ~0.219 V (versus RHE) at a scan rate of 100 mV s −1 . The presented LSV curves in macro-electrochemical cell were iR-corrected by subtracting the ohmic resistance loss (about 9 Ω), the value of which is obtained from electrochemical impedance spectroscopy measurement (Supplementary Fig.  32 ). To evaluate the electrochemically active surface area of the catalysts, the Cu underpotential deposition method 67 , 68 , 69 , 70 was applied, as shown in Supplementary Fig.  30 and Supplementary Table  3 .

To test the faradaic efficiency, the H 2 products were analyzed with an online gas chromatography (GC) setup (Agilent 7890B) equipped with a thermal conductivity detector. Argon (≥99.999%) was used as the carrier gas. The Faradaic efficiency (FE) was calculated by comparing the measured amount of H 2 generated by cathodal electrolysis with the calculated amount of H 2 (assuming an FE of 100%), and the equation is given by:

where the moles of H 2 is measured by GC (the calibration is needed in advance through injecting high-purity H 2 to GC, and the H 2 volume is linearly dependent on GC peak area), Q is obtained from the electrochemical measurements.

Material characterization

The microstructures and morphologies based on MoS 2 were characterized by optical microscopy, SEM (FEI 4200), and Raman spectroscopy (Renishaw inVia). Raman spectroscopy was performed with a 514.5 nm laser (with spot size about 1 μm in diameter) at room temperature. The micro-electrochemical measurements were performed using two source meters (Keithley 2400 and 2450). STEM measurements were performed at the SuperSTEM Laboratory, Daresbury, UK, on a Nion UltraSTEM100 aberration-corrected dedicated STEM. The Nion UltraSTEM has a cold field emission gun with a native energy spread of 0.35 eV and was operated at 60 kV acceleration voltage. The beam was set up to a convergence semiangle of 30 mrad with an estimated beam current of ~100 pA. Under these operating conditions, the estimated probe size is ~1.1 Å, providing the perfect tool for atom-by-atom chemical analysis. HAADF imaging was employed to produce atomically resolved images whose intensity is approximately proportional to the square of the average atomic number Z of the material under investigation. This chemically sensitive “Z-contrast” mode is ideally suited to directly identify the nature of individual atoms. All the HAADF images presented here were filtered for visual clarity using a wiener filter with coefficient sigma ranging from 2–5 and 50 iterations in MEM (Maximum Entropy Methods) with fourth order Gaussian filter, as implemented in the STEM-CELL software 71 , 72 . Geometrical Phase Analysis (GPA) was carried out using a cosine mask of suitable size and 1 binning base. HRTEM and low-resolution annular dark field (ADF)-STEM images were obtained on a FEI Tecnai F20 field emission gun microscope with a 0.19 nm point-to-point resolution at 200 kV, equipped with an embedded Quantum Gatan Image Filter for EELS analyses. Some STEM characterizations were carried out on a JEOL ARM-200F (S)TEM equipped with CEOS CESCOR aberration corrector, operated at an accelerating voltage of 80 kV. The convergence semiangle and acquisition semiangle were 28–33 and 68–280 mrad for the ADF imaging. The dwell time per pixel was set to 12–20.

Data availability

The data that support the findings of this study are available from the corresponding author upon reasonable request.

Yu, Q. et al. Control and characterization of individual grains and grain boundaries in graphene grown by chemical vapour deposition. Nat. Mater. 10 , 443 (2011).

ADS   CAS   PubMed   Google Scholar  

Najmaei, S. et al. Vapour phase growth and grain boundary structure of molybdenum disulphide atomic layers. Nat. Mater. 12 , 754–759 (2013).

van der Zande, A. M. et al. Grains and grain boundaries in highly crystalline monolayer molybdenum disulphide. Nat. Mater. 12 , 554–561 (2013).

ADS   PubMed   Google Scholar  

Chang, S.-Y. & Chang, T.-K. Grain size effect on nanomechanical properties and deformation behavior of copper under nanoindentation test. J. Appl. Phys. 101 , 033507 (2007).

ADS   Google Scholar  

Son, D.-Y. et al. Self-formed grain boundary healing layer for highly efficient CH 3 NH 3 PbI 3 perovskite solar cells. Nat. Energy 1 , 16081 (2016).

ADS   CAS   Google Scholar  

Li, C. et al. Grain-boundary-enhanced carrier collection in CdTe solar cells. Phys. Rev. Lett. 112 , 156103 (2014).

Bowman, W. J., Zhu, J., Sharma, R. & Crozier, P. A. Electrical conductivity and grain boundary composition of Gd-doped and Gd/Pr co-doped ceria. Solid State Ion. 272 , 9–17 (2015).

CAS   Google Scholar  

Guo, C. F., Sun, T., Liu, Q., Suo, Z. & Ren, Z. Highly stretchable and transparent nanomesh electrodes made by grain boundary lithography. Nat. Commun. 5 , 3121 (2014).

Sangwan, V. K. et al. Gate-tunable memristive phenomena mediated by grain boundaries in single-layer MoS 2 . Nat. Nanotechnol. 10 , 403 (2015).

Feng, X., Jiang, K., Fan, S. & Kanan, M. W. Grain-boundary-dependent CO 2 electroreduction activity. J. Am. Chem. Soc. 137 , 4606–4609 (2015).

CAS   PubMed   Google Scholar  

Li, G. et al. All the catalytic active sites of MoS 2 for hydrogen evolution. J. Am. Chem. Soc. 138 , 16632–16638 (2016).

Seh, Z. W. et al. Combining theory and experiment in electrocatalysis: insights into materials design. Science 355 , eaad4998 (2017).

Voiry, D., Yang, J. & Chhowalla, M. Recent strategies for improving the catalytic activity of 2D TMD nanosheets toward the hydrogen evolution reaction. Adv. Mater. 28 , 6197–6206 (2016).

Liu, Y. et al. Self-optimizing, highly surface-active layered metal dichalcogenide catalysts for hydrogen evolution. Nat. Energy 2 , 17127 (2017).

Li, H. et al. Activating and optimizing MoS 2 basal planes for hydrogen evolution through the formation of strained sulphur vacancies. Nat. Mater. 15 , 48–53 (2016).

Kibsgaard, J., Chen, Z., Reinecke, B. N. & Jaramillo, T. F. Engineering the surface structure of MoS 2 to preferentially expose active edge sites for electrocatalysis. Nat. Mater. 11 , 963 (2012).

Ji, Q., Zhang, Y., Zhang, Y. & Liu, Z. Chemical vapour deposition of group-VIB metal dichalcogenide monolayers: engineered substrates from amorphous to single crystalline. Chem. Soc. Rev. 44 , 2587–2602 (2015).

Jaramillo, T. F. et al. Identification of active edge sites for electrochemical H 2 evolution from MoS 2 nanocatalysts. Science 317 , 100–102 (2007).

Xie, J. et al. Defect-rich MoS 2 ultrathin nanosheets with additional active edge sites for enhanced electrocatalytic hydrogen evolution. Adv. Mater. 25 , 5807–5813 (2013).

Dai, X. et al. Co-doped MoS 2 nanosheets with the dominant CoMoS phase coated on carbon as an excellent electrocatalyst for hydrogen evolution. ACS Appl. Mater. Interfaces 7 , 27242–27253 (2015).

Voiry, D. et al. The role of electronic coupling between substrate and 2D MoS 2 nanosheets in electrocatalytic production of hydrogen. Nat. Mater. 15 , 1003–1009 (2016).

Ye, G. et al. Defects engineered monolayer MoS 2 for improved hydrogen evolution reaction. Nano Lett. 16 , 1097–1103 (2016).

Ouyang, Y. et al. Activating inert basal planes of MoS 2 for hydrogen evolution reaction through the formation of different intrinsic defects. Chem. Mater. 28 , 4390–4396 (2016).

Li, S. et al. Vapour–liquid–solid growth of monolayer MoS 2 nanoribbons. Nat. Mater. 17 , 535–542 (2018).

Zhou, J. et al. A library of atomically thin metal chalcogenides. Nature 556 , 355–359 (2018).

Gong, Y. et al. Synthesis of millimeter-scale transition metal dichalcogenides single crystals. Adv. Funct. Mater. 26 , 2009–2015 (2016).

He, Y. et al. Layer engineering of 2D semiconductor junctions. Adv. Mater. 28 , 5126–5132 (2016).

Wang, X. et al. Chemical vapor deposition growth of crystalline monolayer MoSe 2 . ACS Nano 8 , 5125–5131 (2014).

Gao, Y. et al. Ultrafast growth of high-quality monolayer WSe 2 on Au. Adv. Mater. 29 , 1700990 (2017).

Google Scholar  

Li, S. et al. Halide-assisted atmospheric pressure growth of large WSe 2 and WS 2 monolayer crystals. Appl. Mater. Today 1 , 60–66 (2015).

Dathbun, A. et al. Large-area CVD-grown sub-2 V ReS 2 transistors and logic gates. Nano Lett. 17 , 2999–3005 (2017).

Zhou, J. et al. Large-area and high-quality 2D transition metal telluride. Adv. Mater. 29 , 1603471 (2017).

Naylor, C. H. et al. Monolayer single-crystal 1T’-MoTe 2 grown by chemical vapor deposition exhibits weak antilocalization effect. Nano Lett. 16 , 4297–4304 (2016).

ADS   CAS   PubMed   PubMed Central   Google Scholar  

Cui, F. et al. Epitaxial growth of large-area and highly crystalline anisotropic ReSe 2 atomic layer. Nano Res. 10 , 2732–2742 (2017).

Gao, Y. et al. Large-area synthesis of high-quality and uniform monolayer WS 2 on reusable Au foils. Nat. Commun. 6 , 8569 (2015).

Empante, T. A. et al. Chemical vapor deposition growth of few-layer MoTe 2 in the 2H, 1T’, and 1T phases: tunable properties of MoTe 2 films. ACS Nano 11 , 900–905 (2017).

Feng, Q. et al. Growth of large-area 2D MoS 2(1-x) Se 2x semiconductor alloys. Adv. Mater. 26 , 2648–2653 (2014).

Wu, S. et al. Vapor–solid growth of high optical quality MoS 2 monolayers with near-unity valley polarization. ACS Nano 7 , 2768–2772 (2013).

Huang, J.-H. et al. Large-area 2D layered MoTe 2 by physical vapor deposition and solid-phase crystallization in a tellurium-free atmosphere. Adv. Mater. Interfaces 4 , 1700157 (2017).

Kang, K. et al. High-mobility three-atom-thick semiconducting films with wafer-scale homogeneity. Nature 520 , 656 (2015).

Kang, K. et al. Layer-by-layer assembly of two-dimensional materials into wafer-scale heterostructures. Nature 550 , 229 (2017).

Lee, S.-H., Kwak, E.-H. & Jeong, G.-H. Dewetting behavior of electron-beam-deposited Au thin films on various substrates: graphenes, quartz, and SiO 2 wafers. Appl. Phys. A 118 , 389–396 (2015).

Seguini, G. et al. Solid-state dewetting of ultra-thin Au films on SiO 2 and HfO 2 . Nanotechnology 25 , 495603 (2014).

Li, H. et al. From bulk to monolayer MoS 2 : evolution of Raman scattering. Adv. Funct. Mater. 22 , 1385–1390 (2012).

Huang, P. Y. et al. Grains and grain boundaries in single-layer graphene atomic patchwork quilts. Nature 469 , 389 (2011).

Shi, J., Ji, Q., Liu, Z. & Zhang, Y. Recent advances in controlling syntheses and energy related applications of MX 2 and MX 2 /graphene heterostructures. Adv. Energy Mater. 6 , 1600459 (2016).

Shi, Y., Li, H. & Li, L.-J. Recent advances in controlled synthesis of two-dimensional transition metal dichalcogenides via vapour deposition techniques. Chem. Soc. Rev. 44 , 2744–2756 (2015).

Ruffino, F. & Grimaldi, M. G. Controlled dewetting as fabrication and patterning strategy for metal nanostructures. Phys. Status Solidi A 212 , 1662–1684 (2015).

Choi, H.-J. Vapor–Liquid–Solid Growth of Semiconductor Nanowires (Springer Berlin Heidelberg, 2012).

Kresse, G. & Hafner, J. Ab initio molecular-dynamics simulation of the liquid-metal-amorphous-semiconductor transition in germanium. Phys. Rev. B 49 , 14251–14269 (1994).

Lauritsen, J. V. et al. Size-dependent structure of MoS 2 nanocrystals. Nat. Nanotechnol. 2 , 53–58 (2007).

Nørskov, J. K. et al. Trends in the exchange current for hydrogen evolution. J. Electrochem. Soc. 152 , J23–J26 (2005).

Zhou, W. et al. Intrinsic structural defects in monolayer molybdenum disulfide. Nano Lett. 13 , 2615–2622 (2013).

Hansen, L. P. et al. Atomic-scale edge structures on industrial-style MoS 2 nanocatalysts. Angew. Chem. Int. Ed. 50 , 10153–10156 (2011).

Yu, W. J. et al. Vertically stacked multi-heterostructures of layered materials for logic transistors and complementary inverters. Nat. Mater. 12 , 246–252 (2013).

Allain, A., Kang, J., Banerjee, K. & Kis, A. Electrical contacts to two-dimensional semiconductors. Nat. Mater. 14 , 1195 (2015).

Kang, J., Liu, W., Sarkar, D., Jena, D. & Banerjee, K. Computational study of metal contacts to monolayer transition-metal dichalcogenide semiconductors. Phys. Rev. X 4 , 031005 (2014).

Wang, J. et al. Field effect enhanced hydrogen evolution reaction of MoS 2 nanosheets. Adv. Mater. 29 , 1604464 (2017).

Ding, M. et al. An on-chip electrical transport spectroscopy approach for in situ monitoring electrochemical interfaces. Nat. Commun. 6 , 7867 (2015).

He, Y. et al. Self-gating in semiconductor electrocatalysis. Nat. Mater. 18 , 1098–1104 (2019).

Voiry, D. et al. Conducting MoS 2 nanosheets as catalysts for hydrogen evolution reaction. Nano Lett. 13 , 6222–6227 (2013).

Nsimama, P. D., Herz, A., Wang, D. & Schaaf, P. Influence of the substrate on the morphological evolution of gold thin films during solid-state dewetting. Appl. Surf. Sci. 388 A, 475–482 (2016).

Borciaa, R., B., I. D. & Bestehorn, M. Static and dynamic contact angles- a phase field modelling. Eur. Phys. J. Spec. Top. 166 , 127–131 (2009).

Artyukhov, V. I., Hu, Z., Zhang, Z. & Yakobson, B. I. Topochemistry of Bowtie- and star-shaped metal dichalcogenide nanoisland formation. Nano Lett. 16 , 3696–3702 (2016).

Steinbach, I. & Pezzolla, F. A generalized field method for multiphase transformations using interface fields. Phys. D Nonlinear Phenom. 134 , 385–393 (1999).

ADS   MathSciNet   MATH   Google Scholar  

Ingo, S. Phase-field models in materials science. Modell. Simul. Mater. Sci. Eng. 17 , 073001 (2009).

Zheng, Y. et al. High electrocatalytic hydrogen evolution activity of an anomalous ruthenium catalyst. J. Am. Chem. Soc. 138 , 16174–16181 (2016).

Mahmood, J. et al. An efficient and pH-universal ruthenium-based catalyst for the hydrogen evolution reaction. Nat. Nanotechnol. 12 , 441 (2017).

Green, C. L. & Kucernak, A. Determination of the platinum and ruthenium surface areas in platinum−ruthenium alloy electrocatalysts by underpotential deposition of copper. I. unsupported catalysts. J. Phys. Chem. B 106 , 1036–1047 (2002).

Colmenares, L., Jusys, Z. & Behm, R. J. Electrochemical surface characterization and O 2 reduction kinetics of Se surface-modified Ru nanoparticle-based RuSe y /C catalysts. Langmuir 22 , 10437–10445 (2006).

Grillo, V. & Rotunno, E. STEM-CELL: A software tool for electron microscopy: Part I—simulations. Ultramicroscopy 125 , 97–111 (2013).

Grillo, V. & Rossi, F. STEM-CELL: A software tool for electron microscopy. Part 2 analysis of crystalline materials. Ultramicroscopy 125 , 112–129 (2013).

Zhan, Y., Liu, Z., Najmaei, S., Ajayan, P. M. & Lou, J. Large-area vapor-phase growth and characterization of MoS 2 atomic layers on a SiO 2 substrate. Small 8 , 966–971 (2012).

Kong, D. et al. Synthesis of MoS 2 and MoSe 2 films with vertically aligned layers. Nano Lett. 13 , 1341–1347 (2013).

Lin, Y.-C. et al. Direct synthesis of van der Waals solids. ACS Nano 8 , 3715–3723 (2014).

Song, J.-G. et al. Layer-controlled, wafer-scale, and conformal synthesis of tungsten disulfide nanosheets using atomic layer deposition. ACS Nano 7 , 11333–11340 (2013).

Yim, C. et al. High-performance hybrid electronic devices from layered PtSe 2 films grown at low temperature. ACS Nano 10 , 9550–9558 (2016).

Liu, K.-K. et al. Growth of large-area and highly crystalline MoS 2 thin layers on insulating substrates. Nano Lett. 12 , 1538–1544 (2012).

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Acknowledgements

Z.L. gratefully acknowledges funding supports from Ministry of Education (MOE) under AcRF Tier 1 (M4011782.070 RG4/17 and M4011993.070 RG7/18), AcRF Tier 2 (2015-T2-2-007, 2016-T2-1-131, 2016-T2-2-153, and 2017-T2-2-136), AcRF Tier 3 (2018-T3-1-002), National Research Foundation – Competitive Research Program (NRF-CRP21-2018-0092), and A*Star QTE programme. P.T., S.M.S., J.R.M., and J.A. acknowledge funding from Generalitat de Catalunya 2017 SGR 327 and 1246 and the Spanish MINECO coordinated projects between IREC and ICN2 VALPEC (ENE2017-85087-C2-C3). ICN2 acknowledges support from the Severo Ochoa Program (MINECO, Grant SEV-2017-0706). IREC and ICN2 are funded by the CERCA Programme/Generalitat de Catalunya. Part of the present work has been performed in the framework of Universitat Autònoma de Barcelona Materials Science PhD program. S.M.S. acknowledges funding from ‘Programa Internacional de Becas “la Caixa”-Severo Ochoa’. J.R.M. recognizes also its affiliation to University of Barcelona. Part of the electron microscopy aspects of this work were supported by the EPSRC (UK), as the SuperSTEM Laboratory is the EPSRC National Research Facility for Advanced Electron Microscopy. Q.J.W. acknowledges the supports from MOE, Singapore grant (MOE2016-T2-2-159, MOE2016-T2-1-128, and MOE Tier 1 RG164/15) and National Research Foundation, Competitive Research Program (NRF-CRP18-2017-02) and NSFC (61704082) as well as Natural Science Foundation of Jiangsu Province (BK20170851). X.W. and B.K.T. gratefully acknowledge funding support from MOE, Singapore (grant no. MOE2015-T2-2-043). H.Z. acknowledges the supports from MOE under AcRF Tier 2 (MOE2015-T2-2-057; MOE2016-T2-2-103; and MOE2017-T2-1-162), AcRF Tier 1 (2016-T1-002-051; 2017-T1-001-150; and 2017-T1-002-119), Agency for Science, Technology and Research (A*STAR) under its AME IRG (Project No. A1783c0009), and NTU under Start-Up Grant (M4081296.070.500000) in Singapore. The authors would like to acknowledge the Facility for Analysis, Characterization, Testing, and Simulation, Nanyang Technological University, Singapore, for their electron microscopy and X-ray facilities. H.Z. also thanks the support from ITC via Hong Kong Branch of National Precious Metals Material Engineering Research Center, and the Start-Up Grant from City University of Hong Kong. Theory and simulations work at Rice university (Z.H., L.W., and B.I.Y.) was supported by the Office of Naval Research grant N00014-18-1-2182.

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These authors contributed equally: Yongmin He, Pengyi Tang, Zhili Hu, Qiyuan He.

Authors and Affiliations

School of Materials Science and Engineering, Nanyang Technological University, Singapore, 639798, Singapore

Yongmin He, Qiyuan He, Chao Zhu, Qingsheng Zeng, Prafful Golani, Wei Fu, Zhiqi Huang, Chao Zhu, Ao Zhang, Hua Zhang & Zheng Liu

Center for OptoElectronics and Biophotonics, School of Electrical and Electronic Engineering & The Photonics Institute, Nanyang Technological University, Singapore, 639798, Singapore

Yongmin He & Qi Jie Wang

Catalan Institute of Nanoscience and Nanotechnology (ICN2), CSIC and BIST, Campus UAB, Bellaterra, Barcelona, 08193, Catalonia, Spain

Pengyi Tang, Sara Martí-Sánchez & Jordi Arbiol

Catalonia Institute for Energy Research (IREC), Jardins de les Dones de Negre 1, Sant Adrià del Besòs, Barcelona, 08930, Catalonia, Spain

Pengyi Tang & Joan Ramon Morante

College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing, 210016, China

Department of Materials Science and NanoEngineering, Rice University, Houston, TX, 77005, USA

Zhili Hu, Luqing Wang & Boris I. Yakobson

Paul-Drude-Institut für Festkörperelektronik Leibniz-Institut im Forschungsverbund Berlin Hausvogteiplatz, 5-7, 10117, Berlin, Germany

Guanhui Gao & Achim Trampert

Centre for Micro-/Nano-electronics (NOVITAS), School of Electrical & Electronic Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore, 639798, Singapore

Caitian Gao & Zheng Liu

Institute of Fundamental and Frontier Sciences, University of Electronic Science and Technology of China, Chengdu, 610054, China

CNRS-International-NTU-THALES Research Alliance, Nanyang Technological University, Singaproe, 637553, Singapore

Xingli Wang & Beng Kang Tay

Institute of Flexible Electronics, Northwestern Polytechnical University, Xi’an, 710072, China

Xuewen Wang

SuperSTEM Laboratory, SciTech Daresbury Campus, Keckwick Lane, Daresbury, WA44AD, UK

Quentin M. Ramasse

School of Chemical and Process Engineering, University of Leeds, Leeds, LS29JT, UK

College of Materials Science and Engineering, Beijing University of Technology, Beijing, 100124, People’s Republic of China

Boxing An & Yongzhe Zhang

School of Environmental and Chemical Engineering, Shanghai University, Shanghai, 200444, China

Liang Wang & Minghong Wu

Department of Chemistry, City University of Hong Kong, Tat Chee Avenue, Kowloon, Hong Kong, China

CINTRA CNRS/NTU/THALES, UMI 3288, Research Techno Plaza, Singapore, 637553, Singapore

Qi Jie Wang & Zheng Liu

ICREA, Pg. Lluís Companys 23, Barcelona, 08010, Catalonia, Spain

Jordi Arbiol

Environmental Chemistry and Materials Centre, Nanyang Environment and Water Research Institute, Singapore, Singapore

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Contributions

Z.L. and Y.H. conceived and initiated the project. Z.L., Q.W., J.A. and M.W. supervised the project and led the collaboration efforts. Y.H. designed the experiments, grew nanograin film, fabricated the electrochemical devices, and performed micro-electrochemical HER measurement. P.T., C.Z. (Dr. Chao Zhu), S.M.S., J.A., J.M. and Q.M.R. performed the TEM and STEM measurements. Z.Hu did the phase-field simulation of the growth process. Y.H., Q.H. and C.Z. (Mr. Chao Zhu) made the micro-electrochemical measurement setup. L.W. (Luqing Wang) and B.Y. performed the first-principles calculations of the catalytic activity. Q.Z. did the SEM measurement. G.G. and A.T. conducted the cross-sectional TEM measurement. Z. Huang and Q.H. performed gas chromatography measurement. P.G., X.W. (Xuewen Wang), W.F., C.G., X.W. (Xingli Wang), B.A., Y.Z., L.W. (Liang Wang), and B.T. assisted the growth of nanograin film. J.X. and A.Z. did the Raman measurements. Y.H., P.T., Z.Hu, Q.W., H.Z., J.A. and Z.L. wrote the paper. All authors discussed the results and commented on the manuscript.

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Correspondence to Minghong Wu , Qi Jie Wang , Jordi Arbiol or Zheng Liu .

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He, Y., Tang, P., Hu, Z. et al. Engineering grain boundaries at the 2D limit for the hydrogen evolution reaction. Nat Commun 11 , 57 (2020). https://doi.org/10.1038/s41467-019-13631-2

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Broadening the design space of engineering materials through “additive grain boundary engineering”

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• Published: 22 January 2022
• Volume 57 , pages 9530–9540, ( 2022 )

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• Matteo Seita   ORCID: orcid.org/0000-0002-1852-2195 1 , 2 &
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Grain boundary engineering (GBE) is one of the most successful processing strategies to improve the properties of polycrystalline solids. However, the extensive thermomechanical processes involved during GBE restrict its use to selected applications and materials. In this viewpoint paper, we discuss the opportunity provided by additive manufacturing (AM) technology to broaden the applicability of the GBE paradigm and, consequently, the design space for engineering materials. By integrating specially-designed thermomechanical processing within AM, it would be possible to produce bulk, near-net-shape parts with complex geometry and GBE microstructure. We discuss the major challenges in this endeavor and propose some possible strategies to achieve this goal, which we refer to as “additive-GBE”.

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The promise of grain boundary engineering

Amongst the many processing strategies that have been conceived to improve the properties of polycrystalline materials, grain boundary engineering (GBE) deserves a special mention. By manipulating a small fraction of the atoms in the solid—namely, those which are located at grain boundaries (GBs)—GBE leads to dramatic changes in properties [ 1 , 2 ], including ductility [ 3 ], fatigue [ 4 ], creep [ 5 ], hydrogen embrittlement [ 6 , 7 ], and corrosion behavior [ 8 ].

GBE involves applying a sequence of thermomechanical processes to a target material, which typically consist of cyclic plastic deformation and high temperature treatments (Fig.  1 ). As such, GBE is a prototypical metal processing strategy, much like those employed in ancient times to improve the strength of metal alloys. The resulting microstructure exhibits a significantly different distribution of GBs, with higher fractions of low energy GBs [ 9 , 10 ]. The change in the GB character distribution, which can be thought of as a “survival of the fittest”, is the result of microstructure recovery and recrystallisation upon heat treatment [ 10 , 11 ]. During recrystallisation, new, strain free grains nucleate and grow into the surrounding microstructure, reducing the stored energy which was introduced through the mechanical deformation step in GBE. As these grains grow, they promote the formation of new GBs. Amongst those, low energy GBs tend to remain in the recrystallized microstructure because they typically exhibit reduced mobility [ 12 , 13 ].

Schematic of conventional grain boundary engineering

In materials that are characterized by a low-stacking fault energy, GBE yields recrystallized microstructures with copious coherent twin boundaries (TBs). Owing to their perfect atomic registry, coherent TBs have the lowest energy and the highest thermal stability amongst all GBs. As such, they are frequently observed in recrystallized microstructures and may survive the multiple strain-annealing cycles during GBE [ 14 , 15 ]. Coherent TB formation is associated with stacking sequence errors that occur during GB migration upon recrystallisation [ 16 , 17 , 18 ]. It follows that, as a recrystallized grain grows, multiple coherent TBs can form. These highly twinned grain clusters disrupt the connectivity of general, high-energy GBs, improving the GB-governed properties of the material. Besides coherent TBs, GBE may promote the formation of other low-energy GBs, including incoherent TBs and other twin-related GBs. These GBs form as different twinned grain clusters coalesce and interact with one another throughout the repeated strain-annealing cycles in GBE [ 13 ].

The concept of GBE was initially proposed by Watanabe et al . in the 1980s [ 19 ], convincingly demonstrated by Palumbo et al . on a variety of metal alloys a decade later [ 3 , 5 , 20 , 21 ], and widely explained by Randle et al . in the 2000s [ 12 , 13 , 22 , 23 ]. Since then, GBE has proliferated into a myriad of different adaptations and has been applied to a broad range of materials [ 8 , 20 , 23 , 24 ], including non-metals [ 25 , 26 ]. Despite the intense research and large number of success stories, however, very few modern industrial applications employ GBE materials [ 27 , 28 ].

What limits the broad application of GBE?

One of the possible reasons why GBE is not employed ubiquitously in industry is the limited flexibility on part geometry that it provides. Because of the large plastic strain required to trigger recrystallisation, GBE materials generally come in the form of sheets or tubes [ 29 ], as a result of the mechanical processes chosen to yield uniform and controlled deformation (e.g., rolling [ 30 ], drawing [ 31 ] and equal-channel angular pressing [ 32 ]). Thereafter, these materials require additional machining or forming to be shaped into a final product. When combined with the thermomechanical treatments required for GBE, the entire manufacturing process becomes time- and cost-intensive. Moreover, the range of parts that can be produced by sheet or wire forming is limited. For these reasons, GBE is not applied to bulk, three-dimensional (3-D) parts or components with intricate geometries. In these cases, the common practice is to rely on surface or coating technologies to minimize intergranular degradation, especially at high temperature and in corrosive environments [ 33 ].

Another drawback of GBE is the negative impact it has on materials strength, which is one of the main criteria when designing metals and metal alloys for structural applications. Since GBE relies on recovery and recrystallisation, the resulting polycrystals exhibit low densities of dislocations and low-angle GBs, as well as grains made larger by the heat treatment. Thus, the material loses both strain- and GB-hardening. Only rarely have researchers claimed an increase in material strength upon GBE. This trade-off between strength and other GB-controlled properties further restricts the application of GBE strategies to the surface of engineering components to avoid affecting the material’s bulk strength [ 34 , 35 ].

Revamping GBE through additive manufacturing

Additive manufacturing (AM) is regarded as a disruptive technology owing to its unique capability of producing bulk, near-net-shape parts by stacking layers of material into complex 3-D geometries. The unprecedented design freedom provided by AM offers many advantages over traditional manufacturing routes, including part count reduction [ 36 ], incorporation of intricate internal channels and chambers [ 37 , 38 ], and improved strength-to-weight ratio of structural components [ 39 , 40 ]. Beside these geometry-enabled advances in part design, AM opens many new opportunities for a microstructure-based design of materials [ 41 ]. The layerwise nature of the process, in fact, makes it possible to apply materials processing strategies—such as GBE—directly on individual layers as parts are produced (Fig.  2 ). The benefit of a layerwise GBE, which we refer to as additive-GBE (A-GBE), is that it would enable the direct production of bulk metal parts with both GBE microstructure and near-net-shape, topology-optimized geometry. As a result, A-GBE parts could be endowed with lightweight and enhanced resistance to intergranular degradation. This strategy could also be more energy- and cost-effective compared to conventional GBE, owing to the reduced temperature and mechanical deformation required to activate recrystallisation on each layer, as opposed to the entire part.

Conceptual schematic of “additive grain boundary engineering”

Some early studies have explored the possibility of processing materials using hybrid manufacturing approaches, which concurrently combine additive technologies with tooling to do mechanical work on the build [ 42 ]. Some notable examples include in-line rolling [ 43 , 44 ] and forging [ 45 ] to refine the microstructure in directed energy deposition (DED) processes, or in-situ laser or shot peening [ 46 , 47 ] to produce compressive stresses and raise the strain energy in materials produced by laser powder bed fusion (LPBF). The first strategy is restricted to AM parts that tolerate low dimensional accuracy, since the repeated deformation may change the build geometry substantially. The second overcomes this limitation but may only lead to partial recrystallisation of the material due to the relatively shallow depth of the deformation zone.

Another possible approach is to leverage the inherent strain energy formed during AM to activate microstructure recrystallisation, especially in additive processes that involve melting of the material feedstock, such as LPBF and DED. Because of the highly localized melting, steep thermal gradients, rapid cooling rates, and repeated thermal expansion and shrinkage cycles, materials produced by these processes exhibit highly non-equilibrium microstructures containing copious dislocation densities [ 48 , 49 , 50 ], deformation-induced defects [ 51 ], and large residual stress [ 52 , 53 ]. All these features raise the driving force for recrystallisation [ 54 , 55 ], which may be activated via post-production heat treatments. Indeed, parts produced by fusion-based AM are routinely heat treated to relieve residual stresses and to homogenize the microstructure [ 56 , 57 ]. However, microstructure recrystallisation in most of these materials only occurs after exposing them to very high temperatures for long times [ 58 , 59 ]. The shortcoming of these extensive heat treatments is that they may coarsen the microstructure and even yield the formation of unwanted phases, which would impart below-average mechanical performance to the alloy.

While the abovementioned studies delineate a pathway towards A-GBE, we are still far from devising a systematic and robust A-GBE strategy. The major challenge is to produce the necessary energy required to activate recrystallisation without compromising the geometry of the build or introducing detrimental residual strains which could lead to part failure [ 60 , 61 ]. In the following sections, we discuss alternative routes that could lead to A-GBE as well as some intriguing applications of it. We believe that both aspects will be the focus of intense research in this field in the near future.

An outlook on A-GBE in fusion-based AM

It should be noted that the density of coherent TBs and other twin-related GBs—hereafter generally referred to as TBs—in parts produced by AM is very limited in general [ 48 , 51 , 62 ]. Thus, A-GBE must rely on recrystallisation. As mentioned in the foregoing sections, it may be possible to activate recrystallisation and promote the formation of copious TBs by tailoring the non-equilibrium microstructure imparted by AM (and more specifically fusion-based AM). As many other phenomena that underpin the formation of materials and the evolution of their microstructure, recrystallisation requires a driving force and heat to overcome an energy barrier. In other words, the propensity of a material to undergo recrystallisation depends on how much strain energy is stored in the microstructure and how easy it is for new, recrystallized grains to nucleate and grow. In the quest for A-GBE, both aspects may be tuned concurrently through careful selection of AM processing parameters and/or by integrating layerwise mechanical treatments during AM.

Deformation-free A-GBE

Much of the strain energy required for recrystallisation is inherently generated during fusion-based AM. Indeed, the density of geometrically necessary dislocations found in metal alloys produced by LPBF ranges between 10 13  m −2 and 10 14  m −2 [ 50 , 63 ]. In theory, this residual strain should be sufficient to activate recrystallization at temperatures compatible with industrial standards [ 30 , 64 ] without any additional mechanical treatment. In practice, however, most of these AM alloys are thermally stable up to much higher temperatures [ 58 , 65 ]. This thermal stability stems from the presence of a fine solidification structure, which includes pronounced micro-segregation of solute atoms at cell or dendrite boundaries as a result of constitutional undercooling at the solidification front [ 62 ]. This structure hinders the onset and progression of recrystallisation despite the large driving force contained in the microstructure. In a recent work, we have shown how “weakening” this solidification structure by employing AM processing parameters that limit micro-segregation allows for recrystallisation to occur at progressively lower mechanically-induced strain [ 62 ].

Another microstructural feature that hinders recrystallization of alloys produced by fusion-based AM is second phase precipitates, such as oxide nanoparticles [ 66 ]. These particles are thought to originate from the melting of oxidized contaminants contained in the powder feedstock [ 67 , 68 ]. Due to the rapid solidification and high cooling rates the material undergoes, these solutionized impurities precipitate and form nano-scale particles. As in the case of micro-segregation, these particles pin GB motion during recrystallisation. By reducing the oxygen contamination level during the AM process (both in the processing chamber environment as well as in the powder feedstock), or by controlling the material’s cooling rate, it should be possible to limit the presence of these nanoparticles or reduce their size substantially; to a point where they would not refrain the growth of recrystallized grains.

While deformation-free A-GBE has yet to be demonstrated, we believe that devising strategies that simultaneously minimize GB pinning while raising the driving force for recrystallisation may prove successful. The latter could be achieved by employing unconventional laser sources [ 69 ], or laser processing methodologies [ 46 ] that promote higher residual strain in the as-built microstructure. Whatever the approach, a challenge will be to make such strategies scalable. Residual strains could add up and yield failure during production of large-scale parts, such as cracking, delamination, or distortion. Moreover, in these cases it may be more difficult to control the material’s thermal history and thus the solidification structure.

Heat treatment-free A-GBE

Another interesting feature that may facilitate A-GBE in fusion-based AM processes is the intrinsic heat treatment resulting from the repeated melting and solidification of individual layers. As the high-energy source (either a laser or an electron beam) scans the layer, it generates a heat affected zone that starts from the fusion boundary and extends into the solid material surrounding the melt pool [ 70 ]. By selecting different processing parameters, the heat affected zone may be tuned to positively affect the microstructure of the solidified material; for instance by triggering phase transformations [ 71 ], or activating recrystallisation [ 51 , 72 ]. Recently, Laleh et al . [ 73 ] found high fractions of TBs in the as-built microstructure of stainless steel and attributed this unusual phenomenon to poor heat dissipation during LPBF. Although their GB character distribution is dominated by high-mobility incoherent TBs, their work showcases the possibility to capitalize on the cyclic intrinsic heat treatments to activate dynamic recrystallization or recovery during the AM process. The advantage of this approach is that the parts produced would not require a GBE-specific heat treatment to activate recrystallisation, which would decrease production time and cost. Moreover, dynamic recrystallisation could also mitigate long standing problems related to the large residual stresses found in as-built AM parts [ 61 ].

Site-specific A-GBE

Because material and geometry are formed concurrently, point by point, during AM, parts may be produced with dissimilar microstructures using processing parameters that vary site-specifically. When controlled, this microstructure heterogeneity may impart additional functionalities to the build and have positive effects over parts performance. Some notable examples of this strategy can be found in the realm of surface engineering [ 74 ] or thin films technology [ 75 ], where such a heterogeneity can produce additional strengthening mechanisms and even help overcome the strength-ductility trade-off in metallic materials [ 76 , 77 , 78 ]. These “microstructure architectures”, however, are typically restricted to small scale materials because of the limitations associated with the respective manufacturing processes. With AM, these microstructure designs may be extended to bulk materials containing site-specific textures [ 41 ], directional solidification structures [ 79 ], dissimilar grain structures [ 80 ], composition gradients [ 81 ], and multiple phases [ 71 ]. In the context of A-GBE, site-specific microstructure control could be used to engineer the density of nucleation sites for recrystallisation across the build—for instance by selectively weakening the solidification structure. A low nucleation density would lead to the growth of large twin-related grain clusters separated by a sparse and disconnected network of high-angle GBs [ 82 ]. These high TB-density microstructures could exhibit properties comparable to those of materials that undergo several strain-annealing cycles following conventional GBE processes.

By controlling these microstructural features site-specifically, A-GBE could also enable the production of materials that integrate completely different GB character distributions [ 72 ], which would be impossible to attain via conventional GBE routes. One possible approach to achieve this goal is to tune the thermal stability of metal alloys site-specifically to alternate between regions that undergo recrystallisation and regions that do not. Alternatively, mechanical work could be applied only on specific regions of the build during hybrid manufacturing processes [ 72 ]. One benefit these microstructures could bring is to overcome the trade-off between enhanced GB-controlled properties and material strength in GBE materials. By designing the optimum fraction of recrystallized (i.e., soft) and non-recrystallized (i.e., hard) microstructures as well as their spatial distribution, parts could be made with high corrosion resistance and high strength at locations that best suit the constraints imposed by the target applications. We believe that these designs could be of interest for applications that require engineering alloys to operate in harsh environments.

Beyond twin-related GBE

This viewpoint focuses on TB-related GBE. However, TB multiplication through recrystallisation is restricted to materials with low stacking fault energy. While many engineering alloys fall under this category, including nickel, iron, and titanium alloys, others such as aluminum alloys are excluded from it. However, there are other types of GBs which could improve the properties of polycrystalline solids. For instance, some recent studies pinpointed the beneficial effects of low-angle GBs on intergranular corrosion of aluminum alloys [ 83 , 84 , 85 ] and on the strength of stainless steel [ 86 , 87 ]. The possibility offered by AM to control the crystallographic texture and local crystallographic misorientation [ 41 ] in the build opens the path to tailoring the occurrence of different types of GBs to improve the properties of any material. Moreover, this capability would significantly expand the design space of A-GBE materials to include alloys with site-specific regions dominated by high- or low-angle GBs arbitrarily distributed across the build.

Besides controlling the character distribution of GBs, AM may be pivotal to engineer their chemical composition, which provides an additional route to enhancing GB-governed properties of polycrystals. Raabe et al. [ 88 ] demonstrated that solute segregation at GBs may improve boundary cohesion, lower the boundary energy, and even promote local phase transformations. Manipulation of GB segregation during AM has been shown effective at mitigating hot cracking in nickel-based superalloy [ 89 ] and high entropy alloys [ 90 ]. These strategies involve, for instance, designing novel AM alloys that contain solute elements with low solute solubility and high strengthening power [ 91 ], or adjusting the AM process parameters to manipulate the cooling rate and thus tailor the GB segregation level [ 92 , 93 ].

For now, the materials that may be produced via these unconventional processing routes may not have obvious applications or functionalities that can be easily envisaged. However, it is only a matter of time before researchers in academia and industry start considering how to capitalize on these untapped opportunities to address the problems of tomorrow.

Watanabe T (2011) Grain boundary engineering: historical perspective and future prospects. J Mater Sci 46(12):4095–4115. https://doi.org/10.1007/s10853-011-5393-z

Article   CAS   Google Scholar  

Johnson OK, Li L, Demkowicz MJ, Schuh CA (2015) Inferring grain boundary structure–property relations from effective property measurements. J Mater Sci 50(21):6907–6919. https://doi.org/10.1007/s10853-015-9241-4

Lehockey EM, Palumbo G, Aust KT, Erb U, Lin P (1998) On the role of intercrystalline defects in polycrystal plasticity. Scripta Mater 39(3):341–346. https://doi.org/10.1016/S1359-6462(98)00173-0

Gao Y, Ritchie RO, Kumar M, Nalla RK (2005) High-cycle fatigue of nickel-based superalloy ME3 at ambient and elevated temperatures: Role of grain-boundary engineering. Metall Mater Trans A 36(12):3325–3333. https://doi.org/10.1007/s11661-005-0007-5

Article   Google Scholar  

Lehockey EM, Palumbo G (1997) On the creep behaviour of grain boundary engineered nickel 1. Mater Sci Eng A 237(2):168–172. https://doi.org/10.1016/S0921-5093(97)00126-3

Bechtle S, Kumar M, Somerday BP, Launey ME, Ritchie RO (2009) Grain-boundary engineering markedly reduces susceptibility to intergranular hydrogen embrittlement in metallic materials. Acta Mater 57(14):4148–4157. https://doi.org/10.1016/j.actamat.2009.05.012

Seita M, Hanson JP, Gradecak S, Demkowicz MJ (2015) The dual role of coherent twin boundaries in hydrogen embrittlement. Nat Commun 6:6164. https://doi.org/10.1038/ncomms7164

Shimada M, Kokawa H, Wang ZJ, Sato YS, Karibe I (2002) Optimization of grain boundary character distribution for intergranular corrosion resistant 304 stainless steel by twin-induced grain boundary engineering. Acta Mater 50(9):2331–2341. https://doi.org/10.1016/S1359-6454(02)00064-2

Randle V, Owen G (2006) Mechanisms of grain boundary engineering. Acta Mater 54(7):1777–1783. https://doi.org/10.1016/j.actamat.2005.11.046

Kumar M, King WE, Schwartz AJ (2000) Modifications to the microstructural topology in f.c.c. materials through thermomechanical processing. Acta Materialia 48 (9):2081–2091. https://doi.org/10.1016/S1359-6454(00)00045-8

Wang W, Brisset F, Helbert AL, Solas D, Drouelle I, Mathon MH, Baudin T (2014) Influence of stored energy on twin formation during primary recrystallization. Mater Sci Eng A 589:112–118. https://doi.org/10.1016/j.msea.2013.09.071

Randle V (2004) Twinning-related grain boundary engineering. Acta Mater 52(14):4067–4081. https://doi.org/10.1016/j.actamat.2004.05.031

Randle V (1999) Mechanism of twinning-induced grain boundary engineering in low stacking-fault energy materials. Acta Mater 47(15):4187–4196. https://doi.org/10.1016/S1359-6454(99)00277-3

Randle V, Davies H (2002) Evolution of microstructure and properties in alpha-brass after iterative processing. Metall Mater Trans A 33(6):1853–1857. https://doi.org/10.1007/s11661-002-0193-3

Rohrer GS, Randle V, Kim C-S, Hu Y (2006) Changes in the five-parameter grain boundary character distribution in α-brass brought about by iterative thermomechanical processing. Acta Mater 54(17):4489–4502. https://doi.org/10.1016/j.actamat.2006.05.035

Carpenter HCH, Tamura S (1926) The formation of twinned metallic crystals. Proc R Soc Lond Ser A Contain Pap Math Phys Char 113(763):161–182. https://doi.org/10.1098/rspa.1926.0144

Mahajan S, Pande CS, Imam MA, Rath BB (1997) Formation of annealing twins in f.c.c. crystals. Acta Materialia 45 (6):2633–2638. https://doi.org/10.1016/S1359-6454(96)00336-9

Gleiter H (1969) The formation of annealing twins. Acta Metall 17(12):1421–1428. https://doi.org/10.1016/0001-6160(69)90004-2

Watanabe T (1984) An approach to grain boundary design for strong and ductile polycrystals; Duktile Polykristalle hoher Festigkeit durch Einbau geeigneter Korngrenzen. Res mech 11(1):47–84

CAS   Google Scholar  

Palumbo G, King PJ, Aust KT, Erb U, Lichtenberger PC (1991) Grain boundary design and control for intergranular stress-corrosion resistance. Scr Metall Mater 25(8):1775–1780. https://doi.org/10.1016/0956-716X(91)90303-I

Lin P, Palumbo G, Erb U, Aust KT (1995) Influence of grain boundary character distribution on sensitization and intergranular corrosion of alloy 600. Scr Metall Mater 33(9):1387–1392. https://doi.org/10.1016/0956-716X(95)00420-Z

Randle V (2006) ‘Special’ boundaries and grain boundary plane engineering. Scripta Mater 54(6):1011–1015. https://doi.org/10.1016/j.scriptamat.2005.11.050

Kumar M, Schwartz AJ, King WE (2002) Microstructural evolution during grain boundary engineering of low to medium stacking fault energy fcc materials. Acta Mater 50(10):2599–2612. https://doi.org/10.1016/S1359-6454(02)00090-3

Randle V, Coleman M (2009) A study of low-strain and medium-strain grain boundary engineering. Acta Mater 57(11):3410–3421. https://doi.org/10.1016/j.actamat.2009.04.002

Watanabe T, Tsurekawa S (2004) Toughening of brittle materials by grain boundary engineering. Mater Sci Eng, A 387–389:447–455. https://doi.org/10.1016/j.msea.2004.01.140

Krell A, Pippel E, Woltersdorf J, Burger W (2003) Subcritical crack growth in Al2O3 with submicron grain size. J Eur Ceram Soc 23(1):81–89. https://doi.org/10.1016/S0955-2219(02)00072-9

Tan L, Sridharan K, Allen TR (2007) Effect of thermomechanical processing on grain boundary character distribution of a Ni-based superalloy. J Nucl Mater 371(1):171–175. https://doi.org/10.1016/j.jnucmat.2007.05.002

Gabb TP, Telesman J, Garg A, Lin P, Provenzano V, Heard R, Miller HM (2010) Grain Boundary Engineering the Mechanical Properties of Allvac 718Plus™ Superalloy. 7th International Symposium on Superalloy 718 and Derivatives 2010 1:255–269. doi: https://doi.org/10.1002/9781118495223.ch19

Zelinski JA (2005) An evaluation of grain boundary engineering technology and processing scale-up (Master's thesis). Master's thesis, Massachusetts Institute of Technology,

Jones R, Randle V (2010) Sensitisation behaviour of grain boundary engineered austenitic stainless steel. Mater Sci Eng, A 527(16–17):4275–4280. https://doi.org/10.1016/j.msea.2010.03.058

Xia S, Li H, Liu TG, Zhou BX (2011) Appling grain boundary engineering to Alloy 690 tube for enhancing intergranular corrosion resistance. J Nucl Mater 416(3):303–310. https://doi.org/10.1016/j.jnucmat.2011.06.017

Furukawa M, Horita Z, Langdon TG (2005) Processing by equal-channel angular pressing: Applications to grain boundary engineering. J Mater Sci 40(4):909–917. https://doi.org/10.1007/s10853-005-6509-0

Palumbo G, Lehockey EM, Lin P (1998) Applications for grain boundary engineered materials. JOM 50(2):40–43. https://doi.org/10.1007/s11837-998-0248-z

Yang S, Wang ZJ, Kokawa H, Sato YS (2007) Grain boundary engineering of 304 austenitic stainless steel by laser surface melting and annealing. J Mater Sci 42(3):847–853. https://doi.org/10.1007/s10853-006-0063-2

Telang A, Gill AS, Tammana D, Wen X, Kumar M, Teysseyre S, Mannava SR, Qian D, Vasudevan VK (2015) Surface grain boundary engineering of Alloy 600 for improved resistance to stress corrosion cracking. Mater Sci Eng, A 648:280–288. https://doi.org/10.1016/j.msea.2015.09.074

Atzeni E, Salmi A (2012) Economics of additive manufacturing for end-usable metal parts. Int J Adv Manuf Technol 62(9–12):1147–1155. https://doi.org/10.1007/s00170-011-3878-1

Zhang L, Hu Z, Wang MY, Feih S (2021) Hierarchical sheet triply periodic minimal surface lattices: Design, geometric and mechanical performance. Mater Des 209:109931. https://doi.org/10.1016/j.matdes.2021.109931

Li X, Yu X, Zhai W (2021) Additively Manufactured Deformation-Recoverable and Broadband Sound-Absorbing Microlattice Inspired by the Concept of Traditional Perforated Panels. Adv Mater 2104552. https://doi.org/10.1002/adma.202104552

Sercombe TB, Xu X, Challis VJ, Green R, Yue S, Zhang Z, Lee PD (2015) Failure modes in high strength and stiffness to weight scaffolds produced by Selective Laser Melting. Mater Des 67:501–508. https://doi.org/10.1016/j.matdes.2014.10.063

Xu Z, Ha CS, Kadam R, Lindahl J, Kim S, Wu HF, Kunc V, Zheng X (2020) Additive manufacturing of two-phase lightweight, stiff and high damping carbon fiber reinforced polymer microlattices. Addit Manuf 32:101106. https://doi.org/10.1016/j.addma.2020.101106

Sofinowski KA, Raman S, Wang X, Gaskey B, Seita M (2021) Layer-wise engineering of grain orientation (LEGO) in laser powder bed fusion of stainless steel 316L. Addit Manuf 38:101809. https://doi.org/10.1016/j.addma.2020.101809

Sealy MP, Madireddy G, Williams RE, Rao P, Toursangsaraki M (2018) Hybrid Processes in Additive Manufacturing. J Manuf Sci Eng 140(6):060801. https://doi.org/10.1115/1.4038644

Colegrove PA, Donoghue J, Martina F, Gu J, Prangnell P, Hönnige J (2017) Application of bulk deformation methods for microstructural and material property improvement and residual stress and distortion control in additively manufactured components. Scripta Mater 135:111–118. https://doi.org/10.1016/j.scriptamat.2016.10.031

Xu X, Ganguly S, Ding J, Seow CE, Williams S (2018) Enhancing mechanical properties of wire + arc additively manufactured INCONEL 718 superalloy through in-process thermomechanical processing. Mater Des 160:1042–1051. https://doi.org/10.1016/j.matdes.2018.10.038

Li Q, Zhang Y, Chen J, Guo B, Wang W, Jing Y, Liu Y (2021) Effect of ultrasonic micro-forging treatment on microstructure and mechanical properties of GH3039 superalloy processed by directed energy deposition. J Mater Sci Technol 70:185–196. https://doi.org/10.1016/j.jmst.2020.09.001

Kalentics N, Boillat E, Peyre P, Gorny C, Kenel C, Leinenbach C, Jhabvala J, Logé RE (2017) 3D Laser Shock Peening – A new method for the 3D control of residual stresses in Selective Laser Melting. Mater Des 130:350–356. https://doi.org/10.1016/j.matdes.2017.05.083

Book TA, Sangid MD (2016) Evaluation of Select Surface Processing Techniques for In Situ Application During the Additive Manufacturing Build Process. Jom 68(7):1780–1792. https://doi.org/10.1007/s11837-016-1897-y

Wang YM, Voisin T, McKeown JT, Ye J, Calta NP, Li Z, Zeng Z, Zhang Y, Chen W, Roehling TT, Ott RT, Santala MK, Depond PJ, Matthews MJ, Hamza AV, Zhu T (2018) Additively manufactured hierarchical stainless steels with high strength and ductility. Nat Mater 17(1):63–71. https://doi.org/10.1038/nmat5021

Zhang D, Niu W, Cao X, Liu Z (2015) Effect of standard heat treatment on the microstructure and mechanical properties of selective laser melting manufactured Inconel 718 superalloy. Mater Sci Eng, A 644:32–40. https://doi.org/10.1016/j.msea.2015.06.021

Hu Z, Zhang H, Zhu H, Xiao Z, Nie X, Zeng X (2019) Microstructure, mechanical properties and strengthening mechanisms of AlCu5MnCdVA aluminum alloy fabricated by selective laser melting. Mater Sci Eng, A 759:154–166. https://doi.org/10.1016/j.msea.2019.04.114

Sabzi HE, Li X-H, Zhang C, Fu H, San-Martín D, Rivera-Díaz-del-Castillo PEJ (2022) Deformation twinning-induced dynamic recrystallization during laser powder bed fusion. Scripta Mater 207:114307. https://doi.org/10.1016/j.scriptamat.2021.114307

Levkulich NC, Semiatin SL, Gockel JE, Middendorf JR, DeWald AT, Klingbeil NW (2019) The effect of process parameters on residual stress evolution and distortion in the laser powder bed fusion of Ti-6Al-4V. Addit Manuf 28:475–484. https://doi.org/10.1016/j.addma.2019.05.015

Sabzi HE, Aboulkhair NT, Liang X, Li X-H, Simonelli M, Fu H, Rivera-Díaz-del-Castillo PEJ (2020) Grain refinement in laser powder bed fusion: The influence of dynamic recrystallization and recovery. Mater Des 196:109181. https://doi.org/10.1016/j.matdes.2020.109181

Holland S, Wang X, Fang XY, Guo YB, Yan F, Li L (2018) Grain boundary network evolution in Inconel 718 from selective laser melting to heat treatment. Mater Sci Eng, A 725:406–418. https://doi.org/10.1016/j.msea.2018.04.045

Song B, Dong S, Liu Q, Liao H, Coddet C (2014) Vacuum heat treatment of iron parts produced by selective laser melting: Microstructure, residual stress and tensile behavior. Mater Des 1980–2015(54):727–733. https://doi.org/10.1016/j.matdes.2013.08.085

Yadollahi A, Shamsaei N, Thompson SM, Seely DW (2015) Effects of process time interval and heat treatment on the mechanical and microstructural properties of direct laser deposited 316L stainless steel. Mater Sci Eng, A 644:171–183. https://doi.org/10.1016/j.msea.2015.07.056

Sangid MD, Book TA, Naragani D, Rotella J, Ravi P, Finch A, Kenesei P, Park J-S, Sharma H, Almer J, Xiao X (2018) Role of heat treatment and build orientation in the microstructure sensitive deformation characteristics of IN718 produced via SLM additive manufacturing. Addit Manuf 22:479–496. https://doi.org/10.1016/j.addma.2018.04.032

Chen X, Li J, Cheng X, Wang H, Huang Z (2018) Effect of heat treatment on microstructure, mechanical and corrosion properties of austenitic stainless steel 316L using arc additive manufacturing. Mater Sci Eng, A 715:307–314. https://doi.org/10.1016/j.msea.2017.10.002

Kong D, Dong C, Ni X, Zhang L, Yao J, Man C, Cheng X, Xiao K, Li X (2019) Mechanical properties and corrosion behavior of selective laser melted 316L stainless steel after different heat treatment processes. J Mater Sci Technol 35(7):1499–1507. https://doi.org/10.1016/j.jmst.2019.03.003

DebRoy T, Wei HL, Zuback JS, Mukherjee T, Elmer JW, Milewski JO, Beese AM, Wilson-Heid A, De A, Zhang W (2018) Additive manufacturing of metallic components – Process, structure and properties. Prog Mater Sci 92:112–224. https://doi.org/10.1016/j.pmatsci.2017.10.001

Li C, Liu ZY, Fang XY, Guo YB (2018) Residual Stress in Metal Additive Manufacturing. Procedia CIRP 71:348–353. https://doi.org/10.1016/j.procir.2018.05.039

Gao S, Hu Z, Duchamp M, Krishnan PSSR, Tekumalla S, Song X, Seita M (2020) Recrystallization-based grain boundary engineering of 316L stainless steel produced via selective laser melting. Acta Mater 200:366–377. https://doi.org/10.1016/j.actamat.2020.09.015

Song B, Dong S, Deng S, Liao H, Coddet C (2014) Microstructure and tensile properties of iron parts fabricated by selective laser melting. Opt Laser Technol 56:451–460. https://doi.org/10.1016/j.optlastec.2013.09.017

Parvathavarthini N, Dayal RK, Khatak HS, Shankar V, Shanmugam V (2006) Sensitization behaviour of modified 316N and 316L stainless steel weld metals after complex annealing and stress relieving cycles. J Nucl Mater 355(1):68–82. https://doi.org/10.1016/j.jnucmat.2006.04.006

Kouraytem N, Varga J, Amin-Ahmadi B, Mirmohammad H, Chanut RA, Spear AD, Kingstedt OT (2021) A recrystallization heat-treatment to reduce deformation anisotropy of additively manufactured Inconel 718. Mater Des 198:109228. https://doi.org/10.1016/j.matdes.2020.109228

Aota LS, Bajaj P, Zilnyk KD, Jägle EA, Ponge D, Sandim HRZ, Raabe D (2021) Recrystallization kinetics, mechanisms, and topology in alloys processed by laser powder-bed fusion: AISI 316L stainless steel as example. Materialia 20:101236. https://doi.org/10.1016/j.mtla.2021.101236

Saeidi K, Gao X, Zhong Y, Shen ZJ (2015) Hardened austenite steel with columnar sub-grain structure formed by laser melting. Mater Sci Eng, A 625:221–229. https://doi.org/10.1016/j.msea.2014.12.018

Saeidi K, Kvetková L, Lofaj F, Shen Z (2015) Austenitic stainless steel strengthened by the in situ formation of oxide nanoinclusions. RSC Adv 5(27):20747–20750. https://doi.org/10.1039/C4RA16721J

Nassar AR, Reutzel EW (2015) Additive Manufacturing of Ti-6Al-4V Using a Pulsed Laser Beam. Metall and Mater Trans A 46(6):2781–2789. https://doi.org/10.1007/s11661-015-2838-z

Wang T, Zhu YY, Zhang SQ, Tang HB, Wang HM (2015) Grain morphology evolution behavior of titanium alloy components during laser melting deposition additive manufacturing. J Alloy Compd 632:505–513. https://doi.org/10.1016/j.jallcom.2015.01.256

Kurnsteiner P, Wilms MB, Weisheit A, Gault B, Jagle EA, Raabe D (2020) High-strength Damascus steel by additive manufacturing. Nature 582(7813):515–519. https://doi.org/10.1038/s41586-020-2409-3

Gao S, Liu R, Huang R, Song X, Seita M (2022) A hybrid directed energy deposition process to manipulate microstructure and properties of austenitic stainless steel. Mater Des 213:110360. https://doi.org/10.1016/j.matdes.2021.110360

Laleh M, Hughes AE, Tan MY, Rohrer GS, Primig S, Haghdadi N (2021) Grain boundary character distribution in an additively manufactured austenitic stainless steel. Scripta Mater 192:115–119. https://doi.org/10.1016/j.scriptamat.2020.10.018

Liu XC, Zhang HW, Lu K (2013) Strain-Induced Ultrahard and Ultrastable Nanolaminated Structure in Nickel. Science 342(6156):337–340. https://doi.org/10.1126/science.1242578

Wu X, Yang M, Yuan F, Wu G, Wei Y, Huang X, Zhu Y (2015) Heterogeneous lamella structure unites ultrafine-grain strength with coarse-grain ductility. Proc Natl Acad Sci 112(47):14501. https://doi.org/10.1073/pnas.1517193112

Zhu Y, Ameyama K, Anderson PM, Beyerlein IJ, Gao H, Kim HS, Lavernia E, Mathaudhu S, Mughrabi H, Ritchie RO, Tsuji N, Zhang X, Wu X (2021) Heterostructured materials: superior properties from hetero-zone interaction. Mater Res Lett 9(1):1–31. https://doi.org/10.1080/21663831.2020.1796836

Zhu Y, Wu X (2019) Perspective on hetero-deformation induced (HDI) hardening and back stress. Mater Res Lett 7(10):393–398. https://doi.org/10.1080/21663831.2019.1616331

Lu K (2014) Making strong nanomaterials ductile with gradients. Science 345(6203):1455. https://doi.org/10.1126/science.1255940

Tekumalla S, Selvarajou B, Raman S, Gao S, Seita M (2021) The role of the solidification structure on orientation-dependent hardness in stainless steel 316L produced by laser powder bed fusion. Mater Sci Eng A 833:142493. https://doi.org/10.1016/j.msea.2021.142493

Todaro CJ, Easton MA, Qiu D, Zhang D, Bermingham MJ, Lui EW, Brandt M, StJohn DH, Qian M (2020) Grain structure control during metal 3D printing by high-intensity ultrasound. Nat Commun 11(1):142. https://doi.org/10.1038/s41467-019-13874-z

Yeoh YC, Macchi G, Jain E, Gaskey B, Raman S, Tay G, Verdi D, Patran A, Grande AM, Seita M (2021) Multiscale microstructural heterogeneity and mechanical property scatter in Inconel 718 produced by directed energy deposition. J Alloy Compd 887:161426. https://doi.org/10.1016/j.jallcom.2021.161426

Barr CM, Leff AC, Demott RW, Doherty RD, Taheri ML (2018) Unraveling the origin of twin related domains and grain boundary evolution during grain boundary engineering. Acta Mater 144:281–291. https://doi.org/10.1016/j.actamat.2017.10.007

Yan J, Heckman NM, Velasco L, Hodge AM (2016) Improve sensitization and corrosion resistance of an Al-Mg alloy by optimization of grain boundaries. Sci Rep 6(1):26870. https://doi.org/10.1038/srep26870

Orłowska M, Ura-Bińczyk E, Olejnik L, Lewandowska M (2019) The effect of grain size and grain boundary misorientation on the corrosion resistance of commercially pure aluminium. Corros Sci 148:57–70. https://doi.org/10.1016/j.corsci.2018.11.035

Zhang X, Lv Y, Hashimoto T, Nilsson J-O, Zhou X (2021) Intergranular corrosion of AA6082 Al–Mg–Si alloy extrusion: The influence of trace Cu and grain boundary misorientation. J Alloy Compd 853:157228. https://doi.org/10.1016/j.jallcom.2020.157228

Sabzi HE, Hernandez-Nava E, Li X-H, Fu H, San-Martín D, Rivera-Díaz-del-Castillo PEJ (2021) Strengthening control in laser powder bed fusion of austenitic stainless steels via grain boundary engineering. Mater Des 212:110246. https://doi.org/10.1016/j.matdes.2021.110246

Abdi F, Eftekharian A, Huang D, Rebak RB, Rahmane M, Sundararaghavan V, Kanyuck A, Gupta SK, Arul S, Jain V, Hu Y, Nikbin K (2021) Grain boundary engineering of new additive manufactured polycrystalline alloys. Forces Mech 4:100033. https://doi.org/10.1016/j.finmec.2021.100033

Raabe D, Herbig M, Sandlöbes S, Li Y, Tytko D, Kuzmina M, Ponge D, Choi PP (2014) Grain boundary segregation engineering in metallic alloys: A pathway to the design of interfaces. Curr Opin Solid State Mater Sci 18(4):253–261. https://doi.org/10.1016/j.cossms.2014.06.002

Kontis P, Chauvet E, Peng Z, He J, da Silva AK, Raabe D, Tassin C, Blandin J-J, Abed S, Dendievel R, Gault B, Martin G (2019) Atomic-scale grain boundary engineering to overcome hot-cracking in additively-manufactured superalloys. Acta Mater 177:209–221. https://doi.org/10.1016/j.actamat.2019.07.041

Sun Z, Tan X, Wang C, Descoins M, Mangelinck D, Tor SB, Jägle EA, Zaefferer S, Raabe D (2021) Reducing hot tearing by grain boundary segregation engineering in additive manufacturing: example of an AlxCoCrFeNi high-entropy alloy. Acta Mater 204:116505. https://doi.org/10.1016/j.actamat.2020.116505

Thapliyal S, Agrawal P, Agrawal P, Nene SS, Mishra RS, McWilliams BA, Cho KC (2021) Segregation engineering of grain boundaries of a metastable Fe-Mn-Co-Cr-Si high entropy alloy with laser-powder bed fusion additive manufacturing. Acta Mater 219:117271. https://doi.org/10.1016/j.actamat.2021.117271

Shigesato G, Fujishiro T, Hara T (2014) Grain Boundary Segregation Behavior of Boron in Low-Alloy Steel. Metall and Mater Trans A 45(4):1876–1882. https://doi.org/10.1007/s11661-013-2155-3

He XL, Chu YY, Jonas JJ (1989) Grain boundary segregation of boron during continuous cooling. Acta Metall 37(1):147–161. https://doi.org/10.1016/0001-6160(89)90274-5

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Seita, M., Gao, S. Broadening the design space of engineering materials through “additive grain boundary engineering”. J Mater Sci 57 , 9530–9540 (2022). https://doi.org/10.1007/s10853-022-06886-6

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Journal of Materials Chemistry A

Configuration-entropy effects on bifeo 3 –batio 3 relaxor ferroelectric ceramics for high-density energy storage †.

* Corresponding authors

a International PhD Program in Innovative Technology of Biomedical Engineering and Medical Devices, Ming Chi University of Technology, New Taipei City 24301, Taiwan E-mail: [email protected]

b Department of Physics, Silliman University, Dumaguete City, Philippines

c Department of Mechanical Engineering, Hwa Hsia University of Technology, New Taipei City 23567, Taiwan E-mail: [email protected]

d Department of Mechanical Engineering, Ming Chi University of Technology, New Taipei City 24301, Taiwan

e Research Center for Intelligent Medical Devices, Ming Chi University of Technology, New Taipei City 24301, Taiwan

f Department of Physics, Fu Jen Catholic University, New Taipei City 24205, Taiwan E-mail: [email protected]

High energy-storage capability and electric breakdown strength are critical elements in next-generation pulse-power dielectric capacitors. In this report, perovskite (Bi 0.7 Ba 0.3 ) 1− x Na x (Fe 0.7 Ti 0.3 ) 1− x Ta x O 3 relaxor ferroelectric ceramics ( x = 0–0.3) were tailored in terms of configuration entropy from a medium entropy of 1.21 R to a high entropy of 2.07 R to improve energy storage. The integration of paraelectric NaTaO 3 into BiFeO 3 –BaTiO 3 results in breaking of the long-range order and formation of multiple lattice distortions toward relaxor ferroelectric characteristics. Excellent recoverable energy densities of 9.6 J cm −3 and 10.3 J cm −3 with efficiencies of 77% and 68% at 350 kV cm −1 and 550 kV cm −1 (at 10 Hz) were achieved for x = 0.15 and 0.20, respectively. Wide operating frequency (1–100 Hz) and temperature (25 °C–150 °C) stabilities were confirmed at 300 kV cm −1 . Grain boundaries and nanoclusters play critical roles as electric barriers to suppress charge mobility and increase electric breakdown strength. This study presents a promising scheme to utilize high-configuration entropy BiFeO 3 –BaTiO 3 -based ceramics for high energy-density electrostatic capacitors.

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Configuration-entropy effects on BiFeO 3 –BaTiO 3 relaxor ferroelectric ceramics for high-density energy storage

R. Montecillo, C. Chen, K. Feng, R. R. Chien, P. Chen and C. Tu, J. Mater. Chem. A , 2024, Advance Article , DOI: 10.1039/D4TA00921E

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