# Creation of color centers in diamond by recoil implantation through
dielectric films

Yuyang Han<EMAIL_ADDRESS>Department of Electrical and Computer Engineering,
University of Washington, Seattle, Washington 98195, USA  Christian Pederson
Department of Physics, University of Washington, Seattle, Washington 98195,
USA  Bethany E. Matthews  Pacific Northwest National Laboratory, Richland,
Washington 99352, USA  Nicholas S. Yama  Department of Electrical and Computer
Engineering, University of Washington, Seattle, Washington 98195, USA  Maxwell
F. Parsons  Department of Electrical and Computer Engineering, University of
Washington, Seattle, Washington 98195, USA  Kai-Mei C. Fu  Department of
Electrical and Computer Engineering, University of Washington, Seattle,
Washington 98195, USA  Department of Physics, University of Washington,
Seattle, Washington 98195, USA  Pacific Northwest National Laboratory,
Richland, Washington 99352, USA

###### Abstract

The need of near-surface color centers in diamond for quantum technologies
motivates the controlled doping of specific extrinsic impurities into the
crystal lattice. Recent experiments have shown that this can be achieved by
momentum transfer from a surface precursor via ion irradiation, an approach
known as “recoil implantation.” Here, we extend this technique to incorporate
dielectric precursors for creating nitrogen-vacancy (NV) and silicon-vacancy
(SiV) centers in diamond. Specifically, we demonstrate that gallium focused-
ion-beam exposure to a thin layer of silicon nitride or silicon dioxide on the
diamond surface results in the introduction of both extrinsic impurities and
carbon vacancies. These defects subsequently give rise to near-surface NV and
SiV centers with desirable optical properties after annealing.

††preprint: AIP/123-QED

In the past two decades, color centers in diamond have gained significant
attention as atomic-scale sensors [1, 2, 3] and as qubits for quantum
information processing [4, 5, 6, 7, 8]. The negatively-charged nitrogen-
vacancy (NV-) center is particularly attractive for sensing due to its high
internal quantum efficiency [9], room-temperature spin initialization [10],
spin readout [11], and long spin coherence time. The negatively charged
silicon-vacancy (SiV-) center shows promise for quantum network applications
due to its stable, uniform, and strong optical transitions [12, 13] that are
further protected from the electric field noise by its inversion symmetry
[14]. However, the SiV- center also has drawbacks such as lower quantum
efficiency and shorter spin coherence time [15]. Consequently, researchers are
exploring alternative color centers in diamond such as other group IV defects,
including germanium-vacancy (GeV), tin-vacancy (SnV), and lead-vacancy (PbV)
centers [16, 17, 18]. To better study and engineer diamond defects for quantum
technologies, means of controllably creating color centers in diamonds are of
great interest. Particularly, the formation of very near-suface defects (i.e.
$<$ 10 nm) is desirable for their potential applications in nanoscale sensing
[19] and near-field photonics device coupling [20] for ultrafast single photon
sources [21].

The creation of these vacancy defect complexes (i.e., X-V centers) requires
the presence of the extrinsic impurity X – either unintentionally or
deliberately doped – and carbon vacancies, which can be introduced by
electron, ion, or laser irradiation [22, 23, 24]. Among the techniques to
create X-V defects, X-ion implantation with post annealing has been widely
used [25, 26, 27, 28, 29], as bombarding the diamond lattice with selected
ions introduces both carbon vacancies and extrinsic impurities. Subsequent
high-temperature anneal (usually above
$800\text{\,}\mathrm{\SIUnitSymbolCelsius}$) allows vacancies to diffuse to
the impurities and form X-V complexes [30].

During implantation, the lateral position of the ions can be controlled either
by customized implantation masks or focused ion beam (FIB) direct writing [31,
26, 32, 33, 34]; the ion depth distribution by the implantation energy; and
the ion density by the implantation dose, screening layers, and ion detection
systems [25, 28, 35, 36]. Near-surface defects, for example, that are 1 - 20
nm below the diamond surface can be created by low energy (0.4 - 3 keV)
nitrogen ion beam implantation [37, 38, 39, 40, 41] and nitrogen delta-doping
[42]. Nevertheless, challenges remain in the creation of defects with sub-30
nm accuracy and close-to-unity formation yield [43]. FIB implantation faces
additional limitation in its ion source availability.

Recently, an alternative technique known as “recoil implantation” [18, 44, 45,
46, 47, 48] was used to address the limitation of ion sources by harnessing
the momentum transfer from the ion-beam to the surface precursors, including
solid-state films (silicon, germanium, tin, and lead) [18] and nitrogen-based
gases (N2, NF3, and NH3) [44]. To form vacancy complexes, a Xe FIB was used to
displace the atoms from the surface precursors into the diamond lattice,
creating both impurities and vacancies that give rise to the corresponding
color centers after annealing.

In this work, we extend the recoil implantation method to the use of
dielectric precursors and gallium FIB. Specifically, we implanted near-surface
nitrogen and silicon atoms by employing gallium FIB irradiation through a thin
layer of either silicon nitride or silicon dioxide. Following high-temperature
anneal and subsequent film removal, the formed NV and SiV centers display
desirable properties. This demonstration offers a more accessible approach to
generating NV centers by recoil implantation, compared to using gas-phase
precursors. In general, the extension of ion-source options and precursor
materials makes this technique attractive for defect engineering.

Figure 1: Experiment overview: (a) illustration of the experiment process. The
size of exposure-squares are not drawn to scale but their distribution on the
sample is to scale. Distances between the
$$10\text{\,}\mathrm{\SIUnitSymbolMicro
m}$\times$10\text{\,}\mathrm{\SIUnitSymbolMicro m}$$ exposure squares in each
round of FIB exposure are $300\text{\,}\mathrm{\SIUnitSymbolMicro m}$. (b)
Representative confocal images from two of the irradiation squares with
different gallium fluence ($P_{5}$ for upper image and $P_{1}$ for lower
image). Scale bars in the images are $20\text{\,}\mathrm{\SIUnitSymbolMicro
m}$. (c) Representative spectra collected at room temperature (upper plot) and
at 10 K (lower plot). The blue line represents data obtained from the highest-
dose square of the SiO2 sample prior to high-temperature annealing; all other
spectra are acquired post-annealing. The red-line spectrum is from a region
between two squares; and the purple line is taken from one of the squares
irradiated with FIB prior to depositing dielectric layer. Important
wavelengths are labeled by the grey dashed lines. In ascending order, these
correspond to the first-order Raman line, NV0 ZPL (575 nm), second-order Raman
line, NV- ZPL (637 nm), SiV- ZPL (737 nm), and GR1 ZPL (741 nm).

Two $2\times 2\times 0.5$ mm3 single crystal type IIa chemical-vapor-
deposition (CVD) diamond samples were used as substrates. The specified
nitrogen impurity concentration is less than 5 ppm (Element Six). Before the
experiments, we cleaned both samples in acid and solvents (see Supplemental
Materials). Next, a control experiment was performed on both samples to
introduce lattice damage in the absence of surface precursors. Each sample was
exposed to 30 keV Ga+ ions in 12 $$10\text{\,}\mathrm{\SIUnitSymbolMicro
m}$\times$10\text{\,}\mathrm{\SIUnitSymbolMicro m}$$ square regions. Six
different FIB parameters were used for these 12 regions, with each parameter
performed twice to test for variance between identical exposure conditions at
different diamond surface locations (FIG. 1. (a)). The irradiation parameters,
from $P_{1}$ to $P_{6}$, are summarized in TABLE 1.

Parameter | I (nA) | dwell ($\text{\,}\mathrm{\SIUnitSymbolMicro s}$) | Pass | Fluence (cm-2)
---|---|---|---|---
$P_{1}$ | 0.79 | 1 | 10 | $(1.30\pm 0.14)\times 10^{16}$
$P_{2}$ | 0.79 | 1 | 2 | $(2.61\pm 0.28)\times 10^{15}$
$P_{3}$ | 0.43 | 1 | 1 | $(1.14\pm 0.15)\times 10^{15}$
$P_{4}$ | 0.08 | 1 | 1 | $(9.7\pm 2.8)\times 10^{14}$
$P_{5}$ | 0.024 | 1 | 1 | $(5.19\pm 0.50)\times 10^{14}$
$P_{6}$ | 0.0077 | 1 | 1 | $(2.87\pm 0.36)\times 10^{14}$
Table 1:  FIB parameters and estimated ion fluence: dwell time refers to how
long the ion beam is at any location during one raster scan; pass represents
the number of exposures performed on the the square region.

The samples were then solvent-cleaned, after which we deposited a 5 nm SiO2
layer by electron-beam evaporation on one of the samples and a 5 nm SiNx layer
by plasma-enhanced CVD (PECVD) on the other. The deposited layers were used as
surface precursors, where the same FIB exposure (TABLE 1) was performed on
them to create 12 new irradiation squares on each sample (Fig. 1(a)). Six
additional lower-dose squares were added to the SiO2 sample to test a wider
range of ion fluences (see Supplemental Materials.)

The samples were cleaned by acid after the second FIB exposure; the acid clean
does not strip the dielectric films [49] but removes contaminants prior to the
oxygen anneal. This anneal, at $460\text{\,}\mathrm{\SIUnitSymbolCelsius}$ for
four hours in an O2 atmosphere, allows the carbon self-interstitials to
diffuse [50], thus healing some of the irradiation-induced damage. However,
the relatively low temperature does not lead to vacancy complex formation
[51]. Hence, this oxygen-anneal is followed by a high-temperature anneal at
$900\text{\,}\mathrm{\SIUnitSymbolCelsius}$ for 2 hours in a vacuum
environment (see Supplemental Materials). 25% hydrofluoric acid (HF) was used
to strip both dielectric layers after annealing. Finally, a second oxygen-
anneal was performed on both samples to remove the irradiation-induced sp2
carbon and oxygen terminate the surface for charge-state conversion to NV-
[52].

After the first oxygen-anneal and before the high-temperature anneal, PL
confocal imaging and spectroscopy of all irradiation squares was measured at
room temperature with 532 nm excitation. Representative confocal images are
displayed in FIG. 1(b), where a visible halo was found outside each of the 3
highest-dose irradiation squares ($P_{1}$,$P_{2}$, and $P_{3}$), attributed to
neutralized gallium ions that are not fully focused. Spectra acquired from
these fluorescent regions confirmed that the fluorescence comes from the GR1
center (neutral monovacancy defect) emission (FIG. 1(c)). Both the PL spectrum
and intensity remained consistent across squares with identical FIB
parameters, implying no intrinsic variation among different locations on the
sample surface.

After the second oxygen-anneal, the PL emission changes from GR1 to NV and
SiV-. As shown in FIG. 1(c), the SiO2 sample squares display a strong SiV-
ZPL, while the SiNx sample squares have strong NV0, NV-, and SiV- ZPL. This
indicates that both the NV and SiV centers were created from the dielectric
films. The control experiment confirms that recoil implantation is the source
of the doped nitrogen and silicon atoms, rather than diffusion of the
dielectric film into the irradiated diamond. Additionally, a broadband
irradiation-induced fluorescence was observed in all squares (FIG. 1(c),
purple line).

Around the highest-fluence irradiation square from each sample, we took a
hyperspectral image (FIG. 2). Comparing the SiNx sample to the SiO2 sample,
the NV ZPL in the former is notably brighter within the irradiated region,
while the latter has a more homogeneous intensity distribution. This confirms
that the NV centers in the SiNx sample arise from the nitrogen-based
precursor, rather than from the naturally occurring impurities as in the SiO2
sample.

Figure 2: Hyperspectral images of the highest-dose irradiation squares from
the SiO2 sample and SiNx sample. From left to right, the intensity plots are
obtained by: (All) integrating the total PL intensity over the 625 nm-792 nm
confocal bandwidth; (SiV) binning the SiV- ZPL; and (NV) binning the NV- ZPL.
Binning is performed by fitting and integrating over the corresponding ZPL
(see Supplemental Material). The excitation power was 4.2 mW. The plots are
not normalized to the same intensity. The irradiation-square boarder is marked
by the white dashed line in each plot. The scale bars in the plots are
$10\text{\,}\mathrm{\SIUnitSymbolMicro m}$. Figure 3: Formation yield
estimate: TRIM simulations of 30 keV Ga+ incident ion beam on diamond with (a)
5 nm electron-evaporation SiO2 (2.658 g/cm3) and (b) 5 nm PECVD SiNx (2.5
g/cm3). To derive the effective implantation dose of nitrogen and silicon
atoms, integration over the simulation depth was performed, followed by
multiplication with the gallium ion dose. This resulted in 1.83 silicon atom
per gallium ion for the SiO2 sample and 2.18 silicon plus 2.19 nitrogen atoms
per gallium ion for the SiNx sample. (c) Confocal PL intensities for all
irradiation squares measured at 0.2 mW 532 nm excitation. The vertical width
of each data point represents the standard deviation of PL intensities inside
the square; the horizontal width represents the fluence uncertainty. (d)
Formation yield for SiV and NV centers from the SiN and SiO samples that is
computed based on (a), (b), and (c). Figure 4: Defect characterization: (a) 10
K PL spectrum of NV- centers from the SiNx sample. The spectrum was fit with a
Lorentzian function to determine a FWHM of 535 GHz. (b) 10 K PL spectra of
SiV- center(s) from both samples. The spectrum from an irradiation square on
the SiNx sample (labeled “SiN Square”) was fit by eight Lorentzian functions;
the spectrum from a square on the SiO2 sample (labeled “SiO Square”) and the
spectrum from the halo region around this square (labeled “SiO Halo”) were
each Ffit with four Lorentzian functions. The spectra shown in (b) were scaled
and shifted for better visualization. (c) representative ODMR spectrum from NV
centers in the halo region of highest-fluence square on the SiNx sample. The
same experiments were performed when the excitation beam is inside the square,
$10\text{\,}\mathrm{\SIUnitSymbolMicro m}$,
$20\text{\,}\mathrm{\SIUnitSymbolMicro m}$, and
$30\text{\,}\mathrm{\SIUnitSymbolMicro m}$ away from the center of the square.
The first two resonances from these measurements are shown in (d). The
asymmetric spectrum intensity comes from the finite bandwidth of our microwave
antenna (Supplemental Material).

To calculate the formation yield of NV and SiV centers, we first simulated the
nitrogen and silicon atom distribution by TRIM [53] and obtained the number of
recoil-implanted atoms per gallium ion (FIG. 3(a) and (b)). Relative
contributions to the total PL intensity (FIG. 3(c)) from the different color
centers were calculated by fitting the PL spectra with a weighted sum of NV0
spectrum, NV- spectrum, SiV- spectrum, and the “irradiation-induced
fluorescence” spectrum (FIG. 1(d), purple line). After normalizing the
respective PL intensity to that of a sample with known defect density, we
derived the formation yield (FIG. 3(d)) for NV and SiV centers from the two
samples (see Supplemental Materials). The maximum yield is on the order of
$10^{-3}\%$, much lower than 0.5% - 50% that is obtained by direct ion
implantation [28, 26, 54, 55, 56, 35, 25]. However, this is expected as the
implanted atoms are not only very close to the surface but also in regions of
very high ion fluence. Both factors have the effect of reducing the PL
intensity [57] and decreasing the formation yield [26, 37]. In particular, the
yield of the SiV centers is higher in the SiNx sample, which is likely because
the SiV2- (dark state) can capture holes that are photo-generated from the
nearby NV0 centers and become SiV- (bright state) [58]. The PL intensities in
FIG. 3(c) remained “saturated” for the 6 highest-dose squares. This implies
that we used excessive ion fluences for all of the SiNx squares. However, it
is still possible to probe lower-dose conditions by taking measurements in the
halo region, where the fluence of the neutralized gallium ions should be much
lower.

To characterize the optical properties of the formed defects, we measured the
PL spectra from all irradiation squares at 10 K. Most SiV centers from the
squares exhibit four distinct main transition peaks, whereas the NV centers
show large inhomogeneous broadening (FIG. 4(a) and (b)). Among the SiV-
centers, we found the smallest inhomogeneous broadening (FWHM) of 16 GHz from
the halo region of the highest-dose square in the SiO2 sample. This
inhomogeneous linewidth is close to the spectrometer limit (10 GHz), and
similar linewidths from several other irradiation squares are all comparable
to the 15 GHz linewidth observed in the literature by direct silicon ion
implantation [59]. However, the recoil-implanted SiV centers are 10 times
closer to the diamond surface (less than 10 nm compared to around 100 nm),
which typically leads to spectral diffusion due to the unfavorable surface
strain and charge environment [60].

We further investigated the spin coherence of the NV centers from the highest-
dose square on the SiNx sample, by conducting continuous wave optically
detected magnetic resonance (CW-ODMR). In the presence of a 24 G external
magnetic field, we observed four pairs of resonance dips that correspond to
the 4 different crystallographic orientations of NV centers (FIG. 4(c)). As
the excitation beam moved from inside this square to the halo region, the
maximum PL contrast for a single resonance dip increased from 0.47% to 2.94%,
while the average linewidth (FWHM) decreased from 10.3 MHz to 8.37 MHz (FIG.
4(d)). This suggests that the spin-coherence of the NV centers will benefit
from lower ion fluences than what we used in this study. Additionally, these
linewidths are broadened by the 2 mW excitation power [61]; we obtained the
narrowest spectrum with a linewidth of 6.81 MHz with 0.1 mW excitation power
(see Supplemental Material). The observed 2.94% contrast is approximately half
of the expected maximum. This reduced contrast is due to the lower NV-/NV0
ratio and the additional background PL from SiV- centers (Supplemental
Material). The reduction in contrast served as a trade-off for having NV
sensors at these shallow depths and may be improved by additional surface
preparations

Additional TRIM simulations were performed for the SiO2 precursor to explore
the flexibility provided by recoil implantation. As shown in Fig. 5(a) and
(c), both the carbon vacancy and the silicon atom distributions can be
adjusted by varying the implantation ion energy and the SiO2 layer thickness.
This semi-independent control is not possible by direct silicon ion
implantation, where the vacancy and silicon atom ratio is set by the
implantation energy only. Potentially, this new degree of freedom allows for
forming better defects by controlling the distribution of the gallium
impurities. In particular, Fig. 5(d) shows that using a very thin layer of
SiO2 achieves a much smaller range of depth for the silicon atoms, while
sending the Ga+ ions to a different depth from these silicon atoms. This opens
the possibility of forming a 2-d layer of defects on the crystal surface. In
this case, the lateral straggling is expected to be small, which indicates
that the spatial resolution is only limited by the ion beam-profile.

Figure 5: TRIM simulations of Ga+ ion irradiation on diamond with a SiO2 layer
on surface. (a) 5 nm SiO2 layer with variable ion energies. From small to
large the energies are 10 keV, 30 keV, 50 keV, 70 keV, and 90 keV. (c) is the
zoom-in of (a). (b) 30 keV ion energy with variable SiO2 layer thickness. From
small to large the thicknesses are 1 nm, 5 nm, 10 nm, 15 nm, 20 nm, 25 nm, and
30 nm. (d) is the zoom-in of (b.1). The vertical (z) axes of these plots are
density per dose with unit cm${}^{-3}/$cm-2.

To conclude, we demonstrate gallium FIB recoil implantation through dielectric
precursors for generating color centers in diamond. The use of silicon-dioxide
and silicon-nitride films in this experiment expands the range of applicable
surface precursors to include dielectric materials, exemplified by the
formation of near-surface SiV and NV centers. These SiV centers exhibit narrow
spectral linewidths that are promising for performing photoluminescence
excitation and photonics device coupling, while the NV centers display clear
ODMR spectra, highlighting their efficacy in quantum sensing applications.
Simulations also underscore the flexibility of this technique, especially the
degrees of freedom to fine-tune the the spatial distributions and density
profiles of the implanted atoms.

This material is based upon work supported by Department of Energy, Office of
Science, National Quantum Information Science Research Centers, Co-design
Center for Quantum Advantage (C2QA), under contract number DE-SC0012704. The
FIB implantation was performed at the Pacific Northwest National Laboratory.
The dielectric film deposition was performed at the Washington Nanofabrication
Facility, a National Nanotechnology Coordinated Infrastructure (NNCI) site at
the University of Washington with partial support from the National Science
Foundation via awards NNCI-1542101 and NNCI-2025489. Optical and ODMR
measurements was performed in the Quantum Technologies Training and Testbed
(QT3) user facility, supported by the National Science Foundation under award
OMA-1936100 and OMA-1936932. The authors would like to thank Adam Cox and
Chukwuemeka Mordi for their contributions in developing the hyperspectral
imaging scrips.

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