Influence of boron doping level on the basic mechanical properties and erosion behavior of boron-doped micro-crystalline diamond (BDMCD) film

DIAMAT-06719; No of Pages 14 Diamond & Related Materials xxx (2016) xxx–xxx

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Influence of boron doping level on the basic mechanical properties and erosion behavior of boron-doped micro-crystalline diamond (BDMCD) film Xinchang Wang ⁎, Xiaotian Shen, Fanghong Sun, Bin Shen State Key Laboratory of Mechanical System and Vibration, School of Mechanical Engineering, Shanghai Jiao Tong University, Shanghai 200240, China

a r t i c l e

i n f o

Article history: Received 23 May 2016 Received in revised form 27 September 2016 Accepted 27 September 2016 Available online xxxx Keywords: BDMCD film Boron doping level Basic mechanical properties Film-substrate adhesive strength Erosion behavior

a b s t r a c t For the chemical vapor deposition (CVD) of diamond films, boron doping with appropriate boron doping levels can enhance the basic mechanical properties of the as-deposited films, especially the film-substrate adhesive strength. However, the boron doping level, which is thought to play a critical role in modifying the properties of the boron-doped diamond (BDD) film, must be further investigated. In the present investigation, the commonly used reaction-bonded silicon carbide (RB-SiC) material is selected as the substrate upon which boron-doped micro-crystalline (BDMCD) films with similar film thickness (24.8–26.3 μm) are synthesized using mixed reactant gas with different B/C atomic ratios, i.e., different boron doping levels. Systematic characterization and solid particle erosion tests are conducted on all specimens to elucidate the influence of the boron doping level on the diamond films' basic mechanical properties and erosion behavior and to elaborate the impact velocity and impact angle dependence of the erosion behavior. The results demonstrate that moderate boron doping levels (5000 and 8000 ppm) are able to maximize the growth rate, reduce the surface roughness, guarantee diamond quality and hardness, minimize the residual stress, and significantly enhance the film-substrate adhesive strength, thereby providing favorable erosion behavior. By contrast, very high boron doping levels (12,000 and 16,000 ppm) deteriorate the erosion behavior of as-deposited BDMCD films, mainly because of the excessive reduction of the film quality and hardness, together with the deterioration of the film-substrate adhesive strength caused by the transformation of the residual stress from compressive to tensile. With increasing either impact velocity or impact angle, the stable erosion rates of all specimens increase, and the film lifetime of the coated specimens become shorter. Moreover, the impact velocity dependence of stable erosion rates for diamond-coated specimens is considerably stronger than that for the uncoated RB-SiC specimen, as indicated by a much higher velocity exponent. © 2016 Elsevier B.V. All rights reserved.

1. Introduction Diamond films deposited by chemical vapor deposition (CVD) have excellent properties, such as extremely high hardness, high elastic modulus [1], high thermal conductivity [2], high chemical inertness [3], a low thermal expansion coefficient and favorable wear resistance, including erosive wear resistance [4,5]. Thus, these films are demonstrated to be an efficacious protective coating on various components with working surfaces that suffer severe impacts by solid particles or slurries, such as relief valves that control the flow rate in coal liquefaction equipment [6], choke valves widely used in the offshore oil industry [7], and nozzles adopted for spray drying systems [8]. For applications of diamond-coated components, the extremely high hardness of the diamond film can provide sufficient wear resistance for ⁎ Corresponding author at: Mechanical Building B344, School of Mechanical Engineering, Shanghai Jiao Tong University, Dongchuan Road 800, Minhang District, Shanghai 200240, China. E-mail address: [email protected] (X. Wang).

the working surface, while insufficient film-substrate adhesive strength is a major vulnerability that plays a substantial role in the lifetime of the diamond-coated component [9]. For cobalt cemented tungsten carbide (WC-Co) substrates, the presence of Co is well known to cause significant graphitization at the film-substrate interface at the early nucleation stage and to negatively affect the adhesive strength [10]. Another factor that has an adverse effect on the adhesive strength is the accumulation of either intrinsic or thermal residual stress occurring during the nucleation, growth and cooling stages [11,12], which is the main adverse factor for the iron-group-element-free substrates, such as Si, Ta, Ti, Nb, SiC, and Si3N4. In addition to the frequently used traditional or novel pretreatment techniques (such as alkali, acid, and alkali-acid treatments, boronization, and the use of tungsten particles on the substrate surface [13–15]) and intermediate layers (such as SiC, Nb, Cr and Ta [16,17]), dynamic boron-doping technology synchronous with the deposition process is also considered an important method for enhancing the interfacial adhesive quality between the diamond film and the substrate, firstly because the additive boron element can react with the Co and W to form stable Co-W-B compounds that only minimally affect the

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nucleation of the diamond [14]. Moreover, the covalent radius of the B atom (0.085 nm) is larger than that of the C atom (0.077 nm), and thus, the incorporation of B into the sp3 C\\C structures induces a certain amount of intrinsic tensile stress [18]. Because the residual stress in WC-Co or ceramic-based diamond film is mainly compressive residual stress, boron doping technology can reduce the absolute value of the total residual stress [19]. In our previous studies, by selecting a commonly used moderate boron doping level (approximately 5000 ppm in the mixed reactant gas), the enhancement of the film-substrate adhesive strength caused by boron doping has been identified as a positive sign of improving the solid particle erosion behavior of as-deposited diamond film [19] and its application for components that must exhibit erosive wear resistance [8]. In addition, the influences of the substrate material and the film thickness on the erosion performance and mechanism of the boron-doped diamond (BDD) film deposited adopting a boron doping level of 5000 ppm have also been systematically studied [20,21]. Nevertheless, it is well known that BDD films with different properties can be fabricated using different boron doping levels. Boron doping technology is firstly proposed for modifying the electrical properties of the diamond film, mainly changing the insulating diamond film to a conductive film [22]. This observation established that the BDD films fabricated with different doping levels exhibit different conductive pathways and mechanisms [23]. Moreover, the boron doping level also affects the growth rate [24], grain size [25], diamond purity [26] and residual stress [25–27] of as-deposited BDD films with either micro- or nanoscale diamond grains. Thus, there is sufficient evidence to conclude that the boron doping level is likely to influence the erosion behavior. SiC-based diamond films present better adhesion and erosion performance than WC-Co-based diamond films [21], and diamond-coated SiC components have been extensively used in various industries [8, 28,29]. In the present study, reaction-bonded silicon carbide (RB-SiC) is selected as the substrate, upon which boron-doped microcrystalline (BDMCD) films are deposited with different boron doping levels. Further characterization and solid particle erosion tests are conducted on all the specimens for the purpose of clarifying the influence of the boron doping level on their basic mechanical properties and erosion behavior, together with the impact velocity and impact angle dependence of the erosion behavior. 2. Experimental details 2.1. Fabrication of BDMCD-coated RB-SiC specimens The RB-SiC selected in the present research is a typical multicomponent material approximately composed of 80 vol.% α-SiC, 17 vol.% β-SiC and 3 vol.% Si and provided by Debao, Shanghai, China. The deposition of CVD diamond films on SiC or Si3N4 ceramics is generally known to involve fewer requirements for the pretreatment

procedure than that on a WC-Co substrate because there is no element that has harmful effects on the nucleation and growth of the diamond [30]. However, interior or surface defects are unavoidable in the sintered multi-component material [31]; thus, prior to the deposition procedure, the RB-SiC substrate should first be ground to obtain a relatively smooth and integrated surface with an Ra value of approximately 0.5 μm and then ultrasonically cleaned by ethanol for 3 min. Thereafter, the substrate is pretreated by the following simple procedure: scratched by coarse diamond grit (7 μm); ultrasonically cleaned by ethanol for 10 min; ground and seeded by fine diamond grit (250–500 nm), and ultrasonically cleaned by ethanol for 10 min. All the deposition processes are accomplished using a customized bias-enhanced hot filament CVD (BE-HFCVD) apparatus. During the deposition of BDMCD films, trimethyl borate ((CH3O)3B) is selected as the boron source instead of the commonly used boron hydride because of safety concerns. As shown in Fig. 1, the size of the selected RB-SiC substrate is 25 mm × 25 mm × 4 mm; thus, 25–29 samples can be fabricated simultaneously on a ø 220 mm rotational copper worktable. Table 1 lists the common deposition parameters for fabricating different diamond-coated RB-SiC specimens, including the BDMCD-coated samples along with undoped micro-crystalline diamond (UMCD)-coated samples for comparison. The following five important parameters must be controlled during the diamond film deposition process: (1) Flow rate of the reactant gas. To conveniently introduce a minuscule amount of liquid (CH3O)3B into the reactor, liquid acetone is selected as the carbon source, and a mixture of acetone and (CH3O)3B premixed according to a certain B/C ratio is bubbled out and introduced into the reactor by the carrier H2 gas. Because the amount of the (CH3O)3B is small, the flow rate of the mixture can be approximated as the flow rate of the acetone, which can be calculated based on F acetone ¼ F H2

P acetone P−P acetone

ð1Þ

where Facetone is the flow rate of the acetone, FH2 the flow rate of the carrier H2, Pacetone the saturated vapor pressure of the acetone (8.9 kPa at 0 °C), and P the total pressure in the container (101.3 kPa). In addition, the acetone/H2 ratio at the nucleation stage is slightly higher to promote nucleation [32]. (2) Pressure. Within the appropriate range, the relatively lower pressure used at the nucleation stage can promote diamond nucleation because more active groups can move to the substrate surface as a result of their increased mean free paths [13]. (3) Temperature, including the filament temperature and the substrate temperature. The former should be as high as possible to efficiently decompose the reactant gas, especially the H2 because of its high dissociation energy. The filament temperature is limited by the hightemperature stability of the filament material; thus, tantalum wires,

Fig. 1. Schematic of arrangements for substrates and hot filaments during HFCVD deposition processes, together with a picture of the selected RB-SiC substrate.

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Table 1 Common deposition parameters for fabricating different diamond-coated RB-SiC specimens.

Pure H2 flow rate [sccm] Carrier H2 flow rate [sccm] Total H2 flow rate [sccm] Acetone flow rate [sccm] Acetone/H2 proportion [vol.%] Pressure [Pa] Filament number Filament length [mm] Filament diameter [mm] Filament separation [mm] Filament-substrate distance [mm] Total filament power [kW] Filament temperature [°C] Substrate temperature [°C]

Nucleation stage

Growth stage

600 300 900 28.9 3.21 1200 11 240 0.7 22 9.5 ± 0.5 9.0 2200 ± 50 850 ± 25

600 240 840 23.1 2.75 3000

8.5 800 ± 25

which are more resistant to high temperatures than typical wires, are chosen as the filaments. The substrate temperature must be controlled in the range of 700–900 °C. In this study, the filament arrangement was optimized based on previous research to ensure a uniform temperature distribution on all the substrates [33]. Both temperatures are dependent on the total filament power, the filament arrangement and the cooling condition of the apparatus. (4) Time. A special 30 min initial stage with deposition parameters that favor diamond nucleation is employed. Firstly, a constant growth time of 19.5 h is used to deposit all the diamond films and to approximately determine their growth rates. The erosion behavior of diamond films is closely associated with the thickness; thus, for the diamondcoated specimens used in the erosion tests, the growth time is adjusted to obtain a similar film thickness. (5) Boron doping level. The boron doping level is defined as the B/C atomic ratio in the mixed reactant gas. A sufficiently high boron doping level has already been proven to produce metallic conductivity in asdeposited BDMCD films [34,35], and the B/C atomic ratio in the BDMCD film (boron doping concentration) is proportional to the boron doping level in the limited range of 5000–10,000 ppm, but this ratio only slightly increases when the boron doping level is increased from 10,000 to 50,000 ppm [36]. Therefore, in this article five different boron doping levels (2000, 5000, 8000, 12,000 and 16,000 ppm) are selected, which are directly determined by the volume proportion of the (CH3O)3B to the acetone, based on V B ρB

B ¼ 3V

B ρB

 MB

 MB

þ 3V C ρC

ð2Þ

 MC

where B is the boron doping level, VB the volume of (CH3O)3B, ρB the density of (CH3O)3B, MB the molecular weight of (CH3O)3B, VC the volume of acetone, ρC the density of acetone, and MC the molecular weight of acetone. 2.2. Solid particle erosion tests Solid particle erosion tests are conducted in a standard air-sand erosion rig, the schematic diagram of which is shown in Fig. 2, along with a field emission scanning electron microscopy (FESEM) micrograph of the angular silica solid particles with an average diameter of 180 μm (80 mesh) used in the present tests. These particles are provided by Zhiyuan Co., Ltd. in China. In addition to the BDMCD-coated SiC specimens fabricated using different boron doping levels, the UMCD-coated and uncoated RB-SiC specimens are also adopted for comparison. During the erosion test, the air valve can be adjusted to different levels to control the impact velocity v (80–140 m/s), and the actual velocity is measured by the double disk method [37]. The specially designed

Fig. 2. Schematic of the standard air-sand erosion apparatus and FESEM micrograph of angular silica solid particles.

nozzle provides a constricted region to induce the Venturi effect (i.e., increased air speed and reduced pressure) in the compressed air as a result of the reduced cross-sectional area in this region. Therefore, if the suction valve is open, the solid particles in the sand box can be sucked into the nozzle and then carried by the high-speed air to collide with the specimen. The specimen can be rotated to adjust the impact angle α (30–90°). The flux of the solid particles is fixed as 0.55 g/s, and the distance between the exit of the nozzle and the specimen is 20 ± 1 mm. After the impact, both the solid particles and the wear debris are recycled into the recycle box through the recycling holes at the bottom of the test chamber and then separated by the cyclone separator, where the silica particles that exceed the tolerance are recycled into the sand box and the tiny broken particles and wear debris are ejected into the junk box. To compare the erosion behaviors of the as-deposited specimens, the erosion rate ε and the film lifetime lf are adopted as criteria. Herein, the erosion rate ε is defined as the mass loss induced by the impact of a unit mass of solid particles (mg/kg), and the film lifetime lf is defined as the point in time during the erosion test when the erosion rate sharply changes. This time node will be discussed and determined in the following sections based on the variation curve of the erosion rate. 2.3. Detection methodology To characterize the specimens either before or after the erosion tests, a series of detection methodologies are applied. (1) FESEM (ULTRA55, Zeiss, Germany) is employed to observe the surface and cross-sectional morphologies of the as-deposited diamond films, the surface morphology of the uncoated SiC specimen, and the wear topographies after the erosion tests. In addition, the grain size and the thickness of the diamond film are both measured using the FESEM micrographs. (2) The surface roughness values Ra and Ry of all the specimens before and after the erosion tests are measured by a stylus profiler (Dektak 6M, Veeco, USA). Each value is determined by averaging five values obtained by scanning profile curves at five different positions on the film surface, adopting a definite scanning length of 2.0 mm. Because the tip of the profiler can be worn with successive measurements,

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it is regularly checked or calibrated by measuring a standard sample after every 25 measurements, and the tip is changed when necessary. (3) An X-ray diffractometer (XRD, D8 ADVANCE, Bruker, Germany) is used to determine the diamond crystal orientations. The continuous scanning mode is used, and the 2θ scanning range and the step are defined as 20°–90° and 4°/min, respectively. (4) A confocal Raman microscope (Senterra R200-L, Bruker, Germany) containing a laser with a visible excitation wavelength of 532 nm is used to detect the compositions of the diamond films before and after the erosion tests, based on which the approximate residual stress values and boron doping concentrations in the diamond films can also be obtained. (5) The nano-hardness and elastic modulus values of the diamondcoated SiC specimens are measured by the nano-indentation method with n in-situ nanomechanical test system (TI950 TriboIndenter, Hysitron, USA) using the standard Berkovich diamond indenter and a maximum load of 25 mN. Note that the micro-sized diamond film with a relatively rough surface must be polished to an Ra value below 30 nm before the test; therefore, the detection result is unavoidably influenced by the polishing process. The computing method is described elsewhere [38]. (6) The Rockwell hardness tester (HR-150 A, Shanghai Lianer, China) is used to qualitatively evaluate the adhesive strengths of different diamond-coated specimens by observing the corresponding indentation morphologies. The indentation tests are conducted using a diamond indenter (angle = 120°, radius = 0.2 mm) and a static load of 980 N. Based on the indentation morphology, the as-measured nanohardness and the elastic modulus, the fracture strength of the

diamond-coated specimen is calculated as described elsewhere [39–41] to characterize the film-substrate interface characteristics. (7) Mass losses of all the specimens during the erosion tests are measured by a precision balance (ESJ182-4, 32 g/0.01 mg, 182 g/0.1 mg, Shenyang Longteng, China). Before each measurement, the specimen is successively dipped into acetone and deionized water and ultrasonically cleaned to thoroughly remove any solid particles. 3. Results and discussion 3.1. Characterizations and basic mechanical properties 3.1.1. Cross-sectional morphology, growth rate and film thickness Cross-sectional morphologies of BDMCD-coated RB-SiC specimens synthesized using different boron doping levels and the UMCD-coated specimen (growth time = 19.5 h) are firstly observed, as shown in Fig. 3. The actual film thickness is calculated by averaging five values obtained by observing five different scattered sampling points, and the calculated results of both the film thickness and the growth rate (i.e., the film thickness divided by the growth time) are plotted in Fig. 4. Among previous reports, different relationships between the growth rate and the boron doping level have been presented, depending on the CVD technology, substrate, carbon source, boron source, deposition parameters, among other conditions [24,25,42]. In the present study, the existence of (CH3O)3B in the mixed reactant gas (H2 + acetone) is found to enhance the growth of the HFCVD diamond film on the RBSiC substrate. Specifically, with boron doping levels of 0–8000 ppm, the growth rate increases with the boron doping level because the

Fig. 3. Cross-sectional morphologies of BDMCD-coated RB-SiC specimens synthesized with different boron doping levels and the UMCD-coated specimen for comparison (growth time = 19.5 h).

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Fig. 4. Variations of the film thickness and growth rate with the boron doping level (0, 2000, 5000, 8000, 12,000 and 16,000 ppm, growth time = 19.5 h), together with the film thicknesses of the diamond-coated specimens applied for erosion tests.

increased O/C ratio can promote diamond growth [43], and a balanced amount of boron may enhance the activity of the carbon groups involved in the chemical reaction [44]. However, as the level continues to increase from 8000 to 16,000 ppm, the growth rate gradually decreases because a large amount of boron reverses the usual behavior wherein the graphite is difficult to form or is etched by the H atoms when depositing diamond films by CVD technology [45]. Under such conditions, the increased formation of graphite, especially graphitic lamellae, can deteriorate the quality of the diamond film and inhibit the growth of the diamond film. Because the erosion behavior of the diamond film is strongly dependent on the film thickness [20], the growth durations (excluding the nucleation duration) for different diamond-coated specimens adopted for the final erosion tests are defined as 0 ppm–27.5 h; 2000 ppm–19.5 h; 5000 ppm–17.5 h; 8000 ppm–17 h; 12,000 ppm–19 h; 16,000 ppm– 20.5 h, and the actual film thicknesses of the as-synthesized specimens are accurately controlled in the range of 24.8–26.3 μm, as plotted in Fig. 4. 3.1.2. Surface morphology, grain size and surface roughness The surface morphologies of different diamond-coated specimens with similar film thickness (24.8–26.3 μm) and the uncoated RB-SiC specimen are presented in Fig. 5. For each diamond film, the grain size is obtained by counting the measured size values on FESEM micrographs at five different positions. Fig. 6 lists the statistical grain sizes and surface roughness values (Ra and Ry values) of different specimens. The RB-SiC substrate used in the present study has been roughly ground and exhibits relatively low surface roughness (Ra = 58.94 nm, Ry = 71.3 nm). The coverage of diamond films with thicknesses of tens of micrometers will significantly increase the surface roughness, which is mainly attributed to the micro-sized diamond grains on the surface, which protrude to form peaks and valleys. The BDMCD films fabricated using boron doping levels from 0 to 8000 ppm all show well-shaped diamond grains, while clearly distorted grain shapes appear when the boron doping level rises to 12,000 ppm, and many micro cavities exist because of the formation of graphite and the poor film quality caused by the large amount of boron, which will be further established and discussed later (in terms of the Raman characterizations). The grain

size shows a decreasing trend with increasing boron doping because the increase in the O/C ratio and the presence of the boron can enhance the activity of the carbon groups and promote secondary nucleation [44]. When the boron doping level increases from 12,000 to 16,000 ppm, much more grain refinement occurs, which is expected to be useful in many applications and requires further study. Correlatively, the Ra value also exhibits an approximate downtrend with increasing boron doping, but the Ry value instead increases when the boron doping level reaches 12,000 and 16,000 ppm, which is expected to be associated with the distortions of the crystal shapes and the cavities on the surface. 3.1.3. XRD spectra and diamond crystal orientation As presented in Fig. 7, there are similar random crystal textures in all the diamond-coated RB-SiC specimens, mainly including two diffraction peaks corresponding to diamond reflections, two diffraction peaks corresponding to Si reflections and many peaks corresponding to SiC reflections. The RB-SiC used in the present study has been reported to contain 78.2% α-SiC, 18.4% β-SiC and 3.4% Si; thus, the existence of manifold diffraction peaks corresponding to the Si and SiC reflections is logical. (This finding is outside the scope of this research and is not discussed in detail.) Note that in the present 2θ scanning range of 20°– 90°, the two peaks corresponding to the (111) and (220) reflections of the diamond are located at approximately 43.9° and 75.3°, respectively; for different diamond films, expected differences exist between the intensity ratios of the (1 1 1) peak to the (2 2 0) peak (calculated based on the peak area), as listed in Table 2. As the boron doping level is increased from 0 to 12,000 ppm, this intensity ratio gradually decreases, indicating that the dominant texture of the as-deposited diamond film changes from the [111] texture to the [220] texture. However, when the boron doping level is increased further to 16,000 ppm, the [1 1 1] texture predominates again. This variation is consistent with a study on freestanding BDD films deposited on molybdenum substrates [25], but it is slightly different from the research on BDD films grown on WC-Co substrates [46]. Diamond {111} planes can be oxidized more easily than {220} planes [47]; thus, it is suggested that the greater quantity of oxygen provided by the boron doping promotes the preferable appearance of the stable {2 2 0} planes, but the predomination of the

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Fig. 5. Surface morphologies of different diamond-coated specimens with similar film thickness (24.8–26.3 μm) and the uncoated RB-SiC specimen.

[1 1 1] texture again when the boron doping level is 16,000 ppm still requires further investigation. 3.1.4. Raman spectra, boron doping concentration and residual stress The as-obtained Raman spectra of the diamond-coated RB-SiC specimens are plotted in Fig. 8. BDD films fabricated with light boron doping levels are known to often exhibit large photoluminescence (PL) backgrounds in their Raman spectra [23], as shown in Fig. 8b–d, because the boron compensated for the nitrogen-related defects within the diamond that give rise to the majority of the PL; with increasing boron doping level, the PL background gradually declines. To analyze the Raman data, the section between 200 and 1800 cm−1 is selected to ignore the meaningless peak in the section before 200 cm−1. The PL background is subtracted by fitting a polynomial curve, and then the typical Raman peaks are fitted with a combination of Gaussian and Lorentzian line shapes. Note that in the following descriptions, the wavenumber and the full width half maximum (FWHM) values refer to the spectra after the peak-fit processing, as exemplified in Fig. 8g.

As shown in Fig. 8a, for the Raman spectrum of the UMCD film, a pronounced peak located at approximately 1338.29 cm−1 with an FWHM of only 9.51 cm−1 and the absence of clear graphite or amorphous carbon peaks confirm the high diamond purity and quality. Boron doping can induce two new peaks in the Raman spectrum of the diamond film located at approximately 500 cm− 1 and 1200 cm− 1, which are characteristic peaks for the BDD film. In addition, the graphitic G band at approximately 1580 cm−1 is also present and becomes more prominent with increasing boron doping, demonstrating that more nondiamond impurities form as a result of the boron doping, especially at the highest level of doping [26]. As discussed above, a large amount of boron reverses the usual behavior of graphite, which is difficult to form or which can be etched by the H atoms under the condition of CVD diamond film growth, thereby enhancing the formation of the graphite, especially graphitic lamellae [45]. It is significant that for the typical peaks of the BDMCD, there exist asymmetries with an enhancement of one flank of the line and a decrease or a minimum (antiresonance) of the other (referred to as the Fano line shape),

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Fig. 6. Grain sizes and surface roughness values (Ra and Ry values) of different diamond-coated specimens with similar film thickness (24.8–26.3 μm) and the uncoated RB-SiC specimen.

occurring when a discrete transition undergoes quantum mechanical interference with a continuum [48]. The intensities of the boron-doping-related peaks relative to the sp3 diamond peak appear to increase with the boron doping level, and the sp3 diamond peak is nearly covered by the 1200 cm−1 peak when the boron doping level is as high as 16,000 ppm, qualitatively proving that the boron doping concentration can indeed be enhanced by increasing the boron doping level in the mixed reactant gas. For the BDMCD film, it has been proposed that the boron doping concentration can be approximately estimated using the wavenumber position of the Lorentzian component of the 500 cm−1 peak (the 500 peak for short), based on. ½B=cm−3 ¼ 8:44  1030 exp:ð−0:048ωÞ

ð3Þ

where [B] is the boron doping concentration and ω is the exact position of the 500 peak [23]. The referenced study reports that the deviations of the calculated values of the B concentrations for the BDMCD films are all within a factor of 2 of those measured by the secondary ion mass spectroscopy (SIMS), showing that for BDMCD films, the equation based on

the Raman spectrum for assessing the B concentration is accurate. It is apparent that with increasing boron doping, the 500 peak shifts to lower wavenumbers; the corresponding boron doping concentrations are listed in Table 2, indicating that the boron doping concentration increases with the boron doping level in the range of 2000–12,000 ppm, but it only slightly increases when the boron doping level is increased further to 16,000 ppm, in accordance with the reference [36]. The residual stress in the diamond film can be estimated based on the shift of the sp3 diamond peak according to [49,50]: σ ¼ −0:567 ðν−ν d Þ ðGPaÞ

ð4Þ

where σ is the estimated residual stress in the diamond film, v is the wavenumber of the measured diamond peak, νd is the wavenumber of the stress-free diamond peak (1332.4 cm−1), and the negative value indicates that the residual stress is compressive. The estimated residual stress values are also listed in Table 2, indicating that as the boron doping level increases from 2000 to 12,000 ppm, the stress value gradually increases, the stress type transforms from compressive to tensile, and the minimum absolute value of the residual stress occurs when the boron doping level approaches 8000 ppm. The residual stress in diamond films is known to mainly include the growth residual stress caused by distortions or impurities in the film and the thermal residual stress caused by the difference between the thermal expansion coefficients of the diamond film and the substrate. Because the covalent radius of the oxygen atom (0.066 nm) provided by the acetone is smaller than that of the carbon atom (0.077 nm), the C\\O bonds formed in Table 2 Processed data based on XRD and Raman spectra.

Fig. 7. XRD spectra of different diamond-coated specimens with similar film thickness (24.8–26.3 μm).

Boron doping level [ppm]

XRD data Area of the diamond (111) peak

Area of the diamond (220) peak

(111)/(220) intensity ratio

Raman data Boron doping concentration (×1020)

Residual stress (GPa)

0 2000 5000 8000 12,000 16,000

7638 5512 6327 4504 4502 6380

1501 4186 6548 8175 10,396 2245

5.09 1.32 0.97 0.55 0.43 2.84

0 1.757 2.171 3.316 4.345 4.389

−3.34 −2.32 −1.45 −0.08 5.72 4.71

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Fig. 8. Raman spectra of different diamond-coated specimens with similar film thickness (24.8–26.3 μm): (a) UMCD; (b) 2000 ppm BDD; (c) 5000 ppm BDD; (d) 8000 ppm BDD; (e) 12,000 ppm BDD; (f) 16,000 ppm BDD; (g) 16,000 ppm BDD (after the peak-fit processing).

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X. Wang et al. / Diamond & Related Materials xxx (2016) xxx–xxx Table 3 Mechanical properties of diamond films fabricated with different boron doping levels. Boron doping level [ppm]

Mechanical property Nano-hardness [GPa]

Elastic modulus [GPa]

Crack length [mm]

Fracture strength [MPa·m0.5]

0 2000 5000 8000 12,000 16,000

84.77 78.15 79.47 75.29 64.40 66.24

807.90 744.26 720.54 761.82 704.43 719.50

1.66 0.83 0.67 0.69 1.54 1.47

0.02 1.08 1.38 1.40 0.14 0.22

the diamond films can induce significant intrinsic compressive residual stress [18]. In addition, the significantly larger thermal expansion coefficient of the RB-SiC substrate (4.6 × 10−6/K, as reported by the supplier of the substrate material), which is several times larger than that of the diamond (0.8 × 10−6/K), also contributes significantly to the compressive residual stress in the UMCD film deposited on the RB-SiC substrate [51]. Although the boron doping technology can provide extra oxygen atoms, the contribution of the high-efficiency-doped boron atoms with a larger covalent radius (0.085 nm) to the residual stress in the film is much more significant than that of the oxygen atoms, inducing sufficient residual tensile stress to neutralize the residual compressive stress. This effect can also explain why there is no clear difference between the residual stress values of the BDMCD films deposited using the boron doping levels of 12,000 and 16,000 ppm because the boron doping concentrations in both films are almost the same.

9

3.1.5. Mechanical properties The effects of the boron doping level on four main mechanical properties, including the nano-hardness, elastic modulus, adhesive strength and fracture strength, are clarified in this research. Firstly, both the nano-hardness and elastic modulus values are directly calculated from the indentation depth-load curve plotted by the in-situ nanomechanical test system, as listed in Table 3, and show that boron doping can decrease the nano-hardness and the elastic modulus of the as-deposited diamond film, undoubtedly due to the formation of more nondiamond impurities. Secondly, the adhesive strength of the diamond film, which is related to film fracture, film removal and the extension of the crack under the Rockwell indentation test, is qualitatively evaluated. Indentation morphologies of the diamond films are shown in Fig. 9. It is apparent that film fracture, film removal and the extension of the crack occur on all the diamond films, but the different diamond films exhibit rather different areas of film fracture and film removal or different lengths of extended cracks, which are unified as a parameter called crack length, as listed in Table 3. This crack length is determined by averaging four values at different points on the film surfaces and is used to estimate the adhesive strength of the diamond film. It is recognized that the longer the extended crack, the poorer the adhesive strength. Boron doping with appropriate levels (2000–8000 ppm) can enhance the adhesive strength between the diamond film and the RB-SiC substrate, and as the boron doping level is increased from 2000 to 8000 ppm, the adhesive strength gradually improves. In contrast, the adhesive strength deteriorates when the boron doping level reaches 12,000 and 16,000 ppm. It is hypothesized that for the RB-SiC substrate with no element

Fig. 9. Indentation morphologies of different diamond-coated specimens with similar film thickness (24.8–26.3 μm).

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detrimental to the film-substrate adhesive strength (such as the Co element in the WC-Co substrate), the film-substrate adhesive strength is mainly affected by the residual stress. In short, the higher the compressive or tensile residual stress, the easier the extension of initial cracks along the interface and the worse the adhesive strength. Consequently, the BDD films fabricated with boron doping levels of 5000 and 8000 ppm exhibit the best adhesive strength. Thirdly, the fracture strength value of the diamond film is calculated based on the crack length presented in the indentation morphology, the indentation parameters (the indentation force and indenter diameter) and the calculated hardness and elastic modulus values. As listed in Table 3, the adhesive strength calculated from the crack length plays the most critical role in the calculation of the as-defined fracture strength, and thus the variation of this fracture strength with the boron doping level is in good agreement with that of the adhesive strength. 3.2. Erosion behavior 3.2.1. Erosive wear process At an impact velocity v of 140 m/s and under an impact angle α of 90°, variations of the erosion rates ε of seven specimens (the uncoated RB-SiC specimen, the UMCD-coated specimen and BDMCD-coated specimens fabricated using five different boron doping levels) with the test time are plotted in Fig. 10. The uncoated RB-SiC specimens are found to exhibit almost stable ε values during the entire solid particle erosion test, but all diamond-coated specimens show four distinctive stages, as exemplified in the inserted figure of Fig. 10, including the initial stage with a relatively high ε (I), the first stable stage with low and almost stable ε (II), the transitional stage between the two stable stages (III) and the second stable stage with a high ε (IV). As indicated by Fig. 5 and Fig. 6, irrespective of whether the diamond grains are well shaped or defective, all the diamond-coated specimens show significantly rough surfaces composed of diamond grains with apparent sharp peaks and edges. It is easily understood that at the initial stage of the erosion test, the impact of solid silica particles on these sharp peaks and edges will provide much higher Hertz contact stress than those on relatively smooth surfaces (such as the uncoated RB-SiC surface) as a result of the small contact area. Therefore, minor cracks are much easier to initiate and propagate in the impacted diamond

grains, and thus, rapid and clear fractures of sharp peaks and edges of the diamond grains naturally occur [52]. In addition, the removal of the extruded peaks and edges of the diamond grains also plays an important role in reducing the surface roughness of the diamond film to some extent. Adopting the BDMCD (2000 ppm and 5000 ppm)-coated specimens as examples, typical wear topographies after stage I are shown in Fig. 11a and Fig. 11b together with the as-measured surface roughness values, which are approximately 56–72% of the initial values before the erosion tests (as attached in the braces), fully demonstrating the wear mode at this stage. With the gradual fractures of the extruded peaks and edges of the diamond grains and the smoothing of the film surface, ε gradually decreases and then remains stable at a low level (stage II). It is recognized that at this stage, the diamond-coated specimen exhibits a stable wear mode, i.e., the relatively slow removal of material from the diamond film that has a comparatively smooth surface formed in the foregoing stage I. As calculated based on the Hertz impact theory, at an impact velocity v of 140 m/s, the depth of the maximum shear strength caused by the impact of a single angular silica solid particle with a diameter of 180 μm is approximately 14 μm (the Poisson ratio of the solid particle νs = 0.227, the elastic modulus of the solid particle Es = 59 GPa, the Poisson ratio of the diamond νs = 0.07, and the elastic modulus of the diamond Ed = 1140 GPa [20]). Moreover, the contact radius (am) is calculated as 31.22 μm, and in the present study, the coating thicknesses (ct) as a proportion of the contact radius, i.e., ct/am, are in the range of 0.79–0.84. The ratio ct/am can be adopted to evaluate whether the film thickness is adequate to protect the substrate at a certain applied stress [53,54]. A ct/am ratio above 0.4 is considered to assure the high performance of the diamond film, and an as-deposited diamond film with a ct/am above 0.5 can be considered a homogeneous material without the influence of the substrate. The above discussions indicate that at the present stage, especially during the first half of the stage, the impact of the silica solid particles on the film surface has not had a significant influence on the film-substrate interface that is far enough away from the film surface (24.8–26.3 μm), and the influence of the film-substrate adhesive strength can be neglected. Therefore, the stable wear mode as mentioned above dominates in this stable erosive wear process of the diamond-coated specimen, which is mainly affected by the excellent intrinsic mechanical properties of the diamond film itself. It is also known that significant features of the stable wear

Fig. 10. Variation of the erosion rates of seven different specimens with the test time at an impact velocity v of 140 m/s and under an impact angle α of 90°.

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X. Wang et al. / Diamond & Related Materials xxx (2016) xxx–xxx

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Fig. 11. Typical wear topographies of as-fabricated diamond-coated specimens: (a) the wear topography of the BDMCD (2000 ppm)-coated specimen after stage I, together with the asmeasured surface roughness values and initial values before the test (in brackets); (b) the wear topography of the BDMCD (5000 ppm)-coated specimen after stage I, together with asmeasured surface roughness values and initial values before the test (in brackets); (c) the newly formed ring crack on the surface of the BDMCD (8000 ppm)-coated specimen during stage II; (d) the completely penetrated ring crack on the surface of the BDMCD (12,000 ppm)-coated specimen during stage II; (e) the large-area removal of the BDMCD (16,000 ppm)-coated specimen during stage III; (f) the rapid wear of the RB-SiC substrate beneath the BDMCD (8000 ppm) film during stages III and IV.

mode are the formation and the penetration of the ring cracks within the impingement area, as demonstrated by Fig. 11c and Fig. 11d. It has been further established by our and other researchers' previous investigations [20,39] that the formation of the ring crack on the film surface is induced by the requisite level of accumulated tensile stress accumulated by continuous impacts because the kinetic energy of one single impact cannot produce a sufficient tensile stress to form a ring crack. Because the formation of a ring crack is extremely rapid and plays only a subsidiary role in the entire erosive wear process of the diamond-coated specimen, this issue is not further addressed in the present study. After the rapid formation of surface ring cracks, the penetration of these cracks, together with the intersection of the penetrated ring cracks and the propagated transverse cracks, act as the main phenomena of this stable wear mode, namely the typical brittle fracture erosive wear, similar to that of the natural diamond. The low and essentially stable ε at stage II indicates slow but regular material loss of the as-deposited diamond film; thus, in the impacted region, the diamond film gradually becomes thinner, and the depth of the maximum shear stress (approximately 14 μm) gradually approaches the film-substrate interface. As a result, irrespective of whether the film-substrate adhesive strength is satisfactory, the film removal is ensured to appear eventually, which is attributed to the accumulation of the shear stress, the formation of minor cracks, and the propagations and connections of minor cracks formed at the film-substrate interface. However, the ultimate cause is the innate weakness of the adhesive strength between two different materials with relatively weak chemical

bonds or mechanical bonding (i.e., the diamond film and RB-SiC substrate). The beginning of stage III is defined as the film lifetime lf because the beginning of this stage signifies the failure of the diamond-coated components under most circumstances. With the persistent development of stage III, the film removal becomes more severe (as shown in Fig. 11e) because the impact of the solid particles affects the filmsubstrate interface much more directly, and thus the propagation and connection of the minor cracks at or close to the interface occur much faster. Moreover, the continuing impact events on the specimen surface where the diamond film has been removed also cause the rapid wear of the exposed RB-SiC material, as demonstrated in Fig. 11f. Due to the two reasons discussed above, the erosion rate ε gradually increases during this stage and reaches a high and stable value at stage IV. It is expected that the second stable erosion rates of all the diamond-coated specimens are approximately equal to the erosion rate of the RB-SiC specimen, fully demonstrating that the material loss of the exposed RB-SiC substrate dominates in stage IV. 3.2.2. First stable erosion rate εfs As discussed above, the first stable erosion rate εfs, i.e., the low and essentially stable erosion rate during stage II, is intimately associated with the intrinsic mechanical properties of the as-deposited diamond film itself. Therefore, this factor is adopted as the first criterion to evaluate the influence of the boron doping level on the erosion behavior of the MCD film. At the impact velocity v of 140 m/s and under the impact angle α of 90°, the εfs values of the as-fabricated diamond-coated

Please cite this article as: X. Wang, et al., Influence of boron doping level on the basic mechanical properties and erosion behavior of boron-doped micro-crystalline diamond (..., Diamond Relat. Mater. (2016), http://dx.doi.org/10.1016/j.diamond.2016.09.025

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X. Wang et al. / Diamond & Related Materials xxx (2016) xxx–xxx

Fig. 12. The first stable erosion rates εfs and film lifetime lf of different diamond-coated specimens with similar film thickness (24.8–26.3 μm) at an impact velocity v of 140 m/s and under an impact angle α of 90°.

specimens, calculated based on Fig. 10, are plotted in Fig. 12. It has been clarified based on Evans and Gulden's erosive wear model that the stable erosion rate of a typical brittle material caused by brittle fracture erosion exhibits an inverse relationship with the hardness and fracture strength of the brittle material, as quantitatively expressed by [55]: V∝v19=6 r 11=3 ρ19=12 K −4=3 H −1=4 c

ð5Þ

where V is the erosive wear volume of the brittle material (the impacted specimen), v the impact velocity, r the size of the solid particle, ρ the density of the solid particle, Kc the fracture strength of the impacted specimen and H the hardness of the impacted specimen. In the present research, the hardness and fracture strength values of different diamond films are listed in Table 3. Note that in this table, the fracture strength is calculated based on the indentation morphology, which mainly reflects the influence of the adhesive strength but not the intrinsic property of the diamond film itself. The differences between the actual fracture strength values of diamond films related to Eq. (5) may not be as large as those presented in Table 3. Moreover, other theories exist that indicate that different influential factors (mechanical properties of the impacted specimen) have a significant effect on the erosion rate of brittle materials, such as the elastic modulus and the bending strength in the fracture model [56] and the density, plasticity and hardness in the low-cycle fatigue theory [39]. Regardless of the ambiguous fracture strength, the variation trend of the first stable erosion rate εfs with the boron doping level is generally similar to the variation trends of the nano-hardness and the elastic modulus. Specifically, the UMCD-coated specimen with extremely high hardness and elastic modulus exhibits the lowest εfs, and the BDMCD-coated specimens fabricated using relatively low boron doping levels (2000–8000 ppm) exhibit slightly higher εfs, while the εfs values of the heavily doped BDMCD-coated specimens (12,000 and 16,000 ppm) are over 1.3 times higher than that of the UMCD one, mainly because of the excessive reduction of the film quality and hardness. In addition, the erosion rates of all the diamond films measured in the present study are found to be similar to those of different diamond materials in many studies and much lower than those of certain other conventional hard materials or IR-transmitting solids [57–60]. 3.2.3. Film lifetime From the standpoint of applications, the film lifetime is the most intuitive criterion for evaluating the erosive wear resistance of the diamond-coated component. The film lifetime of different diamondcoated RB-SiC specimens in the present research is also plotted in Fig. 12. All the coated specimens present film lifetimes no shorter than

150 min because, whether the films are UMCD or BDMCD during the first 150 min of the erosion test, the as-covered diamond films that are thick enough can provide sufficient protection of the filmsubstrate interface and thus the substrate material. At this stage, the diamond-coated specimen can take full advantage of the excellent properties of the diamond film, mainly the high hardness and elastic modulus. However, over the duration of the erosion test, the gradual erosive wear of the diamond film leads to the reduction of the film thickness, and the maximum shear stress occurs close to the filmsubstrate interface, which features relatively weaker bonding forces than those inside the diamond film. Under this condition, the probability of film removal increases substantially, and the time when the erosion rate sharply changes (i.e., the film lifetime lf) mainly depends on the film-substrate adhesive strength. The UMCD film and the extremely heavily doped BDMCD film fabricated with a boron doping level as high as 16,000 ppm have been proven to exhibit very poor adhesive strength and cannot resist the rapid propagation of the as-formed transverse cracks at or close to the filmsubstrate interface. As a result, when the test time exceeds 150 min, both types of films start to be delaminated and removed from the RBSiC substrate. Note that although the first stable erosion rate εfs of the UMCD film is lower than that of any other film, indicating that the thinning of this film is relatively slower and the formation of the transverse cracks at or close to the film-substrate interface may occur later, it still exhibits the shortest lf, fully demonstrating the dominant role played by the adhesive strength in guaranteeing the erosive performance of the diamond film. In comparison, BDMCD films synthesized using moderate boron doping levels (5000 and 8000 ppm) exhibit substantially enhanced film-substrate adhesive strength. Consequently, both of them exhibit a film lifetime as long as 270 min. In addition, the lf of the BDMCD (2000 ppm)-coated specimen is 240 min, and that of the BDMCD (12,000 ppm)-coated one is 180 min. In conclusion, taking both the εfs and lf into consideration, the BDMCD-coated RB-SiC specimens fabricated using moderate boron doping levels (5000 and 8000 ppm) present the best erosion behavior. 3.2.4. Influence of the impact velocity v and impact angle α Variations of the first stable erosion rate εfs and film lifetime lf of the BDMCD (8000 ppm)-coated specimen as a function of the impact velocity v (α = 90°) and the impact angle α (v = 140 m/s) are plotted in Fig. 13. As presented in Eq. (5), Evans and Gulden's erosive wear model, the mentioned fracture model and low-cycle fatigue theory also depict the substantial relationship between the erosion rate and the impact velocity v, which explains the results shown in Fig. 13a, i.e., with increasing v, εfs significantly increases, mainly as a result of the increase in kinetic energy of the solid particles. Correlatively, the film lifetime lf decreases with increasing v. In addition to the BDMCD (8000 ppm)-coated specimen, the UMCD-coated specimen and BDMCD-coated specimens fabricated with other boron doping levels all present similar εfs-v and lf-v relationships, and the uncoated RB-SiC specimen presents a similar ε-v relationship. The detailed relationship between εfs and v has been shown to be an exponential function [61]: εfs ¼ kvn

ð6Þ

where k and n are constants and n is defined as the velocity exponent. Based on the εfs-v curve as plotted in the double logarithm coordinate system in Fig. 13a, the velocity exponent n for the diamondcoated specimen can be easily obtained as the slope of the fitting curve. Under the impact angle α of 90°, the velocity exponents for all the specimens used in the present research are calculated and are as follows: RB-SiC 1.164, UMCD 2.656, 2000 ppm BDMCD 2.914, 5000 ppm BDMCD 2.72, 8000 ppm BDMCD 2.683, 12,000 ppm BDMCD 2.477, and 16,000 ppm BDMCD 2.49. The velocity exponent for the RB-SiC is found to be much lower than those for all the diamond-coated specimens because there is a clear difference between the initial material

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(8000 ppm)-coated specimen when v = 140 m/s and α = 30° (equivalent normal velocity = 70 m/s) is 0.169 mg/kg and is even higher than that when v = 80 m/s and α = 30° (0.121 mg/kg). 4. Conclusions BDMCD films are deposited on RB-SiC substrates using different boron doping levels (2000, 5000, 8000, 12,000 and 15,000 ppm). Application-oriented studies are conducted on UMCD and BDMCDcoated specimens with similar film thickness (24.8–26.3 μm) and on the uncoated specimen, leading to the following conclusions:

Fig. 13. Variations of the first stable erosion rates εfs and film lifetime lf of the BDMCD (8000 ppm)-coated specimen with (a) an impact velocity v (α = 90°) and (b) an impact angle α (v = 140 m/s).

removal mechanisms of the RB-SiC and diamond film at the stable wear stage, while for the various diamond films, the material removal mechanisms are nearly identical. Because the silica solid particles adopted in the present study are much softer than the diamond, sufficient impact cycles are required to accumulate sufficient stress and damage before the diamond material is removed. Increasing the value of v can significantly reduce the required quantity of such cycles, manifested as a large value of n. However, because the hardness of the silica solid particles is close to that of the RB-SiC specimen, the RB-SiC material can be removed by each impact. As a result, the effect of v on the erosion rate is much smaller, manifested as a very small n. Furthermore, the velocity exponents n for the heavily doped BDMCD (12,000 and 16,000 ppm) films are slightly lower than those for other diamond films, which can also be attributed to the relatively lower hardness and elastic modulus. Variations of εfs and lf for the BDMCD (8000 ppm)-coated specimen with the impact angle α (v = 140 m/s) are plotted in Fig. 13b, and all other coated specimens also show similar εfs-α and lf-α relationships. The monotonic decrease in εfs and the prolongation of lf with decreasing α are easily understood as representing the typical behavior of brittle materials. The erosive wear mechanism of brittle materials only minimally changes with α; thus, for smaller α values, the extent of the strength degradation diminishes, mainly because of the reduction of the normal velocity and the normal force. However, the component velocity and force along the in-plane direction also contribute to the material removal of the diamond film; thus, the value of εfs of the BDMCD

(1) Boron doping technology can increase the growth rate of the diamond film; the maximum growth rate can be attained using moderate boron doping levels of 5000 and 8000 ppm. With increasing boron doping, the grain size gradually decreases. The moderate boron doping levels of 5000 and 8000 ppm produce the lowest surface roughness, minimize the residual stress (absolute value) and maximize the film-substrate adhesive strength. However, under the heavily doped conditions (12,000 and 16,000 ppm), such basic mechanical properties of the asdeposited BDMCD films are significantly degraded, mainly as a result of the formation of many more non-diamond impurities and the increase in tensile residual stress. (2) The UMCD film with the highest diamond quality exhibits the highest nano-hardness and elastic modulus, and both mechanical properties are degraded with increasing boron doping. It is of great significance that the moderate boron doping levels can also guarantee the film quality, hardness and elastic modulus to some extent, which, together with the best adhesive strength make the BDMCD films (5000 and 8000 ppm), exhibit highly favorable erosion behavior, including relatively lower stable erosion rates εfs and the longest film lifetime lf. Note that the film lifetime is the most important and intuitive criterion for evaluating the erosive wear resistance of diamond-coated components in actual applications. Nevertheless, for the highest boron doping (12,000 and 15,000 ppm), the excessive reduction of the film quality and hardness and the deterioration of the film-substrate adhesive strength play significant detrimental roles in the erosion behavior of the as-deposited BDMCD films. (3) For either the uncoated or the diamond-coated specimens, all the stable erosion rates monotonically increase with the impact velocity and the impact angle, and the impact velocity dependences of stable erosion rates of the diamond-coated specimens are much more significant than that of the uncoated specimen, as indicated by a higher velocity exponent. For all diamond-coated specimens, increases in either the impact velocity or the angle dramatically shorten the film lifetime.

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Please cite this article as: X. Wang, et al., Influence of boron doping level on the basic mechanical properties and erosion behavior of boron-doped micro-crystalline diamond (..., Diamond Relat. Mater. (2016), http://dx.doi.org/10.1016/j.diamond.2016.09.025