Silicon-based Yb-doped Y3Al5O12 garnet nanofilms are fabricated by atomic layer deposition, which are polycrystalline after annealing at 1150 °C. The sub-nanometer compositional regulation and the Yb2O3 cladding layers, which also work as the luminescent dopants, are critical for the crystallization. Characteristic Yb3+ luminescence at 1030 nm and 970 nm is identiﬁed under electrical injection, exhibiting the external quantum efficiency of 0.65% and the fluorescence lifetime of 80-200 µs. The doped Yb3+ are impact-excited by hot electrons stemming from Fowler-Nordheim tunneling mechanism within the Y3Al5O12 matrix, with the excitation cross section of 0.7×10−15 to 6.4×10−15 cm2. This work certifies the manipulation of multi-oxide nanofilms with designed composition and crystallinity, revealing the possibility of developing Si-based optoelectronic devices from crystalline garnet films.
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Y3Al5O12 garnet (YAG) crystals doped with Nd3+ or Er3+ ions are promising materials for modern lasers, providing efficient emissions in the near-infrared (NIR) range. Research has shown the possibility of growing this material of good structural quality with high doping concentrations [1 –3]. The component Y3+ ions have a simple electronic structure without radiative recombination route, thus the reduction of luminescence efficiencies due to excited state absorption, cross relaxation and upconversion, is absent . Other attractive qualities of YAG, including isotropic, high thermal conductivity and tolerance of high concentration doping of rare earth (RE) ions, make it desirable for optical applications . Yb3+ ions are widely utilized in the integrated amplifier and tunable laser applications due to their wide gain spectrum around 1 µm [6 –8]. For Yb-doped YAG, the remaining radiative recombination is restricted to the ∼1 µm emission without the competition from other excited energy levels, making it attractive in highly efﬁcient lasers and optoelectronics [8 –11]. In addition, the radii of Y3+ and Yb3+ ions are similar, the replacement should be quite easy and a high-level doping is achievable.
Integration of optical devices into the microelectronics has long been a great scientific and technological challenge [12 –15]. However, efficient Si-based light emitting devices compatible with complementary CMOS technology are still under development and far from application. RE ions embedded in wide bandgap materials are of particularly interest due to their intra-4f transitions providing intense sharp emissions [16 –20]. Such light emitters are CMOS-compatible and represent not only the basis for inter-/intra-chip optical interconnection but also lightning, display, waveguide amplifier, and reagent detection [21,22]. Recently, MOS-structured light emitting devices (MOSLEDs) based on various RE-doped nanofilms have been demonstrated, with competitive external quantum efficiency (EQE) [23 –26]. These devices present several issues not yet resolved such as the efficient carrier injection and the concentration quenching . The RE ions tend to segregate and precipitate above a certain dopant concentration, leading to the formation of optically inactive clusters and the degradation of light emission yield, depending on the host matrix and doping distribution [18,19,24 –27].
Bearing the above-mentioned in mind, the controllable synthesis of Yb-doped YAG nanofilms on silicon substrates with electroluminescence (EL) performance would apparently be interesting and of great potential for optoelectronic applications. However, little achievement has been reported in previous literature. The precise control of the quantity of different elements together with a high temperature are needed for crystallization. Considering the cell dimension of ∼1.2 nm for the YAG garnet crystals, the control of film deposition in a sub-nanometer scale is required, which helps to decrease the temperature for crystallization due to a shorter migration distance. Considering the optoelectronic utilization, although the melting temperature of YAG crystal is as high as 1970 °C, an annealing temperature higher than 1200 °C is not acceptable as it is destructive to the widely used silicon substrate.
We have fabricated Er-doped Yb3Al5O12 (YbAG) garnet polycrystalline nanofilms by atomic layer deposition (ALD) on silicon substrate, which is the first report on the growth of crystallized garnet films on silicon by this technique, giving the possibility of the deposition of crystalline YAG using the same method . ALD is a diverse method capable of depositing high quality nanofilms with excellent consistency over large substrates [29,30]. Since the growth procedure of ALD is governed by self-limiting surface reactions, the precise control of thickness and dopant with an Angstrom level can be conveniently achieved, by alternating the series of deposition cycles of various precursors. The realization of EL from Yb3+ ions in the Si-based YAG ﬁlm by ALD is worth of exploring and provides new ideas for optoelectronic applications.
In this paper, the suitability to fabricate crystallized YAG nanofilms on silicon by ALD with good optoelectronic and morphological properties is validated. The Yb2O3 layers not only work as the luminescent dopants but also promote the crystallization of the deposited films after annealing at 1150 °C. The EL from YAG:Yb MOSLEDs peaks at 1030 nm and 970 nm, with the fluorescence lifetime of 80-200 µs. The emission originates from the infra-4f shell transitions of Yb3+ ions, excited by the impact of hot electrons. The optimal device exhibits the EQE of 0.64% and the power efficiency (PE) of 8×10−5. This work contributes for the development of photonic integrated circuits and solid-state lighting in which crystalline composite nanofilms fully compatible with the CMOS technology are required.
2. Experimental details
The growth parameters of the nanolaminates have been reported previously, using a 4-inch chamber ALD system (MNT Micro and Nanotech Co., Ltd.) [19,25,26,28]. Briefly, different Al2O3/Y2O3/Yb2O3 nanolaminate films were grown on the <100>-oriented phosphorous-doped silicon substrates (n-Si, 2-5 Ω·cm) at the deposition temperature of 350 °C. Trimethylaluminum [Al(CH3)3, maintained at room temperature], Y(thd)3 and Yb(thd)3 (thd = 2,2,6,6-tetramethyl-3,5-heptanedionate, both heated to 195 °C) were used as the precursors for Al2O3, Y2O3 and Yb2O3, respectively, with ozone acting as the oxidant. The growth rates of the Al2O3, Y2O3 and Yb2O3 films are 1.00, 0.21 and 0.19 Å/cycle, respectively. Firstly, a 2 nm Al2O3 buffer layer was deposited, then the nanolaminate composed of several Y2O3 and Yb2O3 cycles together with 1 nm Al2O3 was deposited repeatedly, the number of Y2O3/Yb2O3 cycles was adjusted to change the stoichiometry and doping level in these nanolaminates. The schematic structures of the ALD nanolaminates are shown in Fig. 1(a). The total cycle numbers were adjusted correspondingly to obtain the designed thickness. After the deposition, the films were annealed at 1150 °C in N2 atmosphere for 1 h to enable crystallization and Yb activation. Then the ∼100 nm TiO2-Al2O3 dielectric nanolaminates composed of ∼2.0 nm Al2O3 and 8.0 nm TiO2 sublayers were deposited by ALD on the luminescent films at 250 °C, which work as the protection layers to enhance the device stability during operation. Afterwards, ∼100 nm transparent conductive ZnO:Al films were deposited on the TiO2-Al2O3 nanolaminates by ALD at 150 °C, which were lithographically patterned into 0.5 mm circular electrodes. Finally, ∼100 nm Al films were thermally evaporated on the backsides of n-Si substrates. The resulting devices have a multi-layered structure of ZnO:Al/TiO2-Al2O3/Y3Al5O12:Yb/Si/Al.
All the characterization of the luminescent films and devices were performed as previously reported at room temperature [11,19,25,26]. The film thickness was measured by an ellipsometer with a 632.8 nm He-Ne laser beam at an incident angle of 69.8°. The crystal structures of the ﬁlms were characterized by the glancing angle X-ray diffraction patterns (XRD, D/max 2500/pc, Rigaku, Cu Kα radiation, λ=1.5406 Å). The sample composition was characterized by the Rutherford backscattering spectroscopy (RBS, 5SDH-2, NEC, Japan), using the 2 MeV 4He ion beam at a scattering angel of 160°. The surface morphologies of these ﬁlms were characterized by the scanning electron microscope (SEM, JSM-7800F, JEOL). The microstructures of the nanoﬁlms were characterized by the transmission electron microscope (TEM, Tecnai G2 F20 S-TWIN, FEI) equipped with selected area electron diffraction (SAED).
To activate EL from the MOSLEDs, appropriate forward bias was applied with the negative voltage connecting to n-Si substrates. EL and Current-Voltage (I-V) characteristics were recorded by a Keithley 2410 SourceMeter unit. The EL signal was collected by a 0.5 m monochromator and detected by an InGaAs detector connected to a Keithley 2010 multimeter. The absolute EL power from the device was measured by a calibrated optical power-meter (1830-C with 818-IR Sensor, Newport). The decay time of the EL emission was measured by an InGaAs single-photon avalanche diode (id220-FR-MMF, IDQ) connected to a SR430 multi-channel scaler (Stanford Research Systems) equipped with a waveform generator (DG5072, RIGOL) and a high-voltage amplifier.
3. Results and discussion
The thickness of the as-deposited films is 54.5-55.5 nm, thinner than the designed value of ∼60 nm, which attributed to the affected growth for the first ALD cycles while switching precursors for different deposition sequences, especially for the RE2O3 films, whose low deposition rates result in the susceptibility to changes in surface condition and vapor pressure. In comparison, the growth velocity for Al2O3 always maintain a stable value. Herein the growth rate of Al2O3 and the final thickness of the film are preferably used to calculate the RE2O3 thickness and corresponding nominal RE:Al ratio in the deposited nanofilm. These calculated film components from the thickness present a linear relationship with the designed RE contents, but slightly lower than the componential values obtained by fitting the RBS spectroscopy of these nanolaminates, which might be ascribed to the measurement error of different elements and the different densities of the deposited nanofilms from the standard values.
The crystallization of the luminescent films after annealing at 1150 °C mainly depends on two factors, the RE:Al ratio and the position of the Yb2O3 dopant cycles in the nanolaminates. For the former, more Y3+ ions are needed as the actual density of Y2O3 might be lower than theoretical one, therefore the calculated RE:Al ratios are higher than actual values. For the latter, the Yb3+ ions present higher reactivity with the Al2O3 interlayers to transform to the garnet structure. In the following, XRD patterns of the luminescent nanolaminates are compared, all of which were measured under the same condition. Firstly, for the conventional doped structure Y-Yb-Y (with the Yb cycles buried in the middle of Y2O3 layers, shown as RE2O3 structure 1 in Fig. 1(a)), we found the nanolaminate films annealed at 1150 °C were hard to crystalize into garnet phases, while in our previous report, the Yb2O3 and Al2O3 nanolaminates could react to transform to YbAG grains above 1100 °C . As shown in Fig. 1(b), the RE:Al ratio should be higher than 0.85 to present diffraction peaks assigned to YAG grains, further increase of the RE:Al ratio gradually reduces the diffraction peaks, showing an adverse effect to the grain crystallization. The only prominent diffraction peak comes from the (420) crystal plane, with low intensities. According to their SEM images shown in Fig. 2, the films with RE:Al ratio above 0.85 present prominent grains on the substrates after annealing, but further increase of the ratio results in less grains, which is reasonable as the Al content is not sufficient for Y3Al5O12 crystallization, and the grains seem unable to completely cover the substrates. In addition, we found that the crystallization has less to do with the interlayer thickness. For these nanolaminates with the RE:Al ratio of 0.85 and changing the Al2O3 layers from 0.5 to 1.5 nm, their XRD patterns present no apparent difference (shown in Supplement 1, Fig. S1).
In comparison, we found that for the RE2O3 layers composed of Yb-Y-Yb (Y2O3 layers sandwiched within Yb2O3 dopant cycles, shown as RE2O3 structure 2 in Fig. 1(a)), the RE:Al ratio needed for good crystallization is much smaller, and the full coverage of substrates is easier to achieve. Figure 1(c) shows the XRD patterns of the nanolaminates with RE2O3 layers composed of Yb-Y-Yb but different RE:Al ratios. Apparently, the Yb2O3 layers manifest stronger reactivity with the Al2O3 interlayers, not only the main (420) peaks are stronger, but also more diffraction peaks present in the patterns, such as the (211) peaks. The SEM images of these Yb-Y-Yb nanolaminates are consistent with the XRD results, as shown in Fig. 3, well-crystallized grains cover the entire surface for the nanolaminate with a RE:Al ratio of 0.69. Further increase of the RE:Al ratio to 0.96 still obtains similar crystalline film, with grains stacked slightly looser, also indicating the deficiency of Al2O3 for crystallization. Moreover, from the cross-sectional SEM image (shown in Supplement 1, Fig. S2), after the high temperature annealing, further oxidation and reconstruction of the garnet grains bring about the increased film thickness to ∼100 nm. The annealed nanofilms are densely stacked by grains and although the surficial morphology gives high roughness, no exposure of substrate could be seen from the SEM images. The composition of the film demonstrates distinct contrast with the underneath Si substrate, as shown in the energy dispersive spectrometer (EDS) mapping image. Apparently, the Yb-Y-Yb nanolaminates transform to YAG grains easier and more prominent than the conventional Y-Yb-Y nanolaminates (Yb-doped Y2O3 layers). It should be mentioned that even the 3-cycle Yb2O3 cladding layers (∼0.57 Å) could achieve the above-mentioned promotion of crystallization, and the highest Yb/Y ratio among the samples is lower than 10.2% so the main component should still be YAG.
The microstructure of the luminescent nanolaminates were further characterized using TEM equipped with SAED. As shown in Fig. 4(a) and (b), the crystallized grains exhibit the lattice fringes separated by 0.489 nm and 0.269 nm, which are assigned to the (211) and (420) crystal planes of the YAG phase, respectively [17 –19]. In Fig. 4(c), the SAED pattern focused on a selected crystalline area demonstrates the existence of the (200) and (211) planes of YAG grains, also illustrating the formation of polycrystalline YAG nanofilms. The long-range ordered arrangement of atoms together with the series of single-crystal diffraction dots manifest the well-crystallized garnet grains. These results are consistent with the XRD and SEM ones.
The melting temperature of YAG is ∼1950 °C, a little lower than that of YbAG (2000-2200 °C). However, in this work we find the crystallization of Yb-Y-Yb nanolaminates (Yb2O3 sandwiched Y2O3 nanofilms) with the Al2O3 interlayers is easier than that of Y-Yb-Y nanolaminates (Y2O3 nanofilms with Yb2O3 cycles doped inside). The reaction of Yb2O3 with Al2O3 is realized at 1100 °C while for Y2O3 is negligible, and a higher Y content is needed for crystallization of the Y-Yb-Y nanolaminates with Al2O3 interlayers . Considering the difference on YAG and YbAG garnet crystals, the isolated ionic radius for Y3+ ion is 0.90 Å, larger than the Yb3+ (0.868 Å) and Al3+ (0.535 Å) ions. The ionic radius of Y3+ and Yb3+ ions in the dodecahedral positions of the garnet structure are 1.019 Å and 0.985 Å, respectively. Considering that the unit cell sizes of YAG and YbAG are 12.01 Å and 11.93 Å, these two ions occupy a similar space ratio. The formation enthalpies for these two garnet crystals are −7159 and −6987.9 kJ/mol, respectively, and the driven force for generation of YAG is slightly higher. The standard formation energy / atom for YAG is −3.702 eV, still larger than that for YbAG, which is −3.244 eV, derived from the density functional theory calculations . Therefore, the leftover possible reason is the Yb3+ ions are more likely to migrate and penetrate into Al2O3 matrix to crystallize into garnet grains, as they maintain a smaller radius and electronegativity than Y3+ ions.
Figure 5(a) shows the EL spectra from the MOSLED based on the polycrystalline YAG:Yb nanofilms under different injection currents. The EL spectra mainly contain two sharp peaks at the wavelength of 970 nm and 1030 nm, with the latter stronger. In addition, some shoulder peaks appear in the spectra at around 943, 1007 and 1048 nm. These EL emissions are well assigned to the radiative intra-4f transitions from the 2F5/2 excited state to the 2F7/2 ground state in Yb3+ ions, originating from different Stark levels split by the spin-orbit interaction into two energy manifolds [32,33]. The spectra are quite different from the previous report on MOSLEDs based on Yb-doped Al2O3 films, which present only the 977 nm main EL peaks with weaker emission bands at 1000-1030 nm, again manifesting the unique crystal field of YAG matrix which affects the proximity of Yb3+ ions [25,26]. We have found a similar phenomenon in the Er-doped Yb3Al5O12 polycrystalline nanofilms . For the MOSLEDs fabricated with the Y-Yb-Y nanolaminates, the electron injection is not stable, leading to a breakdown current of ∼10−5 A, which pertains to the poor crystallization and morphology (shown in Supplement 1, Fig. S3). Only the best crystallized device with the RE:Al ratio of 0.85 presents comparable EL performance with the MOSLEDs fabricated with the Yb-Y-Yb nanolaminates.
Figure 5(b) shows the dependence of the EL intensity on the current for the MOSLEDs composed with polycrystalline YAG:Yb nanofilms. Generally, the devices based on the Yb-Y-Yb nanolaminates work compatible with the best MOSLEDs fabricated by the Y-Yb-Y nanolaminates. Moreover, considering the different RE:Al ratios, RE2O3 ratio of around 0.86 is more preferable, the optical power density from the best devices reach 2.45 mW/cm2, pertaining to both the better morphologies of polycrystalline films and the suitable doping concentration. Since these Yb-Y-Yb nanolaminates maintain the same Yb2O3 cladding cycles while the Y2O3 cycle numbers were adjusted to tune the RE:Al ratios, here the calculated Yb-doping concentrations for these nanolaminates with the RE:Al ratio of 0.69, 0.86 and 0.96 are 5.1, 4.4 and 3.6 at%, respectively. The optimal 4.4 at% dopant concentration manifests a high tolerance of the YAG matrix for the doped Yb3+ ions, the similar radii of the luminescent Yb3+ ions with the component Y3+ ions in the YAG grains restrict the clustering and cross-relaxation among adjacent Yb3+ ions, avoiding the reduction of the excitable ions [34,35]. It is believed that demonstrating sub-nanometer control over the superlattice period allows for placing the dopant atoms in their more favorable environment for luminescence and opens new routes to the engineering materials with targeted optical and electrical properties. Again, the appropriate RE:Al ratio serves as an important factor on the EL performance, MOSLEDs fabricated by the Yb-Y-Yb nanolaminates with a moderate RE:Al ratio present stronger emission. Therefore, the recipe and doping structure discussed in this work is critical not only for the deposition of well-crystallized integral YAG:Yb films, but also the EL performance from MOSLEDs based on these nanofilms.
Figure 5(c) presents the dependence of representative 1030 nm EL intensity, together with the current, on the applied bias voltage for the MOSLEDs based on the polycrystalline YAG:Yb nanofilms with different RE:Al ratios. The EL-V and I-V characterizations show no apparent divergence from previous reports on YbAG and RE-oxides devices. All devices feature the typical I-V characteristic of a MOS structure, the injection currents increase approximately exponential within the EL enabling voltage (electric field), which start with a low background one for low electric fields and end with the dielectric breakdown [11,17,28,36]. The EL intensities increase almost linearly with the injection currents above the threshold current of ∼0.1 µA. As shown in the SEM images, the luminescent films are composed of irregular crystal grains measuring several microns, which are stacked densely together. No prominent short-circuit appears during the device measurement, in comparison with our previously-reported MOSLEDs based on RE-doped amorphous Al2O3 nanofilms. Therefore, all the nanolaminates are quite dense without pinholes that causing leakage.
In comparison of the EL lifetime of different YAG:Yb MOSLEDs, Fig. 6(a) and (b) illustrate the decay traces of the ∼1030 nm EL emissions, which can be characterized as the single exponential decay function. For the relatively weak EL from the Y-Yb-Y nanolaminate devices, the ﬁttings of the decay traces derive the lifetime of 180∼300 µs, which are close to that for Yb-doped Al2O3 films but lower than that for Yb-doped YAG grown by other methods [8,25,37,38]. In comparison, for the polycrystalline devices based on the Yb-Y-Yb nanolaminates, the derived EL decay time decreases to 85-120 µs, partially due to the higher doping concentrations (3.6-5.1 at%, 2.3-3.6 at% for the former Y-Yb-Y series). The increase of the RE:Al ratio which further disperse the Yb dopant in the RE2O3 nanolaminates could increase the EL lifetime, as the cross-relaxations (non-radiative decay channels) are further suppressed, beneficial for both the EL intensity and lifetime. This conclusion is consistent with the difference on the EL lifetime of the Yb-doped Y2O3 ceramics .
For the MOSLEDs based on RE-doped oxides, the emissions always originate from the direct impact excitation of RE ions by hot electrons in the conduction band of oxide matrix, while the conduction mechanism varies, mainly depends on the interfacial condition resulting from the annealing process [13,27,40]. According to previous research, the conduction in these high-temperature annealed crystalline oxides are attributed to the Fowler-Nordheim (FN) tunneling mode which can be simply expressed as J = AE 2 exp(-B/E), where J is the current density, and E is the electric field. In simplicity, the plot of the ln(J/E 2) versus the reciprocal electric field 1/E presents linear relationship in this conduction mode [11,28]. Here the plots of the I-V characteristics derived from the polycrystalline YAG:Yb nanofilms with relatively better EL performance are shown in Fig. 6(c). Well-defined linearity is established for all the MOSLEDs in the EL-enabling voltage region. The minor difference on the curve slopes and threshold electric fields are mainly ascribed to the current discrepancy brought about by the variation of RE:Al ratio and the fabrication procedures. Therefore, for the EL from Y3Al5O12:Yb films, the electrons tunnel into the conduction band of Y3Al5O12 through potential barrier under a sufficient electric field, governed by the FN mechanism. These injected electrons are accelerated by the electrical field to gain kinetic energy, which subsequently excite the doped Yb3+ ions through inelastic impact, the following de-excitation process results in the characteristic emissions. Compared with the previous AC-driven thin-film EL (ACTFEL) devices, the device structure and excitation mechanism of the MOSLEDs reported here are similar. The previous achievements on the ACTFEL devices, including the optimization on the dielectric layers, are instructive to the improvement of EL performance in this work, which governs the distribution of electric field within the luminescent nanofilms and promotes the impact excitation by hot electrons [19,36].
The excitation cross section of Yb3+ ions is evaluated to quantify the efficiency of the excitation process via electrical pumping, according to the previous established method by simulating the dependence of the EL intensity on the injection current . The excitation cross sections of the EL from YAG:Yb nanofilms are in the range of 0.7×10−15 to 6.4×10−15 cm2, which are generally consistent with the tendency of EL intensity shown in Fig. 5. The EQE and PE versus the injection current are compared in Fig. 6(d), in evaluation of the performance of these YAG:Yb MOSLEDs. For the best MOSLED fabricated by the Y-Yb-Y nanolaminates, it presents a declining efficiency which is lower at the high current. In comparison, the EL efﬁciencies from the MOSLEDs fabricated by the Yb-Y-Yb nanolaminates gradually increase to a broad maximum, and ﬁnally fall down at the high currents, among which the best device presents a maximum EQE of 0.65%, together with the PE of 8.1×10−5. These lower excitation efficiencies compared with the MOSLEDs based on Al2O3:Yb nanofilms should be mainly ascribed to the higher defect and trap density contributing to the non-radiative de-excitation channels, which results from the polycrystalline garnet films with abundant grain boundaries. Further optimization on the film thickness, dopant distribution and crystal quality could increase the efficiencies [19,25].
In summary, by means of ALD technique and 1150 °C annealing, polycrystalline YAG:Yb nanolaminates are fabricated on silicon. The dopant Yb2O3 cycles deposited on both sides of Y2O3 layers are beneficial for the crystallization of these nanofilms, promoting the reaction of the RE2O3 nanolaminates with the Al2O3 interlayers and the resultant crystallization of YAG grains. Higher RE:Al ratios than stoichiometry are needed for the crystallization of YAG phase, while the morphology deteriorates with too low Al content. The MOSLEDs based on YAG:Yb nanolaminates present intense NIR EL at ∼1 µm, originating from the intra-4f 2F5/2→2F7/2 transitions of Yb3+ ions, with the EQE of 0.65%, the PE of 8.1×10−5, and the fluorescence lifetime of 80-200 µs. The EL is attributed to the impact excitation of the Yb3+ ions by hot electrons stemming from FN tunneling mechanism within the Y3Al5O12 matrix. The excitation cross section for the Yb3+ emission is in the range of 0.7×10−15 to 6.4×10−15 cm2. This work manifests the manipulation of multi-oxide nanofilms with designed composition and crystallinity, exploiting an available method of developing Si-based optoelectronic devices from crystalline garnet films.
National Natural Science Foundation of China (61674085, 61705114); Fundamental Research Funds for the Central Universities (63191409).
The authors declare no conflicts of interest.
See Supplement 1 for supporting content.
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