Enhancing surface production of negative ions using nitrogen doped diamond in a deuterium plasma

A deuterium plasma, figure 1(a), is produced via a helicon source operated in capacitive mode (2 Pa, 26 W), which then expands into a diffusion chamber [54]. The pressure of the diffusion chamber, as measured by a Baratron gauge (MKS), is regulated via a mass flow controller (7.6 sccm, BROOKS 5850TR) in combination with a 150 mm diameter Riber gate valve installed in front of a turbo molecular pump (Alcatel ATP400). To reduce experimental drifts, the experiment source chamber and lower spherical diffusion chamber have a base pressure of 10⁻⁵ Pa, which is lower than the base pressure of a previous setup of 10⁻⁴ Pa [54].

The relatively low power coupled to the plasma source results in plasma densities of approximately 10¹⁴ m⁻³ in the spherical diffusion chamber [54]. The choice of power and pressure was for similarity with previous work [34]. The positive ion composition of the deuterium plasma is measured by the mass spectrometer, described below, to be (84 ± 2) % D$_3^{+}$ ions, (14 ± 2) % D$_2^{+}$ ions and (1.1 ± 0.2) % D⁺ ions. The measurement uncertainty represents the day-to-day variation of the measured plasma composition, however the actual error in the plasma composition due to the internal settings of the mass spectrometer may be higher [55].

The sample holder is shown in figure 1(b). It is attached to a DC voltage source (Equipment Scientific Alimentation de Laboratoire CN7C) that can negatively bias the frame that holds the diamond sample. The voltage applied to the sample is defined as V_DC, which is distinct from the voltage at the sample surface, V_S. The sample is positioned 37 mm away from, and perpendicular to, the plane of the mass spectrometer orifice. This is the closest distance that the sample can be placed in front of the mass spectrometer orifice, and is assumed to be sufficiently small to achieve minimal negative ion signal loss. It has previously been demonstrated that this distance has negligible effect on the shape of the negative ion distributions measured by the mass spectrometer [28, 33].

It is worth noting that the angular dependence of the NIEDFs for carbon materials has previously been shown to be similar [33, 35, 36]. Therefore, a single measurement can be used to compare between samples. A misalignment of the sample surface normal to the mass spectrometer would produce spurious results, so to prevent this, the alignment is regularly checked by rotating the sample and maximising the negative ion signal.

As shown in figure 1(b), a tungsten heating element is built into the sample holder, which is used to heat the back of the sample. The heating element is controlled by a PID controller (designed and built by AXESS tech) using a K-type thermocouple inside the the frame of the sample holder. By fixing a second thermocouple to the surface of the sample, its temperature is calibrated against the temperature measured by the PID. The heating element behind the sample increases the temperature of the sample's surface up to (750± 20) °C.

2.3.1. Mass spectrometer setup.

2.3.2. Procedure for measurement of negative ion energy distribution functions (NIEDFs).

2.4.1. Micro-crystalline boron doped diamond and micro-crystalline diamond.

2.4.2. Micro-crystalline nitrogen doped diamond.

The samples were analysed using confocal microscopy and Raman spectroscopy prior to plasma exposure to characterise their properties.

2.5.1. Surface morphology and crystal structure.

A laser confocal microscope (S neox, Sensofar) was used to observe the diamond surface morphology as shown in figure 2. From visual inspection, the crystal grains are observed to have grown to exhibit (111) crystal faces, (100) crystal faces, or a mixture of both, dependent on the concentration of nitrogen dopant introduced in the gas phase during sample growth [60]. As the gas phase nitrogen concentration is increased from 0 ppm to 50 ppm the crystals are observed to exhibit an increased proportion of (111) faces, with predominantly (111) faces observed at 50 ppm. As distinct from these, the crystal grains of the diamond film with 200 ppm gas phase doping displays predominantly (100) faces. The diamond films grown with 200 ppm nitrogen concentration in the gas phase are different to the other samples, due to be large crystals interspersed with regions of what appears to be much smaller crystals with a less pronounced crystal orientation. The size of the crystals for the 0 ppm sample, figure 2(a), are approximately 10 µm, whilst the 10 ppm sample, figure 2(b), has a much smaller average crystal size, at approximately 1 µm. The 50 ppm sample, figure 2(c), has a crystal size similar to the 0 ppm sample, at approximately 10 µm. The 200 ppm sample, figure 2(d), as previously described, appears to have a distribution of large crystals separated by smaller crystals, here the average crystal size of the larger crystals is approximately 5 µm.

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2.5.2. Measurement of the relative quantity of nitrogen doping within the films.

Raman spectroscopy was undertaken to measure the relative concentration of nitrogen dopant that is introduced into the MCD films with respect to the gas phase nitrogen concentration present during the PECVD process. Raman spectra were generated using a Horiba Jobin Yvon HR800 setup. The measurements were undertaken in air using an excitation wavelength of $\lambda_{\rm L}$ = 514 nm, $\times$ 100 objective (numerical aperture of 0.9, i.e. theoretical spot radius of 0.34 µm), 600 grooves/mm grating (resulting in a resolution of about 1 cm⁻¹), and 5 mW laser power. Measurements were taken using 5 acquisitions of 1 second intervals over a range of 0 cm⁻¹ to 3000 cm⁻¹, of which 750 cm⁻¹ to 2500 cm⁻¹ is presented, with the measurement taken at the centre of a dominant crystal near the centre of the sample. A brief optical microscopic inspection of the samples prior to taking Raman measurements showed relatively uniform crystal distribution across the surface of all the samples and allowed for precise targeting of a dominant crystal near to the centre of the sample to be the subject of the Raman measurement. A quantitative comparison between Raman spectra has not been carried out as this requires the measurement of the polarisation of the Raman emission and a good understanding of the grain orientation [61].

Figure 3 shows the Raman spectra from samples of nitrogen doped diamond with gas phase nitrogen doping of 0 ppm to 200 ppm. The spectra have been presented with the background fluorescence removed and normalised to the 1333 cm⁻¹ peak, which can be attributed to sp3 bonded carbon (the diamond bond of carbon) [47, 62]. The normalisation to this peak is justified due to the transparency of the diamond films such that the measurement is integrated across the sample thickness. The normalised spectra can therefore enable a comparison between samples that accounts for any change in the thickness of the MCNDD film [62].

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The broad peak centred at 2100 cm⁻¹ observed in figure 3 can be attributed to nitrogen vacancy centres (NV⁰) that have been introduced into the diamond [63]. This broad peak appears, not due to vibrational modes, but due to the electronic signature attributed to nitrogen vacancy centres and in reality lies at an energy level of 2.15 eV. As Stokes Raman spectroscopy is energy loss spectroscopy, this peak appears arbitrarily at 2100 cm⁻¹ when using a 514 nm laser. Using another laser to perform the Raman spectroscopy results in a change in the wavenumber of this peak [64].

As the measurement configuration is the same for all samples, a relative comparison of the number of nitrogen centres in the diamond can be made using the broad 2100 cm⁻¹ peak. This can then be used to infer relative nitrogen concentration [65]. As shown in figure 3, the ratio of the NV⁰ peak to the peak centred at 1333 cm⁻¹ increases with increasing gas phase dopant concentration, for samples 0 ppm to 50 ppm (200 ppm will be discussed below). This is consistent with previous work, which showed a similar increase in the magnitude of the NV⁰ characteristic peak with an increase in the gas phase nitrogen doping [66].

In figure 3, the Raman spectrum of the 200 ppm nitrogen doped diamond has a peak at 1500 cm⁻¹ that has a much higher intensity than the other samples. This peak is of particular interest as it is associated with the sp2 bond of carbon that has previously been associated to graphite-like bonds [62]. The ratio of the peaks at 1333 cm⁻¹ and 1500 cm⁻¹ implies that there is a higher ratio of graphite in the 200 ppm diamond film compared to the other samples [28, 67, 68]. The 200 ppm nitrogen doped diamond sample also exhibits a NV⁰ centre peak at 2100 cm⁻¹, which is slightly lower than the 50 ppm sample, suggesting a reduction in the number of nitrogen vacancies, and therefore, a reduction in the concentration of nitrogen in the diamond.

The observed increase in intensity of the NV⁰ peaks, increasing from 0 ppm to 50 ppm, may be attributed to both the increase in nitrogen introduced in the gas phase and by a change in the crystal face from a mix of (100) and (111), figure 2(a) and (b), to a primarily (111) face for which impurity incorporation is higher than that for (100) crystals, figure 2(c). This is consistent with the results of previous work [66, 69, 70] This same process may then account for the slightly lower 2100 cm⁻¹ peak for the 200 ppm sample compared to the 50 ppm samples despite a four fold increase in the nitrogen gas phase content. This decrease could be attributed to the change in the crystal orientation (see figure 2(c) and (d)), from a (111) dominant crystal surface for the 50 ppm sample to predominantly (100) crystal orientation for the 200 ppm sample.

The surface characterisation of the samples show that the incorporation of nitrogen into the PECVD process has multiple effects on the diamond produced, aside from only substitutional or interstitial incorporation of nitrogen into the diamond lattice. For these samples, separating the difference in negative ion yield due to the influence of the crystal face or the nitrogen content in the diamond is not possible because of the interrelated nature the presence of nitrogen in the gas phase has with the crystal face orientation and the measurable number of nitrogen vacancy centres. This is an active area of research [71]. However for this study, as nitrogen doping is the main influencing factor that generates the differences between the samples, it is reasonable to suggest that it is possible to associate the nitrogen gas phase doping with the negative ion yield and this is how the samples will be defined in the next section.

Figure 4 presents the negative ion yield from MCNDD for different dopant concentrations, as measured in the gas phase during sample preparation. In both figures 4(a) and (b) the yield profile for MCNDD has a distinct shape. For example, at 50 ppm, the measured yield is practically zero between a temperature of 30 °C and 400 °C. At 450 °C, the yield rapidly increases by several orders of magnitude to a maximum at 550 °C. This transition is similar for all nitrogen doped samples with the transition occurring at temperatures ranging from 250 °C to 450 °C. This is somewhat unlike MCD which produces measurable negative ions for all temperatures. The trend for MCD is a gradual increase to a maximum yield at a temperature of 450 °C (−130 V, high energy bombardment) or 400 °C (−20 V, low energy bombardment), there is then a decrease in yield from this maximum yield as the temperature is increased further to the maximum temperature of 750 °C. It can also be observed that between ∼ 30 °C and ∼ 150 °C, MCD does undergo a transition, though this is smaller than that seen for MCNDD.

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In order to understand these trends, it is important to note that the negative ion yield measured using the technique described in this article relies on a conductive sample surface. A non-conductive sample would not allow negative ions to be accelerated to the mass spectrometer at an energy which the mass spectrometer is tuned for. The magnitude of the transitions seen in figure 4 is a feature of this experimental technique which inadvertently highlights the temperature at which samples become conductive, such that V_DC = V_S.

The difference in trends between MCNDD samples and the un-doped diamond can be attributed to the differences in conductivity between MCNDD and MCD. Previous work with MCD has shown that it has poor conductivity close to room temperature, which explains the increase in yield occurring between 30 °C and 150 °C [26]. Regarding MCNDD, the temperature at which the sharp increase in yield occurs appears to be dependent on the nitrogen doping of the diamond sample. The results of previous work suggest that the level of interstitial nitrogen doping influences the conductivity of diamond [49, 72, 73] and this sharp increase is consistent with increasing nitrogen dopant concentration and its influence on the conductivity of the diamond, supporting the argument that the nitrogen incorporated into the diamond increases as gas phase nitrogen is increased during its production. An exception to this trend are the results for 200 ppm in both figure 4(a) and (b), which does not exhibit a significant increase in the sample temperature for which the film becomes conductive relative to the 50 ppm MCNDD sample. This can be explained by considering figure 3. As discussed in section 2.5.2, the nitrogen content measured using Raman spectroscopy suggests a nitrogen content that is similar for 200 ppm and 50 ppm MCNDD samples meaning, in the absence of other influences, a similar conductivity for these two samples could reasonably be expected.

The maximum yield from each sample occurs at temperatures between 400 °C and 550 °C, which is highlighted in the insets of figures 4(a) and (b). For the 0 ppm and 10 ppm samples, the maximum yield occurs at 400 °C, whilst for 20 ppm (and 200 ppm) it occurs at 500 °C and for 50 ppm, at 550 °C. As mentioned previously, a conductive sample surface is necessary to hold a DC surface bias which is necessary for the acceleration of negative ions into the mass spectrometer. The trend of increasing temperature for maximum yield as dopant increases (excluding 200 ppm) could be related to the maximum yield in these experimental conditions being restricted by the conductivity of the samples. For example, the maximum yield for MCD and MCNDD (10 ppm) is at the peak of a gradual increase and decrease in yield as temperature is increased from ∼ 30 °C to ∼ 400 °C and then from ∼ 400 °C to ∼ 750 °C respectively. This is most clearly observed in figure 4(b) for the 10 ppm sample. This sample is distinct from the other MCNDD samples as the 10 ppm sample exhibits an increase in yield as the temperature is increased (due to a change in conductivity) then a further smaller increase up to a maximum negative ion yield at ∼ 400 °C. The yield then gradually decreases as the temperature is increased further. The other MCNDD samples also undergo an increase in yield due to a change in conductivity, but no further gradual increase in yield is observed as temperature is increased.

It could therefore be reasonable to suggest that the peak yield conditions are not observed due to a lack of conductivity for samples with more than 20 ppm gas phase nitrogen doping. The trends observed for the 0 ppm and 10 ppm samples suggest that a temperature of approximately 400 °C may be the temperature at which these MCNDD with more than 20 ppm gas phase doping produce the highest yield. A technique to measure negative ions that does not require a conductive surface would be necessary to explore this further.

In figure 4(a), the effect of the nitrogen doping on the maximum yield is not readily observed when a bias voltage V_DC = −130 V is applied to the sample. This is unlike figure 4(b) in which a bias voltage of V_DC = −20 V is used. In this data there is an observed difference between the nitrogen doped and non-doped diamond. The yield in figure 4(a) for the MCNDD samples and MCD samples is also lower than the yields from all of the samples in figure 4(b). A higher bombardment energy as a result of the high magnitude bias is associated with an increase in sp2 bond formation in diamond [36]. It is reasonable to suggest that the reduction in yield for the higher magnitude bias creates more sp2 defects which decreases the yield. Additionally, if the yield is not changing with the addition of nitrogen to the diamond, it is also reasonable to suggest that the nitrogen doped diamond may be more susceptible to defect formation due to high energy bombardment which would result in a surface state that does not enhance the negative ion yield through nitrogen doping. Additional work would be necessary to characterise this process.

The apparent influence of nitrogen doping on the measured negative ion yield is observed in figure 4(b) where a sample bias voltage V_DC = −20 V is applied. When comparing the yield at temperatures above 550 °C, i.e. when all of the MCNDD films are conductive, it is observed that the negative ion yield is higher at similar temperatures, when increasing nitrogen dopant concentration for the 0 ppm to 50 ppm cases. The mechanism for such an increase in yield is not immediately clear and future work will be necessary to identify the specific cause of this increase. For example, it could be solely due to interstitial nitrogen, or a change of crystal orientation or a combination of both. In any case, the increase is correlated to the amount of nitrogen dopant with the exception of the result observed for 200 ppm gas phase doping, which produces a comparatively lower yield compared to the 50 ppm case. Should interstitial nitrogen content be the main cause of an increase in negative ion yield, this result can be explained by the Raman measurements shown in figure 3. As discussed in section 2.5.2, the Raman measurement suggests that the diamond has a similar amount of nitrogen doping when comparing the 2100 cm⁻¹ peak for the 50 ppm and 200 ppm samples. However the Raman measurement also suggests that the 200 ppm MCNDD sample has more carbon sp2 bonds (graphite-like) than the 50 ppm sample. The reduction in yield observed for the MCNDD (200 ppm) sample compared to MCNDD (50 ppm) sample is therefore consistent with previous work, which observed that an increased number of sp2 bonds is less favourable to negative ion production [28, 33]. This work suggests that this is still the case with nitrogen doped diamond samples.

NIEDFs for MCBDD and MCNDD are presented in figure 5 to compare negative ion production processes between MCBDD, a previously studied material [26], and the MCNDD samples. In this figure the NIEDFs are normalised to the modal negative ion energy at temperatures where MCBDD and MCNDD samples are conductive.

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When considering a normalised NIEDF, a reduction in the proportional magnitude of the NIEDF peak at low energies will result in an increase in the apparent proportion of negative ions at high energies. The NIEDFs in figure 5(a), show that for MCBDD the proportion of high energy ions increases as the surface temperature increases. This is because the main contribution to the measured yield is low energy ions, which are predominantly created through the sputtering process, as distinct from backscattering, due to the acceptance angle of the mass spectrometer [26, 54]. Previous work has confirmed this interpretation through comparison of experimental results with SRIM simulations [36]. The physical interpretation for the decrease in the sputtering contribution is that this is due to a decrease in the amount of sub-surface deuterium available for sputtering as a result of out-gassing caused by the increase in temperature. [35, 36].

At a surface bias of V_S = −20 V, i.e. the 'low energy' bombardment condition described in section 2.3, the high energy tail observed in the NIEDFs in figure 5 is not produced. Without a high energy tail, the normalised NIEDF shapes are not strongly dependent on the deuterium surface content [33], and comparison of the ratio of sputtered to backscattered particles cannot be readily inferred using this approach. For this reason, only results with a surface bias of V_S = −130 V are presented.

Comparing figure 5(a) to figure 5(b), which presents NIEDFs for MCNDD at 550 °C and 750 °C, i.e. temperatures at which the sample is conductive, it is observed that MCNDD displays a similar increase in the proportion of high energy negative ions as the sample's surface temperature increases. This implies that MCNDD has similar negative ion production properties to MCBDD.

The trends for MCNDD and MCD observed in figure 4 and discussed in the previous section can be explored in the context of figure 5. Figure 4 shows that the yield for MCD increases up to sample temperatures of ∼ 400 °C and decreases as its temperature is increased further. This is similar to the trends observed for samples of MCNDD when they are conductive. The increase and then decrease in yield as temperature is increased, from ∼ 30 °C to ∼ 400 °C and then from ∼ 400 °C to ∼ 750 °C respectively, can be attributed to two processes that combine to generate the observed trend in figure 4. The first process is the removal of defects on the sample surface. The heating of the sample results in an enhancement of the etching of sp2 bonds created by the bombarding positive ions resulting in a surface which results in a higher ratio of sp3 bonds [28]. The increased proportion of diamond bonds on the surface increases the negative ion yield, as explored in previous work through Raman spectroscopy [26, 28, 32, 33]. The second process is the previously discussed decrease in the sputtering contribution to the negative ion yield due to out-gassing of deuterium from the sample surface, as observed in the measurements of figure 5. As temperature is increased, the influence of each of these processes on the measured negative ion yield is observed to vary significantly. At temperatures below ∼ 400 °C, the reduction in defects increases the yield, whilst the outgassing does not cause a significant decrease in the sputtering contribution. At temperatures above ∼ 400 °C, the decrease in sputtering contribution reduces the yield by a greater extent than the reduction in defects caused by the elevated temperature, causing a reduction in the the measured negative ion yield [28].

For the samples of nitrogen doped diamond with more than 20 ppm nitrogen added in the gas phase, the MCNDD film is not conductive at temperatures where the previously mentioned reduction in the defects can increase yield, i.e. between 30 °C and 400 °C. A more thorough exploration of the resulting interplay between the reduction of defects and the decreasing sputtering contribution is beyond the scope of this experimental study. However, figure 5(b) suggests that the decrease in yield due to a decrease in the sputtering contribution is consistent with current understanding of the behaviour of negative ion formation on micro-crystalline doped diamond.

The negative ion yield with respect to sample surface temperature of the MCD, MCBDD and MCNDD samples is shown in figure 6, with high energy ion bombardment (V_DC = −130 V) shown in figure 6(a) and low energy bombardment (V_DC = −20 V) shown in figure 6(b). 50 ppm MCNDD is chosen as a comparison to MCD and MCBDD as this produced the highest relative negative ion yield of all the nitrogen doped diamond samples, as shown in figure 4.

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In figure 6(a), V_DC = −130 V, the trends for MCD and MCBDD are similar, with an increase in yield by a factor of 6 from 150 °C to 450 °C observed for MCD and a factor of 2 observed for MCBDD from 150 °C to 550 °C. The yield then decreases gradually as temperature is increased further. These results are consistent with previous work using MCD and MCBDD citeCartry12017. The negative ion yield from MCNDD at sample temperatures below 400 °C is effectively zero. After 400 °C there is a rapid increase in yield by several orders of magnitude up to 550 °C, as discussed in section 3.1. After 550 °C, the trend agrees with MCD and MCBDD. In the high energy bombardment regime, the yield from MCNDD is found to be lower than MCBDD and comparable to MCD. This suggests that the higher positive ion bombardment energy is having a larger influence on MCNDD than MCBDD, though a mechanism for such a difference is beyond the scope of this study.

In figure 6(b) for V_DC = −20 V, the trends for MCD and MCBDD are also observed to be qualitatively similar, showing an increase in yield by a factor of 2 and a factor 1.5 from 150 °C to 400 °C respectively, and a gradual decrease in yield above 400 °C, which has been discussed in section 3.1 [26, 33]. Figure 6(b) has a similar trend as figure 6(a) where the yield from MCNDD at temperatures below 400 °C is effectively zero. The yield increases by several orders of magnitude between 400 °C to 550 °C, after which it decreases gradually. At temperatures above 550 °C the general trend of decreasing yield is consistent with both MCD and MCBDD, and agrees with current understanding of these diamond films as discussed in the previous section. Of particular interest is that the yield for MCNDD in this low energy ion bombardment condition is observed to be higher than MCD, and also higher than the previously best performing type of diamond, MCBDD [26]. At 550 °C, the maximum yield observed, MCNDD has a higher negative ion yield than MCD and MCBDD by a factor of 2 and 1.5, respectively. This therefore suggests that controlled addition of nitrogen during the growth of diamond using the PECVD process could be an avenue for increasing the negative ion yield from diamond.