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. 2020 Dec 27;7(2):318–326. doi: 10.1021/acscentsci.0c01365

Laser Cladding Induced Spherical Graphitic Phases by Super-Assembly of Graphene-Like Microstructures and the Antifriction Behavior

Fei Weng †,‡, Cheng Hu §, Runhao Zhang ∥, Huijun Yu †,‡,*, Jiaqing Liu ∥, Meng Wang ∥, Yong Li ∥, Lei Xie ⊥, Chuanzhong Chen §,*, Kang Liang #, Dongyuan Zhao ⊥, Biao Kong ⊥,*
PMCID: PMC7908040  PMID: 33655069

Abstract

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Laser cladding coatings with excellent wear resistance behaviors are prepared on a titanium alloy substrate with a new precursor material system comprising nanoscale B4C and Ni60A self-fluxing alloy powder. Structural analysis reveals the existence of micron-size spherical or nearly spherical graphitic phases in the prepared coatings, which are composed of graphene-like microstructures closely associated with other reinforcement phases of high hardness such as TiC and CrB. The formation mechanism of these graphitic phases involves in situ superassembly of uncombined C atoms via repeated growth and reorientation of the graphene-like microstructures and is closely related to the laser processing parameters as well as the precursor compositions. The coexistence of these heterogeneous phases enable the obtained coatings with high wear resistance and low friction coefficient. It was found that the wear resistance of the coating has a remarkable 43.67 times enhancement than that of the titanium alloy while simultaneously showing a low friction coefficient (∼0.35). The understanding of the formation mechanism on the graphene-related novel microstructures with significantly improved mechanical properties is expected to lay the foundation for future developments and applications of graphene and its related carbon materials, such as large-scale production and further incorporation into composite materials with desired local structures.

Short abstract

Surplus carbon atoms in the molten pool nucleate at the bubble/melt interface and grow into graphitic layers. The spherical graphic structure is formed by superassembly of these graphitic layers.

Introduction

Graphene and its related carbon materials have been proved to play significant roles in various applications ranging from energy storage and conversion to bioimaging and biosensors.1−13 Strategies to synthesize graphene-related structures have been developed to facilitate their use in many applications,14 while their large-scale production and further incorporation into composite materials with desired local structures still remain a challenge. A laser beam is an excellent heat source used in a wide range of materials processing operations, which enables the in situ formation of novel micro- and nanostructures toward a diverse range of application requirements.15−17 In this respect, much attention has been devoted to the surface modification of alloys, and techniques such as pulsed laser deposition,18 laser shock peening,19,20 laser surface melting,21,22 and laser cladding23,24 have been studied. Among these, laser cladding is the most popular surface modification technique, providing improved processing efficiency, smaller heat-affected zones, stronger metallurgical bonding to the substrate, and controllable thickness of the resulting coatings.

In the laser cladding process, precursor materials play important roles in tailoring the final laser cladding coating properties.25,26 Carbides such as B4C27 is a class of widely used laser cladding precursors for titanium alloy surface modification. B4C supplies B and C atoms, which react with Ti to form hard and wear-resistant phases (e.g., TiB2, TiC). Graphite has also been used as a laser cladding precursor for titanium alloy surface modification toward friction reduction. However, related studies suggest that graphite is hard to be restrained after laser processing owing to the high affinity of C with Ti under elevated temperatures.28 To date, in situ synthesized graphitic (sp2 hybridized) phases from laser cladding on titanium alloys, as well as their detailed structure and impact on the coating properties are rarely reported. Recent advancements see the use of nanoscale CeO2,29 Al2O3,30 TiO2,31 SiC,32 and WC33 as laser cladding precursors with significantly improved tribological properties obtained. Moreover, the use of nanoscale carbide precursors combined with laser processing also provide the possibility for the in situ formation of graphitic phases with graphene-like structures, which in turn would lead to additional enhancement in mechanical properties.

In this study, spherical or nearly spherical graphitic phases comprised of graphene-like few-layer microstructures are reported for the first time in laser cladding coatings on titanium alloy with precursors containing nanoscale B4C particles. Using the newly designed material system with proper process parameters, composite laser cladding coatings with excellent tribological properties are obtained. The wear resistance of the coatings has a remarkable 43.67 times improvement than that of the bare titanium alloy substrate while the coating friction coefficient is as low as ∼0.35. The laser cladding induced spherical graphite phases by superassembly of the graphene-like few-layer microstructures and their influence on the coating properties are discussed. The composite coatings reported in this study can be used in surface strengthening of titanium alloy, further broadening their application into harsh friction and wear occasions. Though only a single layer coating was fabricated, it can be broadened into multiple layers (3D printing). The significantly improved mechanical properties as well as the understanding on the formation mechanism of the graphene-related novel microstructures are expected to lay the foundation for future developments of graphene and its related carbon materials, as well as their large-scale production.

Results and Discussion

Coating Preparation and Interfacial Morphologies

In a typical laser cladding treatment process, B4C nanoscale particles were first mixed with Ni60A Ni-based self-fluxing alloy powder to form the precursor materials and preplaced on Ti-6Al-4V titanium alloy substrates. A fine laser beam with controlled power and spot size was used to scan across the sample area at a preset scanning velocity. Upon laser irradiation, the decomposed precursors were mixed with the surface layer of the substrate in the molten pool where they react to form a composite coating that is metallurgically bonded to the substrate (Figure 1). For more experimental details, please refer to the Supporting Information. The bonding quality of the coating to the substrate is mainly influenced by the dilution rate (η), calculated via η = A1/(A1 + A2), where A1 is the melting area of the substrate and A2 is the cladding area. A total of 11 specimens are studied (Table S1), including nine specimens from a group of orthogonal experiments (L9, 33), where the laser power, scanning velocity, and nanoscale B4C fraction are the variables, as well as another two as supplementary to specimen 2 to investigate the effect of the laser power. The maximum content of nanoscale B4C in the precursor materials was 10.0 wt %, which is in consideration of the high specific surface area of the nanoscale particles.

Figure 1.

Figure 1

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Schematic illustration of the laser cladding process and coating formation mechanism. Precursors are first mixed and preplaced on the substrate. Laser beam is used as a heat source and scan across the specimen surface. Upon laser irradiation, precursors decompose while the surface layer of the substrate melts, forming the molten pool. The molten pool forms and solidifies dynamically as the laser beam scans, during which complicated reactions occur to form various reinforcement phases in the resulted coating.

Scanning electron microscopy (SEM) was employed to visualize the cross-section morphologies of the individual laser cladding tracks obtained under different process parameters and precursor material compositions (Figure 2). The results indicate that all the coatings are bonded well with the substrates with well-defined transition zones between the coating and the substrate. Taking Specimen 2 as an example, the transition zone (annotated by the dotted line, Figure S1a) is mainly composed of NiTi/NiTi2 dendrites (Figure S1b), derived from the reaction between Ni and Ti.34Figure 2J presents a schematic illustration of A1 and A2 for dilution rate calculation. To enable successful metallurgical bonding of the coating and the substrate, dilution rate should be carefully controlled by tuning laser power, spot size, and scanning velocity. Low laser specific energy typically leads to low dilution rates (Specimens 1, 2, 5, and 6 in Figure 2A, B, D, E respectively). In contrast, Specimen 7 exhibits the highest dilution rate (61.24%) owing to the high laser specific energy used. Pores are observed near the surface layer of the coating specimens shown in Figure 2B, D, and I. During the cladding process, gas can be formed due to the reaction/convection of the molten pool with ambient atmosphere. Under a lower laser power or a higher scanning velocity, the formed gas phase could not have enough time to exit from the molten pool, leading to formation of pores.

Figure 2.

Figure 2

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Cross-section morphologies of the laser cladding tracks in different specimens. Specimen number (No.), laser power (P, W), scan velocity (v, mm/min), laser specific energy (E, kJ·cm–2), fraction of nanoscale B4C (f, wt %) together with the calculated dilution rate (η) are provided as inset separately. The schematic illustration of the melting area in the substrate (A1) and the cladding area (A2) in the prepared coatings for dilution rate calculation is shown in (J). All the scale bars are 1.0 mm. The cross-section morphology of Specimen 3 is not shown here as no continuous coating were obtained with the cladding parameters used (see Figure S2).

Microstructures

Cross-sectional morphology of the prepared coating was examined by SEM. Carbon-rich zones of a few microns in spherical or nearly spherical morphologies exist in Specimens 2 and 5 (Figure 3A–D), while the same features are not seen in any other specimens. Both linear and point elemental analyses by energy-dispersive spectrometry (EDS) demonstrated the elevated amount of C elements in these areas (Figure S3 and Figure S4). Areas close to the spherical carbon-rich zones are found to be rich in Cr and B (Figure 3C and Figure S4), which therefore can be assigned to chromium boride (CrB). Considering the low friction coefficient of Specimens 2 and 5 (see Tribological Properties), these carbon-rich zones could consist of graphitic microstructures, which shows good self-lubricating characteristic.

Figure 3.

Figure 3

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Microstructure of carbon-rich zones under SEM and the graphene-like few-layer microstructure under HRTEM. (A–D) SEM morphologies showing the spherical or nearly spherical carbon-rich zones in the prepared coatings for Specimen 2 (A and B) and Specimen 5 (C and D) (B and D are magnified from areas in A and C, respectively, as annotated). Point elemental analysis results of P1, P2 and P3 are shown in Figure S3. (E–H) HRTEM images of the graphene-like few-layer microstructures in the carbon-rich zones of specimen 2 showing a graphene-like basal planes (E), spherical carbon-rich zones (F), (002) lattice image of graphene-like layers coexisting with an amorphous phase close to TiC (G), and (002) lattice image of the graphene-like layers coexisting with CrB (H). (Brown ribbons or spheres in the sketch models inset to the images stand for graphene-like structures.)

X-ray diffraction (XRD) was conducted on Specimen 2 to confirm the possible existence of any graphitic phase in the coatings in correlation with the SEM-observed carbon-rich zones. Though no graphitic phase can be identified in the XRD result (Figure S5), various phases in the coating were detected, including γ-Ni, TiB2, TiB, TiC, NiTi2, NiTi, Ni2B, Ni3B, Cr2B, CrB, Cr7C3, and Cr23C6. The composite phase constituent is mainly due to the various reactions between the Ni60A and nanoscale B4C precursor with the substrate elements in the molten pool, consistent with our previous results.35 The nanoscale B4C in the precursor materials should undergo sufficient transformation during the laser cladding, judging from the XRD analysis (Figure S5) together with the previous mentioned SEM and EDS results (Figure 3, Figures S3 and S4).

High-resolution transmission electron microscopy (HRTEM) was conducted to further study the microstructures, especially the carbon-rich zones observed by SEM. The HRTEM image in Figure 3E clearly shows the existence of the graphene-like few-layer microstructure. The visible curvatures of the observed graphene-like layers suggest a significant level of in-plane defects. During the sample preparation process, breakup of the robust coating structure surrounding the carbon-rich zones could lead to the release of local constraining. The graphene-like layers then flatten out as a result, which enables their identification under HRTEM at the right alignment with the electron beam. Moreover, the carbon-rich zones displayed a spherical morphology (Figure 3F), which is similar as observed by SEM (Figure 3B).

Close-up examination of the carbon-rich zone edges clearly reveal the lattice images of the graphene-like layers (indicated by arrows in Figure 3G, H). The interplanar spacing (marked as G 0002 in Figure 3G, H) is slightly larger than the standard (0002) value (∼0.34 nm) of graphite. This can be attributed to the aforementioned intrinsic defect-induced curvatures. Further examination suggests that the graphene-like layers are closely associated with CrB and amorphous phases (marked in Figure 3G, H). Nanoscale TiC phases are also found to exist close to the graphene-like layers, separated by the amorphous phase as shown in Figure 3G. The coexistence of such nanoscale amorphous and crystalline phases originates from the rapid cooling after laser processing.

Beyond the carbon-rich zones, γ-Ni and TiB2 phases are also identified by HRTEM (Figure S6a–d). Weak diffraction spots can be seen in the γ-Ni diffraction pattern (see Figure S6b for selected area electron diffraction), suggesting the coexistence of other microcrystals, possibly Cr23C6. The lattice image (Figure S6e) of the Cr23C6 structures is observed close to γ-Ni (indicated by arrow, Figure S6a). In addition, CrB and Cr3C2 are also observed (Figure S6f) by HRTEM, which both serve as reinforcement phases in the cladding coatings.

Tribological Properties

We next examined the tribological properties of the prepared specimens (Figure 4A–C). Specimen 2 shows the lowest wear mass loss and the lowest friction coefficient among the series of specimens. High wear mass loss was found on Specimen 4, 7, and 11, which can be attributed to the high dilution rate under high laser specific energy. Higher laser power (3.0 kW) leads to higher friction coefficients among the prepared specimens while the lowest friction coefficient is found on Specimen 2 prepared using 1.0 kW laser power. The average friction coefficient of Specimen 2 is 0.35 during the initial 20 min test and the value gradually increases to 0.41 at 30 min (Figure 4B). Besides, Specimen 5 also shows a relatively low wear mass loss and friction coefficient, but large pores are found in the coating (Figure 2D), which could induce local stress concentration. The existence of these large pores can be attributed to the inadequate flowability of the molten pool caused by the high B4C amount (10 wt %) that consumes a significantly higher amount of energy during melting.36 Though the laser specific energy is twice that in specimen 2, more B4C particles are added in specimen 5 (10 wt %) than that in specimen 2 (5 wt %). The decomposition of massive nanoscale B4C particles consumes relatively more energy in specimen 5, resulting in inadequate flowability of the molten pool and even large pores (Figure 2D).

Figure 4.

Figure 4

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Wear property, XRD analysis, and Ti–B–C phase diagram. (A) Wear mass loss of the substrate and the prepared coatings. (B, C) Friction coefficient curves of the prepared coatings. (D) XRD pattern of the nanoscale B4C precursor. (E) Calculated isothermal Ti–B–C ternary phase diagram at 1400 °C.37

Taking Specimen 2, 10, and 11 into consideration, the wear mass loss increases with laser power (from 1.0 kW to 3.0 kW), while the opposite is seen in the microhardness distributions (Figure S7). As previously mentioned, the dilution rate for Specimen 11 is as high as 49.04%, which indicates a massive amount of Ti diffuses into the molten pool from the substrate. A surplus of low hardness Ti weakens the intrinsic mechanical property of the entire Ni-based composite coating, resulting in a “dilution effect”. As a result, Specimen 2 fabricated at 1.0 kW shows the lowest dilution rate and wear mass loss, which is only 1/43.67 that of the Ti-6Al-4V substrate. Among the specimens in the present study, specimen 2 shows superior tribological performance. Comparison on the tribological properties between specimen 2 and some published reports is listed in Table S2, where specimen 2 also stands out with remarkable relative wear resistance and friction coefficient.

It is a great challenge for metal surface modification to achieve high wear resistance and low friction coefficient simultaneously, although they are the most ideal combination for practical applications. A low friction coefficient is beneficial to prevent coatings from delamination.38 Various factors could affect the friction coefficient, including the surface roughness and the inherent properties of the constitute phases that exist within the coatings.39 To reduce the friction coefficient, efforts have been spent by adding self-lubricating additives to the laser cladding precursors. Additives such as CaF2,40 WS2,41 MoS2,42 and h-BN43 have been reported. The emphasis in these studies was to restrain an appropriate amount of self-lubricating phase from the precursors after the laser processing. In the present study, Specimens 2 and 5 demonstrate a comparatively low friction coefficient despite that no self-lubricating additives were introduced to the precursors. It can be concluded that the graphene-like few-layer microstructure within the graphitic phases provides an excellent friction-reducing benefit. The graphitic phases with HCP crystal structure show weak van der Waals force among the basal planes, resulting in low shear strength and friction.44 Their high hardness and wear resistance compared with the others samples placed them among the best of all the specimens studied here.

Formation Mechanism of the Graphitic Phases

According to the traditional cast iron theory, graphitic phases could be obtained under near-equilibrium cooling conditions (relatively low cooling rate45). However, this is not the case in the laser cladding process, where the ultrahigh cooling happens (≥103 k·s–146). As the spherical or nearly spherical graphitic phases composed of graphene-like few-layer microstructures are found to contribute significantly to the coatings’ antifriction behavior, their formation mechanism is studied comprehensively. The origin of carbon atoms is first considered. Both Ni60A and B4C could provide carbon atoms, and the latter is the main carbon source. Besides, it is noteworthy that some uncombined carbon atoms exist in the nanoscale B4C precursor. In the XRD pattern of the nanoscale B4C (Figure 4D), a rhomboidal graphite phase annotated as 3R-graphite can be identified, which is an unstable graphite variant and will transform into hexagonal graphite at temperatures above 1600 °C.47 Since nanoscale B4C is utilized as a precursor in this study, the size of 3R-graphite phase within the raw materials should also be in nanoscale (≤100 nm). However, observation of the carbon-rich zones (Figure 3B–D) indicates that the graphitic spheres are in the size of micrometers (∼1.0–5.0 μm). Nanoscale precursor materials are not able to guarantee a nanoscale product after laser cladding, owing to decomposition of the precursors and accompanied chemical reactions in the molten pool (Figure 1). The final obtained microstructure size relies more on laser process parameters than the particle size of raw materials. The original nanoscale 3R-graphite phase could have transformed into short-range ordered hexagonal regions and thereafter acted as the nucleation core of the spherical or nearly spherical graphitic phases. The 3R-graphite phase is also able to be dissolved into the molten pool under the severe laser heating condition and then react with other alloying elements (e.g., Ti, Cr, etc.) to form carbides.

3R-graphite in the nanoscale precursor is not the sole carbon source to form the graphitic phases. These microstructures in the cladding coatings can also be formed in an in situ manner. Considering the three main elements Ti, B, and C in the prepared coatings, the Ti–B–C ternary phase diagram (Figure 4E)37 suggests that the coexistence of graphitic phase and others is thermodynamically possible (TiB2+C+′B4C′ zone and the TiB2+TiC1–x+C zone). It is reasonable to assume that B and C atoms are released from B4C almost simultaneously under the rapid heating effect of the laser beam, considering the ultrafine size of the B4C particles (80–100 nm). Ti tends to react with B prior to C atoms. Li et al. found that the Gibbs free energy of Ti+2B → TiB2 is lower than that of Ti+C → TiC.48 With B atoms consumed, a surplus of C atoms appear, which provides extra carbon source for the graphitic phase to grow. It is noteworthy that the formation of phases is related with both the thermodynamic and kinetic factors. Despite of the inferior thermodynamic conditions (Figure S8) for some phases (e.g., NiTi2, NiTi, CrB, Cr7C3, and Cr23C6), they are all generated in the present study under the laser cladding condition.

The formation mechanism model of the graphitic phases in the cladding coating is established and illustrated in Figure 5. Under the heating of the laser beam, various atoms (e.g., Ni, Cr, B, C, etc.) arise as a result of Ni60A and B4C precursors decomposition (Figure 5A). Complicated reactions then occur to form multiple phases, including Ni–Ti, Ti–B, TiC, Cr–B, Ni–B, and Cr–C compounds as suggested by the XRD and TEM analysis (Figure 5B). Since the mole fraction of B is significantly higher than C in the precursor and in the molten pool, various borides could first form without long-distance diffusion, consuming a substantial amount of Ti, Ni, and Cr. Subsequently, the elements that are available to react with C decrease, resulting in the generation of uncombined carbon atoms (Figure 5B). Meanwhile, some gas bubbles form in the molten pool due to the reactions with ambient atmosphere during laser cladding. These uncombined carbon atoms tend to precipitate at the bubble/melt interface and grow up according to the heterogeneous nucleation theory (Figure 5C).

Figure 5.

Figure 5

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Schematic illustration of the formation mechanism of the graphitic phase with the graphene-like few-layer microstructures in the cladding coatings. (A) Decomposition of the Ni60A and B4C precursors and the diffusion of Ti atoms from the substrate into the molten pool under the heating of the laser beam. (B) In situ synthesis and the formation of different phases. (C) Surplus C atoms aggregate and precipitate along the interface of gas bubble/melt. (D) Formation mechanism of the spherical graphitic phase, showing the repeated growth and reorientation of the graphitic layers via superassembly.

As to the formation of the graphitic phases in an overall spherical morphology, various theories developed for spheroidal graphitic cast iron can be referred to, including the bubble theory.49 Gas formation in laser cladding is almost inevitable, which is suggested by the pores formed in Specimens 2 and 5 (Figure 2B, D). As gas is formed in the molten pool, graphitic phase tends to nucleate at the interface between the gas and the melt. Further growth follows a superassembly manner as schematically shown in Figure 5D. Owing to its crystal structure, the graphite nuclei then grow along their basal crystal plane (0002), while long-range growth is impossible because of the limited space available. To enable further growth, the later formed graphitic layers will have to reorientate toward the center of the bubble and the growth continues until the whole bubble is filled. This repeated growth and reorientation lead to the formation of the spherical or nearly spherical graphitic phases. As mentioned previously, 3R-graphite present in the nanoscale B4C precursor could also transform into hexagonal graphitic phase and then act as the nucleation core.

Apart from the spherical graphitic phases, nearly spherical morphologies are also observed in Specimens 2 and 5 (indicated by dashed circle in Figure 3B, C). The appearance of nearly spherical graphitic phase is similar to the so-called “graphite degeneration” in ductile irons.50 First, composition fluctuation, which is derived from the mass transfer in the molten pool, leads to the regional difference in composition. At the position with unfavorable composition for nucleation of the spherical graphitic phase, no graphitic phase will be formed or only nearly spherical graphitic phase be generated. Second, the formation of these nearly spherical or degenerated spherical graphitic morphology is also attributed to the harsh cooling conditions during laser cladding. Further, the rapid cooling during the laser cladding process results in nonequilibrium growth of the graphene-like structures, and this, in addition, leads to the high in-plane defect density, which is supported by the highly curved fringes observed by HRTEM (Figure 3G, H).

Appropriate laser specific energy (4.4–8.9 kJ·cm–2) and the amount of nanoscale B4C precursor are critical for the formation of well distributed spherical or nearly sphericalgraphitic phases. A lack of B4C could not supply enough C atoms, leading to an increase in the friction coefficient. A surplus of C atoms, on the other hand, leads to severe reactions between the molten pool and ambient air. The resulting pore formation reduces the mechanical properties of the final coatings. A higher laser power or a lower scanning velocity results in increased Ti dilution that consumes the uncombined C atoms available while the low hardness of Ti weakens the overall wear-resistant property of the prepared coating. Laser cladding parameters together with the precursor compositions can tune the obtained coating microstructure, which further determines their mechanical behaviors.

Conclusions

In this study, Ni60A and nanoscale B4C mixtures are used as the precursor materials for laser cladding on Ti-6Al-4V substrates. With appropriate laser cladding parameters and precursor compositions (i.e., Ni60A base material with 5.0 wt % nanoscale B4C, fabricated using 1.0 kW laser power and 450 mm·min–1 scanning velocity), the wear resistance of the coating shows a remarkable 43.67 times enhancement with excellent friction coefficient than that of the substrate. The simultaneously obtained low friction coefficient (∼0.35) is mainly attributed to the distributed spherical or nearly sphericalgraphitic phases. The formation mechanism of this graphitic phase is discussed from the perspectives of raw material introduction and the laser cladding induced in situ superassembly. Repeated growth and reorientation of the graphene-like microstructure during the laser process is found to play a vital role in the model established for the formation mechanism. To obtain tailored microstructure and properties of the final coatings, the laser process parameters should be strictly controlled with carefully designed precursor material compositions. In the future, the formation mechanism of the graphene-like few-layer microstructures and their superassembly will be further studied with more microstructure characterizations. Extended precursor systems will also be explored to obtain the graphitic structure under the nonequilibrium laser process condition. Furthermore, synthesis of large-scale graphene-related materials and further incorporation into more functional composite materials will also be investigated.

Acknowledgments

This work was supported by the Key Research and Development Program of Shandong Province (grant number 2018CXGC0811 and 2019JZZY010361), the National Key Research and Development Program of China (2017YFA0206901, 2017YFA0206900), and the Natural Science Foundation of Shandong Province (ZR2017MEE063).

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acscentsci.0c01365.

  • Additional tables and figures including the list of detailed laser process parameters, bonding zone morphology, discontinuous cladding track, results of elemental distribution, XRD analysis, supplementary TEM analysis, microhardness distribution (PDF)

Author Contributions

F.W., C.H., and R.Z. contributed equally to this work.

The authors declare no competing financial interest.

Supplementary Material

oc0c01365_si_001.pdf (2.3MB, pdf)

References

  1. Kong B.; Zu L.; Peng C.; Zhang Y.; Zhang W.; Tang J.; Selomulya C.; Zhang L.; Chen H.; Wang Y.; Liu Y.; He H.; Wei J.; Lin X.; Luo W.; Yang J.; Zhao Z.; Liu Y.; Yang J.; Zhao D. Direct Superassemblies of Freestanding Metal-Carbon Frameworks Featuring Reversible Crystalline-Phase Transformation for Electrochemical Sodium Storage. J. Am. Chem. Soc. 2016, 138 (50), 16533–16541. 10.1021/jacs.6b10782. [DOI] [PubMed] [Google Scholar]
  2. Lu Z.; Gao D.; Yi D.; Yang Y.; Wang X.; Yao J. Sp2/Sp3 Hybridized Carbon as an Anode with Extra Li-Ion Storage Capacity: Construction and Origin. ACS Cent. Sci. 2020, 6 (8), 1451–1459. 10.1021/acscentsci.0c00593. [DOI] [PMC free article] [PubMed] [Google Scholar]
  3. Li Y.; Lei X.; Yuan Y.; Wu S.; Han B.; Liu X.; Liu W.; Hu J.; Yang C.; Lin Z.; Lu J. Fe2P-Decorated N,P Codoped Carbon Synthesized via Direct Biological Recycling for Endurable Sulfur Encapsulation. ACS Cent. Sci. 2020, 6 (10), 1827–1834. 10.1021/acscentsci.0c00899. [DOI] [PMC free article] [PubMed] [Google Scholar]
  4. Lu Z.; Chen G.; Siahrostami S.; Chen Z.; Liu K.; Xie J.; Liao L.; Wu T.; Lin D.; Liu Y.; Jaramillo T. F.; Nørskov J. K.; Cui Y. High-Efficiency Oxygen Reduction to Hydrogen Peroxide Catalysed by Oxidized Carbon Materials. Nat. Catal. 2018, 1 (2), 156–162. 10.1038/s41929-017-0017-x. [DOI] [Google Scholar]
  5. Chen P.; Peng H. High-Performance Graphene Fibers Enabled by Hydration. ACS Cent. Sci. 2020, 6 (7), 1040–1042. 10.1021/acscentsci.0c00820. [DOI] [PMC free article] [PubMed] [Google Scholar]
  6. Ma L.; Lin H.; Zhang W.; Zhao P.; Zhu G.; Hu Y.; Chen R.; Tie Z.; Liu J.; Jin Z. Nitrogen-Doped Carbon Nanotube Forests Planted on Cobalt Nanoflowers as Polysulfide Mediator for Ultralow Self-Discharge and High Areal-Capacity Lithium-Sulfur Batteries. Nano Lett. 2018, 18 (12), 7949–7954. 10.1021/acs.nanolett.8b03906. [DOI] [PubMed] [Google Scholar]
  7. Sambyal P.; Iqbal A.; Hong J.; Kim H.; Kim M.-K.; Hong S. M.; Han M.; Gogotsi Y.; Koo C. M. Ultralight and Mechanically Robust Ti3C2Tx Hybrid Aerogel Reinforced by Carbon Nanotubes for Electromagnetic Interference Shielding. ACS Appl. Mater. Interfaces 2019, 11 (41), 38046–38054. 10.1021/acsami.9b12550. [DOI] [PubMed] [Google Scholar]
  8. Rana M.; He Q.; Luo B.; Lin T.; Ran L.; Li M.; Gentle I.; Knibbe R. Multifunctional Effects of Sulfonyl-Anchored, Dual-Doped Multilayered Graphene for High Areal Capacity Lithium Sulfur Batteries. ACS Cent. Sci. 2019, 5 (12), 1946–1958. 10.1021/acscentsci.9b01005. [DOI] [PMC free article] [PubMed] [Google Scholar]
  9. Wang D.; Zhang W.; Drewett N. E.; Liu X.; Yoo S. J.; Lee S.-G.; Kim J.-G.; Deng T.; Zhang X.; Shi X.; Zheng W. Exploiting Anti-T-Shaped Graphene Architecture to Form Low Tortuosity, Sieve-like Interfaces for High-Performance Anodes for Li-Based Cells. ACS Cent. Sci. 2018, 4 (1), 81–88. 10.1021/acscentsci.7b00444. [DOI] [PMC free article] [PubMed] [Google Scholar]
  10. Zheng G.; Lee S. W.; Liang Z.; Lee H.-W.; Yan K.; Yao H.; Wang H.; Li W.; Chu S.; Cui Y. Interconnected Hollow Carbon Nanospheres for Stable Lithium Metal Anodes. Nat. Nanotechnol. 2014, 9 (8), 618–623. 10.1038/nnano.2014.152. [DOI] [PubMed] [Google Scholar]
  11. Qi D.; Liu Z.; Liu Y.; Leow W. R.; Zhu B.; Yang H.; Yu J.; Wang W.; Wang H.; Yin S.; Chen X. Suspended Wavy Graphene Microribbons for Highly Stretchable Microsupercapacitors. Adv. Mater. 2015, 27 (37), 5559–5566. 10.1002/adma.201502549. [DOI] [PubMed] [Google Scholar]
  12. Kong B.; Tang J.; Zhang Y.; Jiang T.; Gong X.; Peng C.; Wei J.; Yang J.; Wang Y.; Wang X.; Zheng G.; Selomulya C.; Zhao D. Incorporation of Well-Dispersed Sub-5-Nm Graphitic Pencil Nanodots into Ordered Mesoporous Frameworks. Nat. Chem. 2016, 8 (2), 171–178. 10.1038/nchem.2405. [DOI] [PubMed] [Google Scholar]
  13. Kong B.; Tang J.; Zhang Y.; Selomulya C.; Gong X.; Liu Y.; Zhang W.; Yang J.; Wang W.; Sun X.; Wang Y.; Zheng G.; Zhao D. Branched Artificial Nanofinger Arrays by Mesoporous Interfacial Atomic Rearrangement. J. Am. Chem. Soc. 2015, 137 (12), 4260–4266. 10.1021/jacs.5b01747. [DOI] [PubMed] [Google Scholar]
  14. Niu Z.; Liu L.; Zhang L.; Shao Q.; Zhou W.; Chen X.; Xie S. A Universal Strategy to Prepare Functional Porous Graphene Hybrid Architectures. Adv. Mater. 2014, 26 (22), 3681–3687. 10.1002/adma.201400143. [DOI] [PubMed] [Google Scholar]
  15. Schmidt M.; Merklein M.; Bourell D.; Dimitrov D.; Hausotte T.; Wegener K.; Overmeyer L.; Vollertsen F.; Levy G. N. Laser Based Additive Manufacturing in Industry and Academia. CIRP Ann. 2017, 66 (2), 561–583. 10.1016/j.cirp.2017.05.011. [DOI] [Google Scholar]
  16. Joe D. J.; Kim S.; Park J. H.; Park D. Y.; Lee H. E.; Im T. H.; Choi I.; Ruoff R. S.; Lee K. J. Laser-Material Interactions for Flexible Applications. Adv. Mater. 2017, 29 (26), 1606586. 10.1002/adma.201606586. [DOI] [PubMed] [Google Scholar]
  17. Alqurashi T.; Butt H. Highly Flexible, Stretchable, and Tunable Optical Diffusers with Mechanically Switchable Wettability Surfaces. ACS Cent. Sci. 2019, 5 (6), 1002–1009. 10.1021/acscentsci.9b00155. [DOI] [PMC free article] [PubMed] [Google Scholar]
  18. Wang H.; Wu Y.; Ma M.; Dong S.; Li Q.; Du J.; Zhang H.; Xu Q. Pulsed Laser Deposition of CsPbBr3 Films for Application in Perovskite Solar Cells. ACS Appl. Energy Mater. 2019, 2 (3), 2305–2312. 10.1021/acsaem.9b00130. [DOI] [Google Scholar]
  19. Lu J.; Lu H.; Xu X.; Yao J.; Cai J.; Luo K. High-Performance Integrated Additive Manufacturing with Laser Shock Peening -Induced Microstructural Evolution and Improvement in Mechanical Properties of Ti6Al4V Alloy Components. Int. J. Mach. Tools Manuf. 2020, 148, 103475. 10.1016/j.ijmachtools.2019.103475. [DOI] [Google Scholar]
  20. Kalentics N.; de Seijas M. O. V.; Griffiths S.; Leinenbach C.; Logé R. E. 3D Laser Shock Peening - A New Method for Improving Fatigue Properties of Selective Laser Melted Parts. Addit. Manuf. 2020, 33, 101112. 10.1016/j.addma.2020.101112. [DOI] [Google Scholar]
  21. Sedao X.; Shugaev M. V.; Wu C.; Douillard T.; Esnouf C.; Maurice C.; Reynaud S.; Pigeon F.; Garrelie F.; Zhigilei L. V.; Colombier J.-P. Growth Twinning and Generation of High-Frequency Surface Nanostructures in Ultrafast Laser-Induced Transient Melting and Resolidification. ACS Nano 2016, 10 (7), 6995–7007. 10.1021/acsnano.6b02970. [DOI] [PubMed] [Google Scholar]
  22. Rakesh K. R.; Bontha S.; Ramesh M. R.; Das M.; Balla V. K. Laser Surface Melting of Mg-Zn-Dy Alloy for Better Wettability and Corrosion Resistance for Biodegradable Implant Applications. Appl. Surf. Sci. 2019, 480, 70–82. 10.1016/j.apsusc.2019.02.167. [DOI] [Google Scholar]
  23. Sun S. D.; Barr C.; Brandt M. In Situ Control of Tempered Martensite during Laser Cladding Repair of Aero-Grade 300M Steel Using AISI 420 Stainless Steel Powder. J. Laser Appl. 2018, 30 (3), 032502. 10.2351/1.5040629. [DOI] [Google Scholar]
  24. Tuominen J.; Kaubisch M.; Thieme S.; Näkki J.; Nowotny S.; Vuoristo P. Laser Strip Cladding for Large Area Metal Deposition. Addit. Manuf. 2019, 27, 208–216. 10.1016/j.addma.2019.01.008. [DOI] [Google Scholar]
  25. Weng F.; Chen C.; Yu H. Research Status of Laser Cladding on Titanium and Its Alloys: A Review. Mater. Eng. 2014, 58, 412–425. 10.1016/j.matdes.2014.01.077. [DOI] [Google Scholar]
  26. Liu J.; Yu H.; Chen C.; Weng F.; Dai J. Research and Development Status of Laser Cladding on Magnesium Alloys: A Review. Opt. Lasers Eng. 2017, 93, 195–210. 10.1016/j.optlaseng.2017.02.007. [DOI] [Google Scholar]
  27. Bai L. L.; Li J.; Chen J. L.; Song R.; Shao J. Z.; Qu C. C. Effect of the Content of B4C on Microstructural Evolution and Wear Behaviors of the Laser-Clad Coatings Fabricated on Ti6Al4V. Opt. Laser Technol. 2016, 76, 33–45. 10.1016/j.optlastec.2015.07.010. [DOI] [Google Scholar]
  28. Chehrghani A.; Torkamany M. J.; Hamedi M. J.; Sabbaghzadeh J. Numerical Modeling and Experimental Investigation of TiC Formation on Titanium Surface Pre-Coated by Graphite under Pulsed Laser Irradiation. Appl. Surf. Sci. 2012, 258 (6), 2068–2076. 10.1016/j.apsusc.2011.04.064. [DOI] [Google Scholar]
  29. Ding L.; Hu S. Effect of Nano-CeO2 on Microstructure and Wear Resistance of Co-Based Coatings. Surf. Coat. Technol. 2015, 276, 565–572. 10.1016/j.surfcoat.2015.06.014. [DOI] [Google Scholar]
  30. Sundaraselvan S.; Senthilkumar N.; Tamizharasan T.; Sait A. N. Surface Modification of AZ61 Magnesium Alloy with Nano TiO2/Al2O3 Using Laser Cladding Technique. Mater. Today Proc. 2020, 21, 717–721. 10.1016/j.matpr.2019.06.745. [DOI] [Google Scholar]
  31. Masanta M.; Shariff S. M.; Roy Choudhury A. Microstructure and Properties of TiB2-TiC-Al2O3 Coating Prepared by Laser Assisted SHS and Subsequent Cladding with Micro-/Nano-TiO2 as Precursor Constituent. Mater. Des. 2016, 90, 307–317. 10.1016/j.matdes.2015.10.135. [DOI] [Google Scholar]
  32. Liu Y.; Gao S.; Zhao W. Laser Cladding of Mg-Zn-Y-Zr Alloy Using Al with Different Ratios of Micro- and Nano-SiC Powder. J. Mater. Sci. 2020, 55 (27), 13270–13293. 10.1007/s10853-020-04893-z. [DOI] [Google Scholar]
  33. Gao Z.; Ren H.; Yu Y.; Guo W.; Zhang C.; Gao Z. Effect of Nano-WC Content on Microstructures and Wear Resistance of Laser Cladding Fe-Based Alloy Coatings. Mater. Res. Express 2020, 7 (9), 096515. 10.1088/2053-1591/abb631. [DOI] [Google Scholar]
  34. Lv Y. H.; Li J.; Tao Y. F.; Hu L. F. Oxidation Behaviors of the TiNi/Ti2Ni Matrix Composite Coatings with Different Contents of TaC Addition Fabricated on Ti6Al4V by Laser Cladding. J. Alloys Compd. 2016, 679, 202–212. 10.1016/j.jallcom.2016.04.037. [DOI] [Google Scholar]
  35. Weng F.; Yu H.; Chen C.; Dai J. Microstructures and Wear Properties of Laser Cladding Co-Based Composite Coatings on Ti-6Al-4V. Mater. Eng. 2015, 80, 174–181. 10.1016/j.matdes.2015.05.005. [DOI] [Google Scholar]
  36. Ma N.; Chen J.; Huang Z.; Liu X.; Chen Z. Densification of C/SiC Composite Surface by the Hybrid Process of Laser Cladding and Subsequent Heat Treatment. Ceram. Int. 2019, 45 (6), 7703–7708. 10.1016/j.ceramint.2019.01.071. [DOI] [Google Scholar]
  37. Duschanek H.; Rogl P.; Lukas H. L. A Critical Assessment and Thermodynamic Calculation of the Boron-Carbon-Titanium (B-C-Ti) Ternary System. J. Phase Equilib. 1995, 16 (1), 46–60. 10.1007/BF02646248. [DOI] [Google Scholar]
  38. Niiyama Y.; Shimizu N.; Kuwayama A.; Okada H.; Takeno T.; Kurihara K.; Adachi K. Effect of Running-in for Delamination and Friction Properties of Self-Mating Diamond-like Carbon Coatings in Water. Wear 2017, 378–379, 27–34. 10.1016/j.wear.2017.01.100. [DOI] [Google Scholar]
  39. Friction, Lubrication and Wear; Chowdhury M. A., Ed.; Intech Open, 2019. [Google Scholar]
  40. Jiang J.; Li R.; Yuan T.; Niu P.; Chen C.; Zhou K. Microstructural Evolution and Wear Performance of the High-Entropy FeMnCoCr Alloy/TiC/CaF2 Self-Lubricating Composite Coatings on Copper Prepared by Laser Cladding for Continuous Casting Mold. J. Mater. Res. 2019, 34 (10), 1714–1725. 10.1557/jmr.2019.57. [DOI] [Google Scholar]
  41. Gao Q.; Yan H.; Qin Y.; Zhang P.; Guo J.; Chen Z.; Yu Z. Laser Cladding Ti-Ni/TiN/TiW+TiS/WS2 Self-Lubricating Wear Resistant Composite Coating on Ti-6Al-4V Alloy. Opt. Laser Technol. 2019, 113, 182–191. 10.1016/j.optlastec.2018.12.046. [DOI] [Google Scholar]
  42. Torres H.; Vuchkov T.; Rodríguez Ripoll M.; Prakash B. Tribological Behaviour of MoS2-Based Self-Lubricating Laser Cladding for Use in High Temperature Applications. Tribol. Int. 2018, 126, 153–165. 10.1016/j.triboint.2018.05.015. [DOI] [Google Scholar]
  43. Lu X.-L.; Liu X.-B.; Yu P.-C.; Qiao S.-J.; Zhai Y.-J.; Wang M.-D.; Chen Y.; Xu D. Synthesis and Characterization of Ni60-HBN High Temperature Self-Lubricating Anti-Wear Composite Coatings on Ti6Al4V Alloy by Laser Cladding. Opt. Laser Technol. 2016, 78, 87–94. 10.1016/j.optlastec.2015.10.005. [DOI] [Google Scholar]
  44. Skinner J.; Gane N.; Tabor D. Micro-Friction of Graphite. Nature, Phys. Sci. 1971, 232 (35), 195–196. 10.1038/physci232195a0. [DOI] [Google Scholar]
  45. Qing J.; Richards V. L.; Van Aken D. C. Growth Stages and Hexagonal-Rhombohedral Structural Arrangements in Spheroidal Graphite Observed in Ductile Iron. Carbon 2017, 116, 456–469. 10.1016/j.carbon.2017.01.063. [DOI] [Google Scholar]
  46. Xiao W.; Li S.; Wang C.; Shi Y.; Mazumder J.; Xing H.; Song L. Multi-Scale Simulation of Dendrite Growth for Direct Energy Deposition of Nickel-Based Superalloys. Mater. Des. 2019, 164, 107553. 10.1016/j.matdes.2018.107553. [DOI] [Google Scholar]
  47. Yi C. Q. O. M. W.; Qunji F. S. X. A Study on the Composition and Structure of Chinese Natural Graphites. Tribology 1992, (3), 9. [Google Scholar]
  48. Li J.; Yu Z.; Wang H.; Li M. Microstructural Characterization of Titanium Matrix Composite Coatings Reinforced by in Situ Synthesized TiB + TiC Fabricated on Ti6Al4V by Laser Cladding. Rare Met. 2010, 29 (5), 465–472. 10.1007/s12598-010-0151-y. [DOI] [Google Scholar]
  49. Yamamoto S.; Chang B.; Kawano Y.; Ozaki R.; Murakami Y. A Proposed Theory of Nodularization of Graphite in Cast Irons. Met. Sci. 1975, 9 (1), 360–369. 10.1179/030634575790444739. [DOI] [Google Scholar]
  50. Li Z.; Li Y. Evaluation of Melt Quality and Graphite Degeneration Prediction in Heavy Section Ductile Iron. Metall. Mater. Trans. A 2005, 36 (9), 2455–2460. 10.1007/s11661-005-0119-y. [DOI] [Google Scholar]

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