In-situ doping of erbium in hydrogenated amorphous carbon by low
temperature metalorganic radio frequency plasma enhanced chemical
vapor deposition
Hui-Lin Hsu
a,⁎
, Keith R. Leong
a, Michael Halamicek
a, I-Ju Teng
b,c, Pratish Mahtani
a, Jenh-Yih Juang
b,c,
Sheng-Rui Jian
d, Li Qian
a, Nazir P. Kherani
a,e,⁎⁎
a
Department of Electrical and Computer Engineering, University of Toronto, ON M5S 3G4, Canada b
Centre for Interdisciplinary Science, National Chiao Tung University, Hsinchu 30010, Taiwan c
Department of Electrophysics, National Chiao Tung University, Hsinchu 30010, Taiwan dDepartment of Materials Science and Engineering, I-Shou University, Kaohsiung 84001, Taiwan e
Department of Materials Science and Engineering, University of Toronto, ON M5S 3E4, Canada
a b s t r a c t
a r t i c l e i n f o
Available online 15 February 2014 Keywords:
Erbium metalorganic compound Hydrogenated amorphous carbon (a-C:H) Fluorination
A significant improvement in the photoluminescence of erbium doped amorphous carbon (a-C:H(Er)) is reported. The effects of the RF power on the anode and cathode a-C:Hfilms were investigated in terms of the microstructural and local bonding features. It was determined that Er doped a-C:Hfilms should be placed on the anode to obtain wider bandgap and lower percentage of sp2carbon bonding. The metalorganic compound,
tris(6,6,7,7,8,8,8-heptafluoro-2,2-dimethyl-3,5-octanedionate) Erbium(+III) or Er(fod)3, was incorporated in-situ into an a-C:H
host by metalorganic rf plasma enhanced chemical vapor deposition. This technique provides the capability of doping Er in a vertically uniform profile. The high erbium concentration (3.9 at.%), partial fluorination of the surrounding ligands, and the large optical bandgap of the host a-C:H are the primary factors that enable enhancement of the photoluminescence.
© 2014 Elsevier B.V. All rights reserved.
1. Introduction
The metal-line based electrical interconnections represent the most important limitation on the performance of Si-based microelectronic (CMOS) devices. The delay in the signal propagation arises from several factors including the parasitic capacitances generated at the metal/ insulator/metal capacitors, the intrinsic resistivity of the metal lines, and the contact resistance at the metal/metal interface. As device features fur-ther shrink, the delay due to the metal interconnection will lead to an un-acceptable bottleneck in device performance[1]. A definitive solution is to employ optical interconnects that are able to transfer data at rates that are orders of magnitude above the limit of traditional electronic technologies. These optical interconnects can be inside a Si chip or between chip-to-chip communication. In order to completely avail optical technologies, it is imperative to develop silicon compatible materials which enable light generation, guiding, switching, detection, modulation and amplification. To realize the co-existence of electrical and optical functions inside a Si chip platform, it is crucial to develop compatible photonic materials.
These materials must possess processing temperatures below 400 °C, in order to meet the Si back end-of-line (BEOL) requirements for the current integrated circuit (IC) fabrication technology.
The development of Si compatible photonic materials includes the observation of an optical gain in Si nanocrystals[2], electrolumines-cence for a Si diode[3], Si nanocrystalfield-effect-transistors[4], imple-mentation of a Si Raman laser[5], and the realization of a high-speed Si electro-optic modulator[6]. However, given silicon's inherent indirect bandgap, crystalline Si is not able to readily emit light. This limits the approaches described above. Thus the lack of an efficient Si-based light source is self-evident.
Erbium (Er) ion implantation in a variety of Si-based[7–9], silica-based[8,10], and ceramic[8]thinfilm hosts has a leading role in the effort to efficiently produce photons from Si. The advantage of this approach is that standard Si technology can be deployed to introduce Er as a dopant. In addition, excited Er3+ions emit at 1.5μm, which is
a strategic wavelength for telecommunications due to the minimum in the absorption for silicafibers. However, photoluminescence is se-verely quenched at room temperature in crystalline Si based hosts
[11,12]. Also, co-implantation of additional O atoms is highly preferred in order to reduce Er precipitation and increase the fraction of active Er3+ions[13]. For silica-based and ceramic thinfilm hosts, a high
pro-cessing temperature is required to grow good quality material. Also,
⁎ Corresponding author.
⁎⁎ Correspondence to: N.P. Kherani, Department of Electrical and Computer Engineering, University of Toronto, ON M5S 3G4, Canada.
E-mail addresses:[email protected](L. Qian),[email protected](N.P. Kherani).
http://dx.doi.org/10.1016/j.tsf.2014.02.038
0040-6090/© 2014 Elsevier B.V. All rights reserved.
Contents lists available atScienceDirect
Thin Solid Films
high post-annealing temperatures (N700 °C) are typically necessary to eliminate the ion implantation-induced damage, to optically activate the Er3 +ions, and/or to enhance the photoluminescence lifetime or
quantum efficiency. These high temperature processes are incompatible with Si BEOL fabrication processes.
While research on Er-implanted silicon and silica-based materials has been extensive, Er doping in amorphous carbon based host has received little attention[14–17]. Hydrogenated amorphous carbon (a-C:H)films can be grown by a low-temperature plasma enhanced chemical vapor deposition (PECVD) method. PECVD methods are compatible with cur-rent CMOS fabrication technology. They allow ease of integration and reproducible processing, and are low-cost. Furthermore, a-C based films possess a number of outstanding properties such as high chemical resistance, biocompatibility, mechanical hardness, and transparency in the infrared[18,19]. Due to their excellent tribological properties, a-C:Hfilms are widely used as protective coatings for hard disks and magnetic media, machine parts, optical windows andfibers, etc.[20]. In the past few years, a-C:H coatings have been also implemented on bio-medical and biosensor products[21]. The specific properties of a-C:H films can be tailored over a wide range by adjusting the sp3to sp2
hybrid-ized carbon ratio, the type of sp3(predominately C\C or a mixture of
C\C and C\H) and sp2
(the number and size of the clusters) bonding configurations, and the amount of incorporated hydrogen in the film via various deposition parameters and deposition methods[19].
The first reported demonstration of room-temperature photo-luminescence (PL) from Er at 1.54μm in a-C:H(Er) thin films was pub-lished in 2002[14]. a-C:H(Er)films were deposited by magnetron sputtering of a graphite target that was partially covered by Er platelets in an Ar/C6H12atmosphere. The Er concentration in the a-C:H(Er)films
could be changed from 0.15 at.% to 1.2 at.%. However, the PL intensity was relatively low. This was attributed to the low optical band gap (~0.5 eV) of the sample and the non-radiative relaxation pathway in-duced by C\H vibrations[15]. In this deposition technique, the Er con-centration highly depends on the degree of magnetron sputtering of the Er/graphite target. Accordingly, high Ar ion energy andflux are required to achieve high Er concentration. However, this causes a high concentra-tion of sp2carbon and a low optical bandgap. The incorporation of an Er
metalorganic compound into a carbon layer by the radio frequency plasma enhanced chemical vapor deposition (RFPECVD) method was demonstrated by Prajzler et al.[16]in 2003. However no PL spectra were presented. In 2009, Tsai et al.[17]grew a-C:H(Er) and a-C:D(Er) films, where D in a-C:D(Er) is deuterium, the isotope of H, via in-situ ther-mal evaporation of the tris(2,2,6,6-tetramethyl-3-5 heptanedionato) erbium(+III), or Er(tmhd)3, compound in a DC saddle-field PECVD
sys-tem. A much higher PL signal was obtained from a-C:D(Er)film compared to a-C:H(Er)films. This was due to the optical quenching from the highly abundant C\H bonds. The Er(tmhd)3metalorganic compound contains a
high percentage of C\H bonds, 58.76 at.%. Hence, this precursor is inher-ently inefficient at promoting Er3+
photoluminescence.
In this work, the feasibility of the in-situ growth of metalorganic Er-doped amorphous carbon (a-C:H(Er)) thinfilms was performed. Films were deposited at low temperature (b200 °C) by a metalorganic radio frequency plasma-enhanced chemical vapor deposition (MO-RFPECVD) system. The properties of the host a-C:Hfilm and the incorporated Er concentration were independently controlled. Prior to the MO-RFPECVD depositions, the effects of the RF power and the placement of substrate in the RFPECVD system were systematically investigated. The structural and optical properties of the host a-C:H material were obtained. These properties were evaluated and discussed with respect to the local bonding features.
2. Experimental details
2.1. MO-RFPECVD and sample preparation
A capacitively coupled MO-RFPECVD system as shown inFig. 1was deployed to deposit hydrogenated amorphous carbon (a-C:H) and Er-doped a-C amorphous carbon (a-C:H(Er)) thinfilms. An ac-powered thermal evaporator was situated next to the RF-powered showerhead electrode (cathode) inside the deposition chamber. The thermal evapo-rator was used to in-situ dope the Er metalorganic compound while commencing a-C:Hfilm deposition via a methane plasma. A thermocou-ple was embedded in the external surface of the bottom of the evapora-tor for feedback temperature control. This measured temperature was
anode showerhead cathode controller Matching box
Deposition pressure gauge Heater assembly 13.56 MHz power supply AC power supply system Temperature controller Low pass filter Precursor gas delivery system To Vacuum system
designated as the nominal evaporator temperature. The temperature of the vapor delivery nozzle was also measured and found to be higher than that of the external bottom surface of the container by 30–50 °C. This temperature difference prevented condensation of the Er metalorganic vapor on the delivery nozzle. Further, it is expected that the temperature inside the container is higher than the evaporation temperature thus enabling the observed sublimation of the Er metalorganic powder. For a-C:Hfilms, the methane (CH4)flow rate
was 40 sccm and the chamber pressure was 16 Pa. The substrates were kept at room temperature, while the RF power was varied from 10 W to 300 W. For the a-C:H(Er) sample, the evaporation temperature was 150 °C. The substrate temperature of 80 °C was used, as opposed to room temperature, so as to aid the uniformity of the preparedfilms as well as to avoid the direct condensation of the metalorganic vapor in powder form; precluding direct condensation ensures vapor-plasma chemistry leading to molecular occlusion of the dopant. The RF power was selected to be 40 W so as to minimize the sp2content in the host
a-C:Hfilm, and to lower the probability of dissociating the Er\O bonds in the pristine metalorganic compound due to energetic ions/ radicals. Theflow rate and pressure was the same as that for the a-C:H films. The c-Si substrates with resistivity of 30 Ω-cm were subjected to the standard CMOS cleaning procedure before being loaded into the chamber.
2.2. Er metalorganic compound
Er(tmhd)3powder has been used as precursor in a-Si:H(Er) samples
prepared by PECVD[22]. The Er in Er(tmhd)3is coordinated to six
oxy-gen atoms, which represents a similar bonding environment to Er2O3.
This bonding environment is thought to be replicated in the a-Si:H(Er) samples since the Er acted as an optically emitting center[22]. Herein, a similar compound, tris(6,6,7,7,8,8,8-hepta fluoro-2,2-dimethyl-3,5-octanedionate) Erbium(+ III), abbreviated (Er(fod)3), with chemical
structure Er(C10H10F7O2)3, as represented inFig. 2, is selected as the
doping candidate for a-C:H(Er)films in this work. The Er(fod)3powder
was obtained from Strem Chemicals Inc. and was loaded into the vacu-um chamber without any special treatment. The Er in Er(fod)3also has a
similar bonding environment to Er2O3. Moreover, the large separation
of the Er ions is expected to reduce concentration quenching. Further-more, thefluorinated ligands are expected to reduce the non-radiative deactivation channels from C\H bonds. This will enable an enhance-ment in the Er3+luminescence efficiency[23,24].
2.3. Film characterization
The thickness (deposition rate), refractive index n, extinction coef fi-cient k, and optical bandgap E04of the a-C:Hfilms were probed through
spectroscopic ellipsometry. The measurements were carried out using a Sopra UV–VIS-NIR spectroscopic ellipsometer. The wavelength range was 350–1700 nm at an incident angle of 75°. The spectra were ana-lyzed by regressionfitting using the linear Levenberg–Marquard algo-rithm method with a maximum of 1000 iterations under a three-layer optical system, void (ambient)/a-C:H layer/c-Si substrate. Afirst-order initial thickness approximation of the a-C:Hfilm was estimated from profilometry measurements. The five constants of Forouhi–Bloomer dispersion model[25]and thickness of the a-C:H layer were allowed to vary during thefitting process. The optical bandgap E04, defined as
the photon energy at which the absorption coefficient α(=4πk / λ) is equal to 104cm−1,whereλ is the wavelength, was determined from
the extinction coefficient k. The coefficient of regression R2~ 0.99, and
the error of the 6fitting parameters was less than +/−10% indicating the model was appropriate for the a-C:Hfilms.
The hydrogen concentration and C\Hxstretching absorption bands
were characterized by Fourier Transform Infra-Red (FTIR) Spectroscopy using a Perkin Elmer 2000 spectrometer with the resolution of 4 cm−1. To calculate the transmission spectra from the thinfilm alone, the back-ground interference pattern due to the multiple reflections in the film was subtracted from the raw transmission spectra. The hydrogen concentration was determined by the following equation[26]
H:conc: ¼AsðCHxÞ
bυN Z
a vð Þ dv ð1Þ
whereα(υ) is the absorption coefficient, As(CHx) is the absorption cross
section of individual stretching mode andb υ N is the average wavenum-ber. The IR spectrum in the wavenumber region 2700–3100 cm−1was
deconvoluted based on the various bond assignments and their corre-sponding wavenumber. As(CHx) was calculated by considering the
cross section in a vacuum environment and extrapolation refractive index n of each sample to 3.3μm[26].
X-ray photoelectron spectroscopy (XPS) was used to quantitatively characterize the elemental composition, bonding configuration, and the depth distribution of the a-C:H and a-C:H(Er)films. The XPS was a Ther-mo Fisher Scientific K-Alpha spectrometer with a monochromatic Al Kα X-ray source. The base pressure of the XPS was of order of 10−7Pa. X-ray excited auger electron spectroscopy (XAES) was employed to estimate the ratio between the sp2-hybridized and sp3-hybridizd carbon atoms
in the host a-C:Hfilms. The percentage of sp2
hybridized carbon is found through Eq.(2) [27]
%sp2¼ Dsample−Ddiamond
Dgraphite−Ddiamond
100 % ð2Þ
where D-parameter can be found in the derivative spectra of XAES measurements.
Photoluminescence (PL) spectra of a-C:H(Er)films were collected at room temperature to verify the optical activity of the Er in the a-C:H(Er) films. A continuous wave diode-pumped solid-state 532 nm laser, with a power density of 80 mW/mm2, was used as the excitation source. The
energy of the laser is nearly resonant with the4S3/2excited level of the
Er ions. The excited Er ions decayed to the4I
13/2level through a fast
non-radiative transition, and then a photon is emitted at 1.54μm through the4I13/2to4I15/2transition. A single pass monochromator,
thermoelectrically cooled InGaAs photodiode, and a standard lock-in amplifier were used.
Fig. 2. Illustration of the Er metalorganic compound, Er(fod)3, with chemical structure Er(C10H10O2F7)3. The large center green atom represents Er, yellow atoms represents O, the dark gray atoms represent C, blue atoms represent F, and white atoms represent H.
3. Results and discussion 3.1. a-C:H host
Fig. 3reveals the dependence of deposition rate, optical bandgap E04,
refractive index n and extinction coefficient k on the applied RF power
for the a-C:Hfilms grown on the bottom powered (cathode, i.e. C-RF) and top grounded (anode, i.e. A-RF) electrodes respectively. Both of the sample sets, C-RF and A-RF, reveal similar trends of an increase in the deposition rate, decrease in the optical bandgap E04, increase in
the refractive index and extinction coefficient with increasing RF power. However, the rate of change of n and k at 532 nm and E04is
more rapid for the C-RF sample set. This indicates that the placement of substrate in this work plays an important role in thefilm deposition process.
The deconvoluted FTIR absorption spectra, normalized on thefilm thickness, provide insight into the presence of the different C\Hx
stretching modes in terms of the hybridization and bond configurations in the a-C:Hfilms. Nine C\H stretching modes have been identified through examination of free molecule vibrational frequencies which comprises the absorption spectra from 2700 cm−1to 3100 cm−1and can be found in Ristein[26].Fig. 4(a) and (b) displays the IR absorption coefficient for the a-C:H films grown on the anode and cathode with an RF power of 60 W. For thefilm grown on the anode inFig. 4(a), there is a wide range of stretching modes. In particular, there is a significant con-tribution from the end groups sp3CH
3and sp3CH2with an estimated
concentration (from Eq.(1)) of 1.78 × 1022and 7.75 × 1021cm−3. In
contrast, for thefilm grown on the cathode inFig. 4(b) there is a compa-rable amount in sp3CH
2(6.82 × 1021cm−3), sp2CH olefinic (3.12 ×
1021cm−3), and sp2CH aromatic (5.18 × 1021cm−3). The total
hydro-gen concentration of the a-C:Hfilms decreases with increasing RF power as displayed inFig. 4(c). Moreover, the hydrogen concentration is lower for the a-C:Hfilms deposited on the cathode. As well, for the films on the cathode, there is larger rate of decrease of the hydrogen concentration with increasing RF power.
As the RF power is increased there is a corresponding increase in the RF voltage and current. This increases the ion energy and the plasma density, and hence there is more ionization, excitation, and dissociation. This leads to an increase in theflux of ions, radicals, and electrons to the substrate surface which corresponds to an increase in the deposition
0 30 60 90 120 150 180 210 240 270 300 0.0 6.4 12.8 19.2 1.82 2.73 3.64 4.55 1.54 1.76 1.98 2.20 10 -4 10-3 10-2 10-1 100 RF power (W) Bandgap (e.V.) (at 532 nm) Extinction coefficient (at 532 nm) A-RF C-RF Refractive index Deposition rate (nm/min)
Fig. 3. The optical properties (E04, n, k) and deposition rate of a-C:Hfilms grown on the anode (solid line), and cathode (dash line), as a function of the applied RF power.
0 60 120 180 240 300 1 2 3 4 5 6 RF Power (W) 0 60 120 180 240 300 0.5 1.0 1.5 2.0 2.5 RF Power (W) 2700 2800 2900 3000 3100 3200 0 200 400 600 800 Wavenumber (cm-1) 2700 2800 2900 3000 3100 3200 0 500 1000 1500 2000 2500 Wavenumber (cm-1)
a
b
c
d
A-RF C-RF sp3CH sp3CH3 sym. sp3CH2 asym. sp3CH 3 asym. sp2CH2 olef sym. sp2CH olef. sp2CH aromatic sp3CH2 sym. sp2CH2 olef asym. A-60W sp3CH2 sym. sp3CH 3 sym. sp3CH2 asym. sp3CH3 asym. sp2CH olef. sp2CH aromatic C-60W A-RF C-RF Hydrogen Conc. (x10 22 cm -3) Absorption constant (cm -1) Absorption constant (cm -1) C-H x sp 2 Conc. (x10 22 cm -3)Fig. 4. Deconvoluted FTIR spectra in the C\Hxstretching region for a-C:Hfilms grown at an RF power of 60 W on the (a) anode and (b) cathode, respectively. (c) The hydrogen concen-tration and (d) the C\Hxsp2bonding concentration as a function of the applied RF power for a-C:Hfilms grown on the anode (solid line) and on the cathode (dash line).
rate for both the cathode and the anode. The deposition rate in the C-RF sample set is 10–15% higher than that in the A-RF sample set. This is attributed to the larger area of the electrically grounded surface and to the higher ion/radical energy impinging on the cathode. In a capacitive-ly coupled RFPECVD system, any asymmetry in the sheath capacitances (anode and cathode) results in a DC bias on the electrodes. Typically, and in the present case, the anode is grounded and the cathode is small-er than the anode. Since the capacitance varies with the electrode area, and the voltage across a capacitor is inversely proportional to its capac-itance, a DC bias is developed on the smaller electrode, the cathode. Hence, the ions and radicals impinging on the cathode possess greater energy than those impinging on the anode. These higher energy ions and radicals can penetrate the surface of a growingfilm and bond to a carbon cluster within the bulk of thefilm, leading to a higher growth rate.
As the RF power increases the ion energy increases since the RF volt-age increases. As well, the DC bias on the cathode is increased from 34 V to 543 V with increasing RF power from 10 W to 300 W. Thus, ions ing the cathode possess a great deal and more energy than those strik-ing on the anode. The increase in ion energy with RF power accounts for the decreasing trend of the hydrogen concentration in the a-C:H film. High energy hydrogen ions/radicals can penetrate into the bulk of thefilm to displace a bonded hydrogen atom, form H2, and desorb
from thefilm[19]. Nevertheless, for the A-RF sample set the hydrogen concentration initially increases then decreases with a peak at 30 W. This is thought to be due to the increased plasma density as the RF power is increased from 10 W to 30 W. Although thefilm density is not taken into account, it is recognized that a relatively low hydrogen atomic density (in atoms/cm3) may actually be transformed into a
rela-tively high atomic percentage (in at.%) if thefilm exhibits a low density. Moreover, the C\Hxsp2bonding modes could provide a qualitative
measure of the change in the relative H bonding configuration with RF power. Hydrogen prefers to bond to sp3hybridized carbon atoms as it
represents a lower energy state than sp2hybridized carbon.
According-ly, the C\Hxsp2would be less likely to occur unless the available sp3
carbon bonds are near saturation due to a high percentage of hydrogen in thefilm. As depicted inFig. 4(d), the C\Hxsp2concentration also
tends to decrease as RF power increases except for the A-RF sample set with low power. Furthermore, the rate of decline of the H concentra-tion and the C\Hxsp2concentration is significantly larger for the C-RF
sample set. This is consistent with the prior discussion about the role of the ion bombardment energy.
Both sample sets (A-RF and C-RF) show an increase in the percent-age of sp2hybridized carbon bonding, from XAES results, as the RF
power increases (plot not shown here). The C-RF sample set possesses much greater sp2bonding than the A-RF sample set as indicated in Fig. 5. The percentage of sp2hybridization increases as the RF power
in-creases due to large ion energy impinging on thefilm and the low film density. High energy ions process enough energy to overcome the pen-etration threshold energy of thefilm, i.e. 32 eV. The excess energy that these ions possess, above the threshold energy, will be transferred to the thermal energy to thefilm. This relaxes C\C sp3bonds to the
more stable C\C sp2configuration[19], leading to a further increase in
%sp2bonding of thefilm. The rapid increase in the % sp2for the C-RF
samples compared to A-RF ones is due to the much larger ion energy impinging on the cathode.
Robertson[19]describes the microstructures of a-C as a contin-uous network of sp3bonded carbon atoms with sp2bonded carbon
localized clusters embedded within the network. The sp3bond
con-figuration forms σ–σ* bands and the sp2sites createsπ–π* bands
which form localized states. The size and quantity of the sp2 clus-ters dominates thefilm's optical properties. Accordingly, the in-creasing % sp2 in thefilm implies an increase in the localized
density of states lying deeper in the gap. This leads to the decrease in E04bandgap and the increase in extinction coefficient
as illustrated inFig. 5.
3.2. a-C:H (Er) PL
As discussed above, the host a-C:Hfilms deposited on the anode exhibit a wider bandgap and less % of sp2carbon bonding. Thus the
hostfilm absorption in 1540 nm wavelength range is expected to be minimized. Hence, the substrates were situated on the anode while per-forming the in-situ doping using the Er(fod)3compound to synthesize
the a-C:H(Er)film.
The room temperature PL spectrum centered at 1540 nm exhibited inFig. 6(a) corresponds to the4I
13/2to4I15/2electronic transition of
Er3+ions. The spectral width of the emission band is due to
inhomoge-neous and homogeinhomoge-neous broadening in addition to Stark splitting of the Er3+excited and ground states. The PL peak is wider than that of other
Er-implanted silicate glasses[8], indicating the Er3+possesses a variety
of local bonding environments in the a-C:H matrix. Its 65 nm of full width at half-maximum (FWHM) suggests the potential of enabling a wide gain band width for optical amplification. From XPS analysis, the concentration of Er of a-C:H(Er)film is estimated to be 3.9 at.%, which is much higher than those prepared by magnetron sputtering[15,16], pulsed laser deposition[28,29], and DC Saddle-Field PECVD[17]. Furthermore, the depth distribution depicted inFig. 6(b) reveals a uni-form concentration of Er throughout thefilm (thickness of 850 nm). This contrasts the ion-implantation of Er where the optically active ions are always located near the surface[8]. Note that the high oxygen at the surface is simply surface contamination.
The prominent PL signal observed from the a-C:H(Er)film is attrib-uted to several factors including a high Er concentration, the large opti-cal bandgap of the a-C:H host, and the decrease in the C\H quenching. The long hydrocarbon ligands of the Er metalorganic compound matches the internal structure of the a-C:H host matrix. This is thought to result in a high solubility of the Er metalorganic compound, and hence promote a high Er concentration. The large optical bandgap of the a-C:H host is thought to increase the pumping efficiency of the Er3+ions without high absorption from a-C:H host itself. To effectively
20 40 60 80 100 10-4 10-3 10-2 10-1 1.76 2.64 3.52 4.40 % sp2 bonding
C-RF
A-RF
Bandgap (e.V.) Extinction coefficient (at 532 nm)Fig. 5. The optical bandgap E04and extinction coefficient k at 532 nm of the a-C:H films grown on the anode (solid line), and cathode (dash line), as a function of the % sp2 bonding.
decrease the C\H quenching, partial fluorination is incorporated into the a-C:H(Er)film from the Er metalorganic compound.
The ratio of the O to Er concentration in the a-C:H(Er)film is approx-imately 1.5 in thefilm instead of 6 in the pristine Er(fod)3compound. Table 1lists the ratio of the atomic concentrations and the relative (and the absolute) atomic concentrations of the Er(fod)3compound,
thermally evaporated Er(fod)3, and the a-C:H(Er) form XPS
mea-surements. The [F]/[O], [C]/[O], and [C]/[F] concentration ratios are approximately 5.9, 9.1, and 1.5, which is greater than the Er(fod)3
stoi-chiometric ratios of 3.5, 5, and 1.4. However, the [F]/[Er] and [C]/[Er] concentration ratios are approximately 9 and 14, all less than the Er(fod)3stoichiometric ratios of 21 and 30. This implies that the plasma
causes significant dissociation of the Er(fod)3metalorganic compound.
It is noted that the thermally evaporated Er(fod)3 film possesses
essentially the same relative concentration as the Er(fod)3compound,
with the exception that the [C] is enhanced. The [F]/[O], [F]/[Er], and [O]/[Er] ratios are similar, while the [C]/[O], [C]/[F], and [C]/[Er] are much larger than the stoichiometric Er(fod)3compound. This increase
in the [C] is thought to be due to the trapping or intermolecular bonding of methane like species in the Er(fod)3film.
For the a-C:H(Er)film,Fig. 6(c) displays the C1s XPS spectra indicating the majority of carbon bonds are C\C and C\H, with some C\F, and a few C\O bonds. Although, relative to [Er], the concentrations of O, F,
and C have decreased in descending order ([C] has decreased the least with respect to [Er]). This is witnessed by the greater [F]/[O], [C]/[O], and [C]/[F] ratios. It is surprising that the [C]/[Er] ratio has decreased, since CHx(x = 0…3) radicals/ions are also being deposited by the
meth-ane plasma. Although this could be due to the loss of large organic frag-ments since the concentration of all species relative to Er is lower (when compared to the Er(fod)3power orfilm). However, the relative
(as [H] is unknown) atomic concentration of F is high. This may indicate that part of the molecule (CmF7, mN 5) accounts for a significant fraction
of thefilm. This is a positive result, since the Er(fod)3compound was
selected for the purposes of incorporating the largefluorinated ligands into the a-C:H matrix. With respect to the change infilm properties after Er metalorganic compound incorporation, the % sp2bonding of the
a-C:H(Er)film is expected to increase compared to that in the a-C:H host due to the addition of Er dopant in the carbon matrix based on the observations from Foong et al.[28,29]. In addition, with the high fraction offluoride composites, as analyzed from XPS measurement, the % sp2
bonding would be expected to be further enhanced as demonstrated pre-viously[30–32]. Accordingly, a decrease in the optical bandgap E04and an
increase in the extinction coefficient k of the film would occur. The de-tailed optical properties of the a-C:H(Er)film are still under investigation. One obvious result is that oxygen is being reduced or omitted from the a-C:H(Er)film. This is clearly seen through the elevated
0 250 500 750 1000 0 3 6 9 12 30 40 50 60
Distance from surface (nm) C F O Er 166 168 170 172 174 176 178 180 Er (fod)3 powder Er (fod)3 film a-C:H(3.9 at %Er)
Binding Energy (eV)
1450 1500 1550 1600 1650
Wavelength (nm)
278 280 282 284 286 288 290 292 294
Binding Energy (eV)
a
d
c
b
C-C/C-H C-CF/ C-O C-F/C=OIntensity (arb. units)
Concentraiton (at%)
Intensity (arb. units)
PL Intensity (arb. units)
Fig. 6. (a) Room-temperature PL spectrum of a-C:H(Er) with peak centered at 1540 nm and FWHM of ~65 nm. (b) Depth profile of C, F, O, and Er concentrations from XPS measurements. (c) Deconvoluted XPS C1s spectrum of the a-C:H(Er)film. (d) XPS Er4d spectra of a-C:H(Er) film, Er(fod)3film (evaporated in the vacuum chamber with CH4precursor gasflowing without plasma ignition), and Er (fod)3powder.
Table 1
The ratio of the atomic concentrations and the relative/absolute atomic concentrations of the Er(fod)3compound, thermally evaporated Er(fod)3, and the a-C:H(Er)film from XPS measurements.
C at.% F at.% O at.% Er at.% [F]/[O] [C]/[O] [C]/[F] [F]/[Er] [C]/[Er] [O]/[Er]
Er(C10H10F7O2)3 34.1a 23.9a 6.8a 1.1a 3.5 5.0 1.4 21.0 30.0 6.0
Er(fod)3thermal evaporatedfilm 71.3 21.5 6.3 1.0 3.4 11.4 3.3 21.7 72.1 6.3
a-C:H(Er)film 54.6 35.4 6.0 3.9 5.9 9.1 1.5 9.0 13.9 1.5
a
concentrations of C, F, and Er relative to oxygen. Although, the obser-vance of a PL signal indicates that some of the erbium is in the 3 + state.Fig. 6(d) shows the XPS Er4d spectra for the a-C:H(Er) film, Er(fod)3evaporatedfilm, and the Er(fod)3powder. The three spectra
are similar indicating that the local bonding environments are also similar. That is, oxygen is bonded to the erbium in the correct con figura-tion. Moreover, the proximity of the C\H bonds is thought to be far enough from Er3+ions to avoid quenching. However, the optimization
of the process and the electronic transfer mechanism are still under investigation.
4. Conclusion
The feasibility of the in-situ growth of Er-doped a-C thinfilms (a-C: H(Er)) on Si substrates at low temperature (b200 °C) by a simple one-step MO-RFPECVD system was demonstrated. A high Er concentration (3.9 at.%) in a-C:H(Er) films was achieved and room-temperature photoluminescence peaking at 1.54μm was observed. By adopting an Er metalorganic precursor, Er(fod)3, the optically active Er3+ions are
preserved without the need for any subsequent high temperature annealing. Furthermore, the in-situ thermal evaporation technique provides the potential of doping Er in a vertically uniform profile. In addition, non-radiative C\H vibrational quenching was significantly reduced by partialfluorination of the surrounding ligands. This was achieved despite the use of hydrogenated amorphous carbon as the host material.
Acknowledgments
This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC) Discovery grants, the NSERC CREATE program in Nanoscience and Nanotechnology, and the Depart-ments of Electrical and Computer Engineering at the University of Toronto and Materials Science and Engineering at the University of Toronto. The authors also acknowledge valuable discussions with Dr. Rana Sodhi for the analysis of XPS results, and help from Dr. Davit Yeghikyan and Dr. Tome Kosteski for the assembly of the MO RF-PECVD system.
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