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Spectroscopic properties of Er3+-doped α-Al2O3
Optical Materials 35 (2013) 821–826
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Spectroscopic properties of Er3+-doped a-Al2O3 T. Sanamyan a,⇑, R. Pavlacka b, G. Gilde b, M. Dubinskii a a b
U.S. Army Research Laboratory, RDRL-SEE-M, 2800 Powder Mill Rd., Adelphi, MD 20783, USA U.S. Army Research Laboratory, RDRL-WMM-E, Building 4600, Aberdeen Proving Ground, MD 21005, USA
a r t i c l e
i n f o
Article history: Available online 5 December 2012
a b s t r a c t We present the results of comprehensive spectroscopic study of the recently successfully synthesized nanopowders and fully densified transparent ceramics of Er3+-doped a-Al2O3 (corundum). Fluorescence and excitation spectra in a range of temperatures of 10–300 K were used to infer the energy level schemes of the ground state and the first three excited Er3+ multiplets (4I15/2, 4I13/2, 4I11/2, and 4I9/2) in this material. Obtained experimental data also allowed us to assess this material as a potential gain medium for high power laser applications in anticipation that it becomes available in the low-loss grain-aligned ceramic or single crystalline form in the near future. Published by Elsevier B.V.
1. Introduction Despite recent significant advances in power scaling of Y3Al5O12 (YAG) based cryogenically cooled lasers with nearly diffraction limited beam quality [1], including power scaling toward laser fission application [2], the extreme pump power densities combined with the unavoidable non-radiative losses in the pump-lase process introduce severe thermally induced optical distortions in the aggressively pumped and aggressively cooled gain medium. Regardless of the sophistication of heat removal techniques and their efficiency, the insufficient thermal conductivity of the currently used gain media is the bottleneck for non-distortive heat removal. The development of new gain media with extremely high thermal conductivity remains the most critical technological challenge preventing power scaling of bulk high energy solid-state lasers with high beam quality. Due to its unique suite of thermal, optical, and electrical properties, rare earth doped aluminum oxide (RE:Al2O3) is the subject of ever-growing interest as the gain material for microelectronics, integrated on-chip devices for telecommunication, and most recently, for high power laser applications. Indeed, single crystal aAl2O3 (sapphire) has a liquid nitrogen temperature (LNT) thermal conductivity of well over 1000 W/m K [3], which is 40 times higher than the LNT thermal conductivity of YAG (the most prevalent solid state laser host) and approximately two orders of magnitude higher than the room temperature (RT) thermal conductivity of YAG. Due to the mentioned above and the favorable combination of other LNT parameters (dn/dT = 0.19 105 K1, and CTE = 0.34 106 K1) which define the thermal figure of merit ⇑ Corresponding author. Tel.: +1 301 394 2044; fax: +1 301 394 0310. E-mail address: [email protected] (T. Sanamyan). 0925-3467/$ - see front matter Published by Elsevier B.V. http://dx.doi.org/10.1016/j.optmat.2012.10.036
of the cryo-cooled doped a-Al2O3 laser material, it can be pumped far more aggressively with far less thermally induced optical distortion. Thus, a cryo-cooled rare earth doped sapphire (RE: Sapphire) laser would possess extremely high power extraction capabilities, far exceeding those possessed by state-of-the-art solid state laser gain media. To date, RE:Al2O3 devices have been synthesized predominantly as thin films for optoelectronics applications using multiple thin film deposition techniques. In most instances these films are amorphous aluminum oxide layers deposited on Si or oxidized silicon substrates aiming at applications in the future monolithic integrated optics, interconnects, and integrated optoelectronics devices in general [4–7]. Amorphous RE:Al2O3, however, is not very useful for high power bulk solid-state laser applications where crystallinity is necessary to access the highest possible thermal conductivity as well as acceptably high absorption and emission cross sections. The latter are largely due to the narrowband single-crystalline absorption and emission features as opposed to wideband features in amorphous materials with much lower peak cross sections. Successful efforts in producing single-crystalline thin films of Nd3+:Al2O3 grown by plasma-assisted molecular beam epitaxy (MBE) have been recently reported [8]. The Yb:Al2O3 laser heated pedestal growth (LHPG) and Eu:Al2O3 combustion synthesis were presented in [9] and [10], respectively. Though above growth and synthesis methods cannot be seriously considered for volume scaling as required for high power laser applications, these results have confirmed that: (i) single-crystalline sapphire can be doped with RE3+ ions to a usable dopant level, (ii) spectral features of the RE:Sapphire are, as expected, ‘‘crystalline-narrow’’, (iii) the spectra can be consistent with single-site doping of the RE3+ ion into the host crystal, and (iv) RE:Sapphire exhibits sufficient peak transition
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cross sections, exceeding those of RE3+-doped YAG. Nevertheless, currently there is no available technology to produce RE:Sapphire with dimensions required for scalable high energy laser applications. Undoped sapphire can be grown in large sizes by a variety of melt-growth techniques. However, due to the significant ionic radius mismatch between the rare earth ions and the aluminum ion of the host, rare earth elements exhibit negligible solubility within the sapphire crystal structure and tend to not incorporate into the melt-grown crystal. Therefore a non-traditional approach to obtaining single-crystalline bulk RE:Sapphire must be found. It has been demonstrated that bulk ceramics possess the capability to accommodate ‘‘less-friendly’’ dopants and/or significantly higher dopant levels than melt-grown crystals while maintaining sufficient optical quality for laser applications [11]. This indicates that this solid state approach allows for more ‘non equilibrium’ dopant incorporation than traditional melt-growth techniques. In many cases (at least for cubic hosts, e.g., YAG) a highly transparent, low loss, ceramic is a suitable replacement for a single crystal for laser applications. Due to its uniaxial crystal structure, however, laserquality transparency cannot be achieved in polycrystalline a-Al2O3 unless some technique of ceramic grain alignment, like magnetic alignment [12], is used. Additionally, the thermal conductivity of polycrystalline a-Al2O3 is far lower than that of sapphire. Therefore, RE-doped a-Al2O3 ceramics are not very promising for high power laser gain media. On the other hand, RE-doped a-Al2O3 ceramics can potentially be used as a precursor to bulk RE:Sapphire. One possibility would be to apply solid state crystal conversion (SSCC) technique, wherein a ceramic is converted to a single crystal via the growth of a single grain to produce RE:Sapphire. In fact, SSCC has already been used to produce large sapphire crystals [13]. Crystals grown by SSCC have also been demonstrated to be capable of ‘non equilibrium’ dopant incorporation. Therefore the synthesis and spectroscopic characterization of RE-doped a-Al2O3 ceramics may enable the development of single-crystalline RE:Sapphire in dimensions and dopant concentrations adequate for high energy laser applications. This paper reports the results of spectroscopic study of Er3+doped a-Al2O3 (Er:Al2O3) per our first attempt toward transparent bulk Er3+-doped sapphire gain material via ceramic route with the objective to develop a new approach to RE:Sapphire as laser gain medium. We have achieved successful incorporation of sufficient amounts of Er3+ into both highly crystalline a-Al2O3 powders and fully densified transparent a-Al2O3 ceramics. Presented here are the results of detailed spectroscopic investigation of Er3+-doped a-Al2O3 at all development stages. The detailed spectroscopic analysis of the ground state and the first three excited multiplets of Er3+ ions in a-Al2O3 host is presented for the first time. These data can be successfully used for further development of high power Er3+doped a-Al2O3 based lasers once the high quality Er3+-doped sapphire material is available.
2. Experimental Erbium doped Al2O3 powder was synthesized by reverse strike co-precipitation of high purity metal nitrate solution (Er:Nitrate, Al:Nitrate, Mg:Nitrate; all >99.9% purity, Alfa Aesar). The metal nitrates were dissolved in deionized water. The metal nitrate solution was slowly dripped into a buffered bath consisting of deionized water, ammonium bicarbonate, and nitric acid with a starting pH of 7. The pH was continuously monitored and adjusted to pH = 7 using an aqueous solution of ammonium bicarbonate and ammonium hydroxide throughout the precipitation process. Prior to the initial dissolution of metal nitrates, the hydration states of the individual metal nitrates were measured using thermogravimetric analysis to ensure that the Er3+ concentration is 0.026 atom-
ic percent (which corresponds to 0.065 cation percent and yields a formula unit of Er0.0013Al1.9987O3. Magnesium nitrate (500 ppm cation basis) was included during co-precipitation, as Mg was found to help stabilize Er incorporation in the lattice (as determined by single-site Er3+ fluorescence) at sintering temperatures. Following precipitation the suspended power was aged for 24 h while stirring, after which the powder was filtered out and rinsed using deionized water and isopropyl alcohol and finally dried in a vacuum drying over at 75 °C overnight. The amorphous precipitated powder was placed in a high purity alumina crucible and calcined in air at 1300 °C for 1 h in order to form a-Al2O3. The phase purity of the calcined powder was tested using X-ray diffraction (Rigaku Miniflex II). The calcined powder was crushed and ground using a mortar and pestle and sintered in vacuum in a ½ inch inner diameter graphite die using a field assisted sintering technology unit (FAST), which is often referred to as spark plasma sintering (SPS). Samples were sintered at 1200 °C for 5 min as measured by an infrared pyrometer on the exterior of the graphite die. A uniaxial pressure of 50 MPa was applied during sintering. Chemical composition was measured by Rutherford backscattering spectrometry (RBS), using a 2.0 MeV He+ ion beam from a National Electrostatics model 5SDH-2 positive ion accelerator. The backscattering angle was 170°, and the data was interpreted using the SIMNRA simulation software package to yield the average composition over a 2mm diameter target area and through the first few micrometers of material depth from the surface. The ACTON 2500i spectrometer was used to obtain fluorescence spectra in the spectral range of 500–1200 nm, and Nicolet 6700 Fourier Spectrometer was used to obtain fluorescence spectra in the wavelength range beyond 1200 nm. Both spectrometers were equipped with a standard commercial cryogenic dewar (Cryo Industries of America, Inc.) allowing temperature control and measurements in the temperature range of 10–300 K. To achieve low thermal resistance of the pulverized samples (for temperature uniformity) and to minimize the self-absorption effects in Er3+ doped a-Al2O3, the powder samples were immersed in the thermoconductive paste and sandwiched between the two sapphire windows. The thickness of the samples used in all spectroscopic measurements was less than 100 lm. The sandwiched samples were attached to the copper holder and mounted on the cold finger of the cryostat. The spectral resolution of the measurement setup was better than 1 nm at room temperature and 0.2 nm at cryogenic temperatures. Excitation of Er3+ energy levels around 10,000– 12,000 cm1 was typically accomplished by diode lasers emitting at 800 and 980 nm. Occasionally a tunable pulsed Ti-Sapphire laser was used for fluorescence excitation at shorter wavelengths. The spectral correction of the fluorescence spectra taken with both spectrometers has been performed using a standard calibrated quartz-halogen tungsten lamp (Optronic Laboratories, LLC). The majority of samples under investigation were Er3+-doped a-Al2O3 powders. Currently available sintered transparent ceramic samples were less than 0.5 mm thick, therefore the information on Er3+:Al2O3 absorption properties was obtained using the excitation spectroscopy technique. In this case a tunable Continuum Panther EX optical parametric oscillator (OPO) pumped by a SpectraPhysics PRO Series Laser was used as a fluorescence excitation source. The OPO was capable of delivering a few millijoules of short, 10 ns, pulses in the spectral range of 400–1600 nm at 10 Hz repetition rate. In order to compensate for pulse-to-pulse energy fluctuations of the excitation source the excitation spectra were normalized to the reference signal from a photodiode which was directly proportional to the excitation pulse energy. The excitation spectra for all transitions were monitored by observing the fluorescence at 1600 nm corresponding to Er 4I13/2 ? 4I15/2 transition, using InGaAs detector. The optically filtered signal from the detector was averaged by a Stanford Research System Model
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SR 250 boxcar integrator. The overall spectral resolution of the excitation setup was limited by the OPO output spectral width and was found to be better than 0.3 nm across the entire range of the excitation wavelengths. In order to double-check the validity of spectral correction, along with the Er3+:Al2O3 excitation spectra the excitation spectra of Er:YAG, the material well known spectroscopically, excited by the same source have been collected as well. The Er3+:Al2O3 lifetime measurements in the temperature range of 10–300 K were conducted by using 10-ns pulses from tunable Ti:Sapphire laser for excitation and detected by a high speed InGaAs or Si detector attached to the ACTON 2500i monochromator. The detector signal was processed using a TDS 2014 digital oscilloscope. The fluorescence spectrum in most instances was detected by a cooled InGaAs detector combined with a Stanford Research System SR 850 digital lock-in amplifier. The equipment control and data acquisition were integrated into a simple LabVIEW based program. For signal acquisition, equipment control and processing the NI DAQmx USB 6210 device was used in LABVIEW environment. Due to the fact that all spectral measurements were performed on samples which were either powders or sintered ceramics with random grain orientation, the measured fluorescence and excitation spectra were all unpolarized (polarization unresolved), though the Er3+-doped a-Al2O3 is an uniaxial spectroscopic object. This means that all of the presented spectra can be treated as a superposition of the r- and p-polarized spectral components.
3. Results The chemical composition of both the calcined powder and sintered ceramic were confirmed by RBS (both yielding Er3+ concentrations of 0.026 ± 0.003 atomic percent), indicating that all of the Er3+ included during powder synthesis was incorporated into the calcined powder and sintered ceramic. The calcined powder is readily sintered to full density with sub-micron grain size. All the samples (powder and sintered ceramic) obtained under optimized conditions exhibited sharp, well resolved emission lines characteristic of RE3+ ions, also indicating predominant crystalline single-site Er3+ occupation. Obtained IR spectra were similar to those previously observed in Er3+-doped Al2O3 powders prepared using the sol–gel technique in [14].Typical room temperature (RT) 4I13/2 ? 4I15/2 fluorescence spectrum for three randomly
Fig. 1. Room temperature fluorescence spectra of Er3+ doped powders of a-Al2O3, YAG, YAlO3 and Y2O3 (from top to bottom).
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chosen experimental samples of Er3+:Al2O3 with such well defined spectral structure is presented in Fig. 1. All experimental data presented and analyzed in this paper have been obtained from Er3+doped a-Al2O3 powders and transparent ceramics with such well defined, crystalline, highly reproducible, nearly single-site fluorescent spectra. No discernible difference was observed in spectral behavior of properly prepared powders and sintered ceramics. Single-phase nature of each and every sample was carefully monitored by XRD in order to confirm that we were working with the phase-pure a-Al2O3. 3.1. Lifetime measurements of the 4I13/2 manifold Depending on the excitation wavelength, trivalent Er ions are capable of fluorescing in a wide spectral range extending from blue to mid-IR. We have attempted to excite and observe fluorescence from all Er3+ levels below 2H11/2 (6000–19,000 cm1 photon energy range) by using for excitation 10-ns pulses from a tunable OPO (described above) in the range of 1520–1600 nm. We observed fluorescence from 4I13/2 manifold through excitation of 4I9/2, as well as 4F9/2 and 4S3/2 manifolds, but due to the relatively low Er3+ ion concentration and significant nonradiative quenching of all the states above 4I13/2 the quality of obtained fluorescence spectra was insufficient for their sensible spectroscopic analysis. For that reason the data presented here are based on the fluorescence measurements from the excited state with the longest lifetime, 4I13/2 (4I13/2 ? 4I15/2 transition). It was found that the fluorescence decay from the 4I13/2 state is single-exponential, which is consistent with the low Er3+ concentration, when one can assume that inter-ionic interactions between the neighboring Er3+ ions, as well as between the Er3+ ions and possible uncontrollable impurity ions are negligible. Typical fluorescence decay from the 4I13/2 level of Er3+:Al2O3 when excited by 10-ns pulses at 800 nm (excitation via the 4 I15/2 ? 4I9/2 absorption band) is shown in Fig. 2 (inset). The temperature dependence of the 4I13/2 level fluorescence lifetime in the range of 10–300 K is shown in Fig. 2. The slight changes in the lifetime with temperature can be explained by the 4I13/2 population redistribution and decay from the Stark sublevels 4I13/2 with larger probability of spontaneous emission. Much larger values of
Fig. 2. Typical temperature dependence of the Er3+ 4I13/2 state fluorescence lifetime in Er3+ doped a-Al2O3 in the temperature range of 10–300 K. Shown on the inset is the experimental decay curve (red) with the exponential fitting indicated in black. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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Fig. 3. 10-K fluorescence spectrum of Er3+ 4I13/2 ? 4I15/2 transition in Er3+ doped aAl2O3. The arrows indicate the spectral lines corresponding to transitions from the lowest Stark sublevel of the 4I13/2 multiplet to the lowest five sublevels of the 4I15/2 multiplet. Indicated on the left inset (in blue) is the blow-up of the weak spectral feature in the 1620–1700 nm spectral range. Shown in red is the same spectrum averaged by using Savitzky–Golay filter. Indicated on the middle inset is the energy level scheme of the 4I15/2 ground state multiplet inferred from the spectrum. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
the 4I13/2 ? 4I15/2 transition peak cross sections observed at lower temperatures (see Fig. 4) indirectly support this statement.
3.2. 4I13/2 manifold – fluorescence and excitation spectra The emission associated with the 4I13/2 ? 4I15/2 transition of Er3+ ion in a-Al2O3 is of particular practical interest since efficient, high power, eye-safe lasers based on this transition have already been successfully demonstrated in other oxide laser hosts. It was shown that these lasers can efficiently operate at relatively low Er3+ concentrations at both room and cryogenic temperatures [15–17]. The first report on successful doping of a-Al2O3 powder with trivalent Er was presented in [14]. Later on, the authors of [18] have indicated that the absorption and fluorescence spectra observed in [14], in fact, could have originated from the residual phase of the Er3+:YAlO3 compound. Moreover, numerous follow-on studies of Er-doped Al2O3 have shown that the residual phases of ErAlO3, ErAG and Er2O3 compounds can be formed from starting materials as uncontrollable impurities during precipitation or calcination, and they may be responsible for the observed Er3+-like fluorescence in Al2O3 in the IR and visible spectral ranges. It is for that reason, in order to assure that the fluorescence and excitation spectra observed in this study undoubtedly originate from Er3+ ions in aAl2O3, along with testing our own Er-doped a-Al2O3 samples we have also performed the fluorescence and excitation testing of Er3+-doped YAG, Y2O3 and YAlO3 samples. Those samples were acquired as single crystalline and were pulverized for spectroscopic testing in order to make them identical to our own samples in terms of experimental conditions. We assumed Er3+-doped YAG, Y2O3 and YAlO3 samples to be a spectroscopic substitute for ErAG, Er2O3 and ErAlO3. It is well known that, spectroscopically, Er3+doped YAG, Y2O3 and YAlO3 are expected to be nearly identical or at least very similar to those of ErAG, Er2O3 and ErAlO3. Fig. 1 also shows the room temperature Er3+ 4I13/2 ? 4I15/2 emission spectra taken from Er3+-doped a-Al2O3, YAG, Y2O3 and YAlO3 powder samples. This direct comparison indicates that within the accuracy
Fig. 4. Emission cross section of the Er3+ 4I13/2 ? 4I15/2 transition at room (300 K) and cryogenic (77 K) temperatures in Er3+ doped a-Al2O3.
of our measurements there is no manifestation of any residual Er3+-doped YAG, Y2O3 and YAlO3 emissions (lines) in the Er3+doped a-Al2O3 samples used in our work. As seen from Fig. 1 the fluorescence spectrum of our Er3+-doped a-Al2O3 powder is distinct in all of its structural details from the spectra obtained from Er3+-doped YAG, Y2O3 and YAlO3 powders. It also has a unique single sharp emission peak at 1535.4 nm, which the others do not have. The most detailed 4I13/2 ? 4I15/2 fluorescence spectrum at 10 K is shown in Fig. 3. It clearly exhibits the five strong spectral components at 1535.4, 1544.2, 1568.1, 1573.6, 1578.3 nm and the three much weaker components at 1662.2, 1665.8 and 1672.2 nm, which are zoomed in and shown on the right inset in Fig. 3. The energy level diagram of the Stark-split ground state multiplet 4I15/2 of Er3+ ion in a-Al2O3 inferred from the measured fluorescence spectrum at 10 K is presented on the top inset in Fig. 3. For the three highest Stark levels of the energy level diagram, which were inferred from the mentioned three weakest fluorescence transitions, the energy (in cm1) is marked by an asterisk next to it. Due to the low strength and diffuse nature of these emission lines we could not identify these lines as related to the 4 I13/2 ? 4I15/2 transition of Er3+-doped Al2O3 with the 100% certainty. Therefore with the high degree of confidence we can state that we observe five out of eight theoretically expected spectral components corresponding to the inter-Stark transitions from the lowest Stark sublevel of the 4I13/2 multiplet to the first five lowest sublevels of the 4I15/2 ground state multiplet. It should be noted that the energy level diagram inferred from the fluorescence spectra in Fig. 3 substantially differs from that reported in [14]. As mentioned in [18], the energy level diagram derived from the 10 K fluorescence spectra for Er3+ 4I13/2 ? 4I15/2 in [14] can be partially related with the fluorescence of Er3+ in the residual YAlO3 compound which was likely contained in their Er3+:Al2O3 powders. From the measured fluorescence spectrum and the radiative lifetime, s of the 4I13/2 multiplet one can derive the stimulated emission cross section using the well known Fuchtbauer– Ladenburg relationship [19]:
remi ¼
gRk5 8pn2 cs kIðkÞdk
ð1Þ
where remi is the emission cross section, g is the fluorescence quantum efficiency, k is the fluorescence wavelength, n and c are the refractive index and the light velocity, respectively. It should
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be noted that for our cross-section derivation the longest measured lifetime, 7.7 ms, was used. This is based on an assumption that in the samples exhibiting the longest fluorescence lifetime the nonradiative quenching processes are negligible and, thus, the observed 4I13/2 ? 4I15/2 transition lifetime is purely radiative. For the same reason, with this longest lifetime chosen for an estimate, the quantum efficiency g of this transition is assumed to be close to 1. The emission cross section spectrum of the 4I13/2 ? 4I15/2 transition obtained from the above derivation is presented in Fig. 4 for both RT and LNT. While on a relative scale, Fig. 4 indicates that with temperature reduction from 300 K to 77 K the maximum emission cross section is going up by a factor of 6. It can be seen that Er3+ ion in a-Al2O3 host has the cross section values quite adequate for an efficient laser. It should be suitable for high power laser operation based on the 4I13/2 ? 4I15/2 transition. The highest laser power scaling potential is expected at cryogenic temperatures where the observed peak cross section is practically on par with that of RT Nd3+-doped YAG. The excitation spectra corresponding to the 4I15/2 ? 4I13/2 intermultiplet transitions are shown in Fig. 5. Similarly to what was observed in the fluorescence spectra, we found four strong excitation lines at 1509.4, 1516.1, 1522.3 and 1535.4 nm and three much weaker components at 1455.8, 1457.1 and 1459.0 nm. The inferred energy level diagram of the 4I13/2 multiplet is shown on an inset of Fig. 5. For the reasons already mentioned above in relevance to the 4 I13/2 ? 4I15/2 fluorescence spectra, for these weak transitions the corresponding energy levels on the diagram have asterisks next to the level energy. Because at the moment we do not have any transparent samples with thicknesses and concentrations making direct absorption measurements possible it behooves us to compare the obtained excitation spectra for the 4I15/2 ? 4I13/2 transition with the absorption cross section for the same transition. The absorption cross section can be inferred from the Fuchtbauer–Ladenburg derived emission cross section by reciprocity from the following well known expression [20]:
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ration between the lowest Stark sublevels of the 4I13/2 and 4I15/2 multiplets, k, c and k are the Planck constant, light velocity, transition wavelength, T is the temperature of the host, respectively. The calculation results are presented in Fig. 6. As can be seen from Fig. 6, there is a satisfactory agreement between the relative excitation peak strength and the absorption cross sections.
3.3. 4I11/2 manifold – excitation spectra
where rabs and remi are absorption and emission cross sections, ZU and ZL are partition functions for the upper 4I13/2 and the lower 4 I15/2 states, EZL is the zero line photon energy, i.e. the energy sepa-
The energy structure of the 4I11/2 manifold is also of great practical importance since most of Er-based solid state lasers are currently pumped through this level. This manifold provides sufficiently strong absorption in the spectral range of emission of the most efficient commercially available InGaAs/AlGaAs 980-nm semiconductor lasers. Moreover, this level in most oxide laser materials is metastable, with the lifetime varying from the shortest (0.1 ms) in Er:YAG to the longest (3.6 ms) in Er:Y2O3 [21]. Lasers emitting around 3 lm have been built based on a number of Er3+-doped oxide hosts (Er:YAG, Er:YAlO3 Er:Y2O3, etc.), where laser action is initiated from this level, and exploits the nominally self-terminated 4I11/2 ? 4I13/2 laser transition [22]. The fluorescence from this level in our Er3+-doped Al2O3 samples was too weak in order to be used for any sensible interpretation. The high multiphonon decay rate of the 4I11/2 level in this material is due to the fact that the energy gap between this and the next level down (4I13/2) is about 3000 cm1 and the maximum phonon frequency of this host material is 750 cm1 (one of the highest among the oxide laser hosts) [23]. The excitation spectra of the Er3+-doped a-Al2O3, corresponding to the transitions from the ground state multiplet 4I15/2 to all Stark sublevels of the 4I11/2 state are shown in Fig. 7 for the range of temperatures 10–300 K. The excitation spectrum taken at 10 K corresponds to the transitions from the lowest Stark sublevel of the ground state multiplet 4I15/2 to all Stark sublevels of the 4I11/2 state. At this temperature we observed five out of six (expected theoretically) spectral components at 964.2, 965.0, 973.8, 975.2 and 979.5 nm. These 4I15/2 ? 4I11/2 transitions are indicated by arrows in Fig. 7. A very weak spectral component at 10,322 cm1 which can only be discerned after thorough data analysis and noise filtering in the 967–969 nm spectral range can actually be interpreted as a missing line of the 4I15/2 ? 4I11/2 transition at 968.8 nm. The energy level diagram of the 4I11/2 state is presented on the inset in Fig. 7.
Fig. 5. Excitation spectra corresponding to the Er3+ 4I15/2 ? 4I13/2 transitions in the temperature range of 10–300 K in Er3+ doped a-Al2O3. Indicated on the inset is the energy level scheme of the 4I13/2 multiplet inferred from these spectra.
Fig. 6. Emission (from Fuchtbaer–Ladenburgh) and absorption (from emission through reciprocity) cross sections along with excitation spectrum for the Er3+ 4I15/2 ? 4I13/2 and 4I13/2 ? 4I15/2 transitions at room temperature in Er3+ doped a-Al2O3.
rabs
ZU ¼ remi exp ZL
hc k
EZL kT
! ð2Þ
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sublevels of the 4I9/2 manifold. We have observed all five absorption lines (expected theoretically) in the 10-K excitation spectrum. These lines are marked by arrows in Fig. 8. The energy level diagram for the 4I9/2 derived from the 10-K excitation spectrum is shown on an inset in Fig. 8. 4. Conclusions
Fig. 7. Excitation spectra corresponding to Er3+ 4I15/2 ? 4I11/2 transition in the temperature range 10–300 K (from top 10 K to bottom 300 K).
3.4. 4I9/2 manifold – excitation spectra Normally, fluorescence from the 4I9/2 level of Er3+ ion is quenched by severe multiphonon relaxation (in almost all oxides) and only in a few hosts (the so called low-phonon hosts) is it observable. The fluorescence signal from this level in Er3+-doped Al2O3 was also too weak in order to be used for any meaningful interpretation. Nevertheless we have extensively explored the excitation spectra corresponding to the 4I15/2 ? 4I9/2 transition in a wide temperature range. Knowledge of the energy level diagram of this level is also important since the 4I15/2 ? 4I9/2 absorption band of Er3+ is very suitable for pumping with another efficient commercially available InGaAs/AlGaAs semiconductor laser emitting around 800 nm. The excitation spectra of Er3+-doped Al2O3 in the 788– 823 nm spectral range are shown in Fig. 8 for a range of temperatures of 10–300 K. The absorption peaks at 791.7, 792.5, 793.9, 808 and 808.7 nm are identified as corresponding to the 4I15/2 ? 4I9/2 inter-multiplet transitions. The arrows indicate the 10-K transitions from the lowest Stark sublevel of the ground state to all Stark
Fig. 8. Excitation spectra corresponding to Er3+ 4I15/2 ? 4I9/2 transition in the temperature range 10–300 K in Er:Al2O3 (from top 10 K to bottom 300 K).
This work presents spectroscopic evidence that the trivalent Er ions can be successfully doped into a-Al2O3 host despite the significant ionic radius mismatch between dopant and ligand. The observed Er3+ ion spectra in a-Al2O3 nanopowders and fully densified ceramics are a manifestation of purely crystalline environment and also indicate the nearly single-site nature of Er3+ activation. The emission cross section (derived via Fuchtbauer–Ladenburg routine) followed by absorption cross section derived from reciprocity are presented for both room and cryogenic temperatures. Presented detailed spectroscopic study of the ground state and the first three excited Er3+ multiplets (4I15/2, 4 I13/2, 4I11/2, and 4I9/2) in this material can be successfully used for future design and development of high power Er:Sapphire based lasers. The currently planned crystal field calculations for Er3+ ion in an octahedral Al3+ position will answer the question whether the observed spectral behavior supports the hypothesis of Er3+ ion incorporation into Al3+ site in a-Al2O3 single crystal structure. Acknowledgments The authors gratefully acknowledge financial support of this work by the High Energy Laser Joint Technology Office. The authors also thank T. Parker and J. Demaree of U.S. Army Research Laboratory for their assistance in conducting and interpreting the RBS results. References [1] T.Y. Fan, D.J. Ripin, R.L. Aggarwal, J.R. Ochoa, B. Chann, M. Tilleman, J. Spitzberg, IEEE J. Sel. Topics Quantum Electron. 13 (2007) 448–459. [2] S. Tokita, J. Kawanaka, M. Fujita, T. Kawashima, Y. Izawa, Appl. Phys. B80 (2005) 635–638. [3] M.G. Holland, J. Appl. Phys. 33 (1962) 2910–2911. [4] E. Cattaruzza, M. Back, G. Battaglin, P. Riello, E. Trave, Opt. Mater. 33 (2011) 1135–1138. [5] Q. Song, C. Li, J. Li, W. Ding, J. Xu, X. Deng, C. Song, Opt. Mater. 28 (2006) 1344– 1349. [6] R. Serna, S. Nunez-Sanchez, F. Xu, C.N. Afonso, Appl. Surf. Sci. 257 (2011) 5204– 5207. [7] L. Agazzi, J.D.B. Bradley, M. Dijkstra, F. Ay, G. Roelkens, R. Baets, K. Wörhoff, M. Pollnau, Opt. Express 18 (2010) 27703–27711. [8] R. Kumaran, T. Tiedje, S.E. Webster, S. Penson, W. Li, Opt. Lett. 35 (2010) 3793– 3795. [9] J.K. Krebs, U. Happek, J. Lumin. 94 (2011) 65. [10] P.A. Tanner, Z. Pan, N. Rakov, G.S. Maciel, J. Alloys Compd. 424 (2006) 347. [11] A. Ikesue, Y.L. Aung, T. Taira, T. Kamimura, K. Yoshida, G.L. Messing, Ann. Rev. Mater. Res. 36 (2006) 397–429. [12] T. Taira, Opt. Mater. Express 1–5 (2011) 1040–1050. [13] C. Scott, M. Kaliszewski, C. Greskovich, L. Levinson, J. Am. Ceram. Soc. 85 (2002) 1275–1280. [14] A.A. Kaplyanskii, A.B. Kulinkin, A.B. Kutsenko, S.P. Feofilov, R.I. Zakharchenya, T.N. Vasilevskaya, Phys. Solid State 40 (1998) 1310–1316. [15] N. Ter-Gabrielyan, V. Fromzel, M. Dubinskii, Opt. Mater. Express 1 (2011) 503– 513. [16] N. Ter-Gabrielyan, L.D. Merkle, G.A. Newburgh, M. Dubinskii, Laser Phys. 19 (2009) 867–869. [17] S.D. Setzler, M.P. Francis, Y.E. Young, J.R. Konves, E.P. Chiclis, J. Sel. Topics Quantum Electron. 11 (2005) 645–657. [18] P.A. Tanner, K.L. Wong, Y. Liang, Chem. Phys. Lett. 399 (2004) 15–19. [19] B.F. Aull, H.P. Jenssen, IEEE J. Quantum Electron. QE-18 (1982) 925–929. [20] S.A. Payne, L.L. Chase, L.K. Smith, W.L. Kway, W.F. Krupke, IEEE J. Quantum Electron. QE-28 (1992) 2619–2629. [21] T. Sanamyan, J. Simmons, M. Dubinskii, Laser Phys. Lett. 8 (2010) 569–572. [22] V.I. Zhekov, V.A. Lobachev, T.M. Murina, A.M. Prokhorov, Sov. J. Quantum Electron. 13 (1983) 1235–1237. [23] S.P.S. Porto, R.S. Krishnan, J. Chem. Phys. 47 (1967) 1009–1012.
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Optical Materials journal homepage: www.elsevier.com/locate/optmat
Spectroscopic properties of Er3+-doped a-Al2O3 T. Sanamyan a,⇑, R. Pavlacka b, G. Gilde b, M. Dubinskii a a b
U.S. Army Research Laboratory, RDRL-SEE-M, 2800 Powder Mill Rd., Adelphi, MD 20783, USA U.S. Army Research Laboratory, RDRL-WMM-E, Building 4600, Aberdeen Proving Ground, MD 21005, USA
a r t i c l e
i n f o
Article history: Available online 5 December 2012
a b s t r a c t We present the results of comprehensive spectroscopic study of the recently successfully synthesized nanopowders and fully densified transparent ceramics of Er3+-doped a-Al2O3 (corundum). Fluorescence and excitation spectra in a range of temperatures of 10–300 K were used to infer the energy level schemes of the ground state and the first three excited Er3+ multiplets (4I15/2, 4I13/2, 4I11/2, and 4I9/2) in this material. Obtained experimental data also allowed us to assess this material as a potential gain medium for high power laser applications in anticipation that it becomes available in the low-loss grain-aligned ceramic or single crystalline form in the near future. Published by Elsevier B.V.
1. Introduction Despite recent significant advances in power scaling of Y3Al5O12 (YAG) based cryogenically cooled lasers with nearly diffraction limited beam quality [1], including power scaling toward laser fission application [2], the extreme pump power densities combined with the unavoidable non-radiative losses in the pump-lase process introduce severe thermally induced optical distortions in the aggressively pumped and aggressively cooled gain medium. Regardless of the sophistication of heat removal techniques and their efficiency, the insufficient thermal conductivity of the currently used gain media is the bottleneck for non-distortive heat removal. The development of new gain media with extremely high thermal conductivity remains the most critical technological challenge preventing power scaling of bulk high energy solid-state lasers with high beam quality. Due to its unique suite of thermal, optical, and electrical properties, rare earth doped aluminum oxide (RE:Al2O3) is the subject of ever-growing interest as the gain material for microelectronics, integrated on-chip devices for telecommunication, and most recently, for high power laser applications. Indeed, single crystal aAl2O3 (sapphire) has a liquid nitrogen temperature (LNT) thermal conductivity of well over 1000 W/m K [3], which is 40 times higher than the LNT thermal conductivity of YAG (the most prevalent solid state laser host) and approximately two orders of magnitude higher than the room temperature (RT) thermal conductivity of YAG. Due to the mentioned above and the favorable combination of other LNT parameters (dn/dT = 0.19 105 K1, and CTE = 0.34 106 K1) which define the thermal figure of merit ⇑ Corresponding author. Tel.: +1 301 394 2044; fax: +1 301 394 0310. E-mail address: [email protected] (T. Sanamyan). 0925-3467/$ - see front matter Published by Elsevier B.V. http://dx.doi.org/10.1016/j.optmat.2012.10.036
of the cryo-cooled doped a-Al2O3 laser material, it can be pumped far more aggressively with far less thermally induced optical distortion. Thus, a cryo-cooled rare earth doped sapphire (RE: Sapphire) laser would possess extremely high power extraction capabilities, far exceeding those possessed by state-of-the-art solid state laser gain media. To date, RE:Al2O3 devices have been synthesized predominantly as thin films for optoelectronics applications using multiple thin film deposition techniques. In most instances these films are amorphous aluminum oxide layers deposited on Si or oxidized silicon substrates aiming at applications in the future monolithic integrated optics, interconnects, and integrated optoelectronics devices in general [4–7]. Amorphous RE:Al2O3, however, is not very useful for high power bulk solid-state laser applications where crystallinity is necessary to access the highest possible thermal conductivity as well as acceptably high absorption and emission cross sections. The latter are largely due to the narrowband single-crystalline absorption and emission features as opposed to wideband features in amorphous materials with much lower peak cross sections. Successful efforts in producing single-crystalline thin films of Nd3+:Al2O3 grown by plasma-assisted molecular beam epitaxy (MBE) have been recently reported [8]. The Yb:Al2O3 laser heated pedestal growth (LHPG) and Eu:Al2O3 combustion synthesis were presented in [9] and [10], respectively. Though above growth and synthesis methods cannot be seriously considered for volume scaling as required for high power laser applications, these results have confirmed that: (i) single-crystalline sapphire can be doped with RE3+ ions to a usable dopant level, (ii) spectral features of the RE:Sapphire are, as expected, ‘‘crystalline-narrow’’, (iii) the spectra can be consistent with single-site doping of the RE3+ ion into the host crystal, and (iv) RE:Sapphire exhibits sufficient peak transition
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cross sections, exceeding those of RE3+-doped YAG. Nevertheless, currently there is no available technology to produce RE:Sapphire with dimensions required for scalable high energy laser applications. Undoped sapphire can be grown in large sizes by a variety of melt-growth techniques. However, due to the significant ionic radius mismatch between the rare earth ions and the aluminum ion of the host, rare earth elements exhibit negligible solubility within the sapphire crystal structure and tend to not incorporate into the melt-grown crystal. Therefore a non-traditional approach to obtaining single-crystalline bulk RE:Sapphire must be found. It has been demonstrated that bulk ceramics possess the capability to accommodate ‘‘less-friendly’’ dopants and/or significantly higher dopant levels than melt-grown crystals while maintaining sufficient optical quality for laser applications [11]. This indicates that this solid state approach allows for more ‘non equilibrium’ dopant incorporation than traditional melt-growth techniques. In many cases (at least for cubic hosts, e.g., YAG) a highly transparent, low loss, ceramic is a suitable replacement for a single crystal for laser applications. Due to its uniaxial crystal structure, however, laserquality transparency cannot be achieved in polycrystalline a-Al2O3 unless some technique of ceramic grain alignment, like magnetic alignment [12], is used. Additionally, the thermal conductivity of polycrystalline a-Al2O3 is far lower than that of sapphire. Therefore, RE-doped a-Al2O3 ceramics are not very promising for high power laser gain media. On the other hand, RE-doped a-Al2O3 ceramics can potentially be used as a precursor to bulk RE:Sapphire. One possibility would be to apply solid state crystal conversion (SSCC) technique, wherein a ceramic is converted to a single crystal via the growth of a single grain to produce RE:Sapphire. In fact, SSCC has already been used to produce large sapphire crystals [13]. Crystals grown by SSCC have also been demonstrated to be capable of ‘non equilibrium’ dopant incorporation. Therefore the synthesis and spectroscopic characterization of RE-doped a-Al2O3 ceramics may enable the development of single-crystalline RE:Sapphire in dimensions and dopant concentrations adequate for high energy laser applications. This paper reports the results of spectroscopic study of Er3+doped a-Al2O3 (Er:Al2O3) per our first attempt toward transparent bulk Er3+-doped sapphire gain material via ceramic route with the objective to develop a new approach to RE:Sapphire as laser gain medium. We have achieved successful incorporation of sufficient amounts of Er3+ into both highly crystalline a-Al2O3 powders and fully densified transparent a-Al2O3 ceramics. Presented here are the results of detailed spectroscopic investigation of Er3+-doped a-Al2O3 at all development stages. The detailed spectroscopic analysis of the ground state and the first three excited multiplets of Er3+ ions in a-Al2O3 host is presented for the first time. These data can be successfully used for further development of high power Er3+doped a-Al2O3 based lasers once the high quality Er3+-doped sapphire material is available.
2. Experimental Erbium doped Al2O3 powder was synthesized by reverse strike co-precipitation of high purity metal nitrate solution (Er:Nitrate, Al:Nitrate, Mg:Nitrate; all >99.9% purity, Alfa Aesar). The metal nitrates were dissolved in deionized water. The metal nitrate solution was slowly dripped into a buffered bath consisting of deionized water, ammonium bicarbonate, and nitric acid with a starting pH of 7. The pH was continuously monitored and adjusted to pH = 7 using an aqueous solution of ammonium bicarbonate and ammonium hydroxide throughout the precipitation process. Prior to the initial dissolution of metal nitrates, the hydration states of the individual metal nitrates were measured using thermogravimetric analysis to ensure that the Er3+ concentration is 0.026 atom-
ic percent (which corresponds to 0.065 cation percent and yields a formula unit of Er0.0013Al1.9987O3. Magnesium nitrate (500 ppm cation basis) was included during co-precipitation, as Mg was found to help stabilize Er incorporation in the lattice (as determined by single-site Er3+ fluorescence) at sintering temperatures. Following precipitation the suspended power was aged for 24 h while stirring, after which the powder was filtered out and rinsed using deionized water and isopropyl alcohol and finally dried in a vacuum drying over at 75 °C overnight. The amorphous precipitated powder was placed in a high purity alumina crucible and calcined in air at 1300 °C for 1 h in order to form a-Al2O3. The phase purity of the calcined powder was tested using X-ray diffraction (Rigaku Miniflex II). The calcined powder was crushed and ground using a mortar and pestle and sintered in vacuum in a ½ inch inner diameter graphite die using a field assisted sintering technology unit (FAST), which is often referred to as spark plasma sintering (SPS). Samples were sintered at 1200 °C for 5 min as measured by an infrared pyrometer on the exterior of the graphite die. A uniaxial pressure of 50 MPa was applied during sintering. Chemical composition was measured by Rutherford backscattering spectrometry (RBS), using a 2.0 MeV He+ ion beam from a National Electrostatics model 5SDH-2 positive ion accelerator. The backscattering angle was 170°, and the data was interpreted using the SIMNRA simulation software package to yield the average composition over a 2mm diameter target area and through the first few micrometers of material depth from the surface. The ACTON 2500i spectrometer was used to obtain fluorescence spectra in the spectral range of 500–1200 nm, and Nicolet 6700 Fourier Spectrometer was used to obtain fluorescence spectra in the wavelength range beyond 1200 nm. Both spectrometers were equipped with a standard commercial cryogenic dewar (Cryo Industries of America, Inc.) allowing temperature control and measurements in the temperature range of 10–300 K. To achieve low thermal resistance of the pulverized samples (for temperature uniformity) and to minimize the self-absorption effects in Er3+ doped a-Al2O3, the powder samples were immersed in the thermoconductive paste and sandwiched between the two sapphire windows. The thickness of the samples used in all spectroscopic measurements was less than 100 lm. The sandwiched samples were attached to the copper holder and mounted on the cold finger of the cryostat. The spectral resolution of the measurement setup was better than 1 nm at room temperature and 0.2 nm at cryogenic temperatures. Excitation of Er3+ energy levels around 10,000– 12,000 cm1 was typically accomplished by diode lasers emitting at 800 and 980 nm. Occasionally a tunable pulsed Ti-Sapphire laser was used for fluorescence excitation at shorter wavelengths. The spectral correction of the fluorescence spectra taken with both spectrometers has been performed using a standard calibrated quartz-halogen tungsten lamp (Optronic Laboratories, LLC). The majority of samples under investigation were Er3+-doped a-Al2O3 powders. Currently available sintered transparent ceramic samples were less than 0.5 mm thick, therefore the information on Er3+:Al2O3 absorption properties was obtained using the excitation spectroscopy technique. In this case a tunable Continuum Panther EX optical parametric oscillator (OPO) pumped by a SpectraPhysics PRO Series Laser was used as a fluorescence excitation source. The OPO was capable of delivering a few millijoules of short, 10 ns, pulses in the spectral range of 400–1600 nm at 10 Hz repetition rate. In order to compensate for pulse-to-pulse energy fluctuations of the excitation source the excitation spectra were normalized to the reference signal from a photodiode which was directly proportional to the excitation pulse energy. The excitation spectra for all transitions were monitored by observing the fluorescence at 1600 nm corresponding to Er 4I13/2 ? 4I15/2 transition, using InGaAs detector. The optically filtered signal from the detector was averaged by a Stanford Research System Model
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SR 250 boxcar integrator. The overall spectral resolution of the excitation setup was limited by the OPO output spectral width and was found to be better than 0.3 nm across the entire range of the excitation wavelengths. In order to double-check the validity of spectral correction, along with the Er3+:Al2O3 excitation spectra the excitation spectra of Er:YAG, the material well known spectroscopically, excited by the same source have been collected as well. The Er3+:Al2O3 lifetime measurements in the temperature range of 10–300 K were conducted by using 10-ns pulses from tunable Ti:Sapphire laser for excitation and detected by a high speed InGaAs or Si detector attached to the ACTON 2500i monochromator. The detector signal was processed using a TDS 2014 digital oscilloscope. The fluorescence spectrum in most instances was detected by a cooled InGaAs detector combined with a Stanford Research System SR 850 digital lock-in amplifier. The equipment control and data acquisition were integrated into a simple LabVIEW based program. For signal acquisition, equipment control and processing the NI DAQmx USB 6210 device was used in LABVIEW environment. Due to the fact that all spectral measurements were performed on samples which were either powders or sintered ceramics with random grain orientation, the measured fluorescence and excitation spectra were all unpolarized (polarization unresolved), though the Er3+-doped a-Al2O3 is an uniaxial spectroscopic object. This means that all of the presented spectra can be treated as a superposition of the r- and p-polarized spectral components.
3. Results The chemical composition of both the calcined powder and sintered ceramic were confirmed by RBS (both yielding Er3+ concentrations of 0.026 ± 0.003 atomic percent), indicating that all of the Er3+ included during powder synthesis was incorporated into the calcined powder and sintered ceramic. The calcined powder is readily sintered to full density with sub-micron grain size. All the samples (powder and sintered ceramic) obtained under optimized conditions exhibited sharp, well resolved emission lines characteristic of RE3+ ions, also indicating predominant crystalline single-site Er3+ occupation. Obtained IR spectra were similar to those previously observed in Er3+-doped Al2O3 powders prepared using the sol–gel technique in [14].Typical room temperature (RT) 4I13/2 ? 4I15/2 fluorescence spectrum for three randomly
Fig. 1. Room temperature fluorescence spectra of Er3+ doped powders of a-Al2O3, YAG, YAlO3 and Y2O3 (from top to bottom).
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chosen experimental samples of Er3+:Al2O3 with such well defined spectral structure is presented in Fig. 1. All experimental data presented and analyzed in this paper have been obtained from Er3+doped a-Al2O3 powders and transparent ceramics with such well defined, crystalline, highly reproducible, nearly single-site fluorescent spectra. No discernible difference was observed in spectral behavior of properly prepared powders and sintered ceramics. Single-phase nature of each and every sample was carefully monitored by XRD in order to confirm that we were working with the phase-pure a-Al2O3. 3.1. Lifetime measurements of the 4I13/2 manifold Depending on the excitation wavelength, trivalent Er ions are capable of fluorescing in a wide spectral range extending from blue to mid-IR. We have attempted to excite and observe fluorescence from all Er3+ levels below 2H11/2 (6000–19,000 cm1 photon energy range) by using for excitation 10-ns pulses from a tunable OPO (described above) in the range of 1520–1600 nm. We observed fluorescence from 4I13/2 manifold through excitation of 4I9/2, as well as 4F9/2 and 4S3/2 manifolds, but due to the relatively low Er3+ ion concentration and significant nonradiative quenching of all the states above 4I13/2 the quality of obtained fluorescence spectra was insufficient for their sensible spectroscopic analysis. For that reason the data presented here are based on the fluorescence measurements from the excited state with the longest lifetime, 4I13/2 (4I13/2 ? 4I15/2 transition). It was found that the fluorescence decay from the 4I13/2 state is single-exponential, which is consistent with the low Er3+ concentration, when one can assume that inter-ionic interactions between the neighboring Er3+ ions, as well as between the Er3+ ions and possible uncontrollable impurity ions are negligible. Typical fluorescence decay from the 4I13/2 level of Er3+:Al2O3 when excited by 10-ns pulses at 800 nm (excitation via the 4 I15/2 ? 4I9/2 absorption band) is shown in Fig. 2 (inset). The temperature dependence of the 4I13/2 level fluorescence lifetime in the range of 10–300 K is shown in Fig. 2. The slight changes in the lifetime with temperature can be explained by the 4I13/2 population redistribution and decay from the Stark sublevels 4I13/2 with larger probability of spontaneous emission. Much larger values of
Fig. 2. Typical temperature dependence of the Er3+ 4I13/2 state fluorescence lifetime in Er3+ doped a-Al2O3 in the temperature range of 10–300 K. Shown on the inset is the experimental decay curve (red) with the exponential fitting indicated in black. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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Fig. 3. 10-K fluorescence spectrum of Er3+ 4I13/2 ? 4I15/2 transition in Er3+ doped aAl2O3. The arrows indicate the spectral lines corresponding to transitions from the lowest Stark sublevel of the 4I13/2 multiplet to the lowest five sublevels of the 4I15/2 multiplet. Indicated on the left inset (in blue) is the blow-up of the weak spectral feature in the 1620–1700 nm spectral range. Shown in red is the same spectrum averaged by using Savitzky–Golay filter. Indicated on the middle inset is the energy level scheme of the 4I15/2 ground state multiplet inferred from the spectrum. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
the 4I13/2 ? 4I15/2 transition peak cross sections observed at lower temperatures (see Fig. 4) indirectly support this statement.
3.2. 4I13/2 manifold – fluorescence and excitation spectra The emission associated with the 4I13/2 ? 4I15/2 transition of Er3+ ion in a-Al2O3 is of particular practical interest since efficient, high power, eye-safe lasers based on this transition have already been successfully demonstrated in other oxide laser hosts. It was shown that these lasers can efficiently operate at relatively low Er3+ concentrations at both room and cryogenic temperatures [15–17]. The first report on successful doping of a-Al2O3 powder with trivalent Er was presented in [14]. Later on, the authors of [18] have indicated that the absorption and fluorescence spectra observed in [14], in fact, could have originated from the residual phase of the Er3+:YAlO3 compound. Moreover, numerous follow-on studies of Er-doped Al2O3 have shown that the residual phases of ErAlO3, ErAG and Er2O3 compounds can be formed from starting materials as uncontrollable impurities during precipitation or calcination, and they may be responsible for the observed Er3+-like fluorescence in Al2O3 in the IR and visible spectral ranges. It is for that reason, in order to assure that the fluorescence and excitation spectra observed in this study undoubtedly originate from Er3+ ions in aAl2O3, along with testing our own Er-doped a-Al2O3 samples we have also performed the fluorescence and excitation testing of Er3+-doped YAG, Y2O3 and YAlO3 samples. Those samples were acquired as single crystalline and were pulverized for spectroscopic testing in order to make them identical to our own samples in terms of experimental conditions. We assumed Er3+-doped YAG, Y2O3 and YAlO3 samples to be a spectroscopic substitute for ErAG, Er2O3 and ErAlO3. It is well known that, spectroscopically, Er3+doped YAG, Y2O3 and YAlO3 are expected to be nearly identical or at least very similar to those of ErAG, Er2O3 and ErAlO3. Fig. 1 also shows the room temperature Er3+ 4I13/2 ? 4I15/2 emission spectra taken from Er3+-doped a-Al2O3, YAG, Y2O3 and YAlO3 powder samples. This direct comparison indicates that within the accuracy
Fig. 4. Emission cross section of the Er3+ 4I13/2 ? 4I15/2 transition at room (300 K) and cryogenic (77 K) temperatures in Er3+ doped a-Al2O3.
of our measurements there is no manifestation of any residual Er3+-doped YAG, Y2O3 and YAlO3 emissions (lines) in the Er3+doped a-Al2O3 samples used in our work. As seen from Fig. 1 the fluorescence spectrum of our Er3+-doped a-Al2O3 powder is distinct in all of its structural details from the spectra obtained from Er3+-doped YAG, Y2O3 and YAlO3 powders. It also has a unique single sharp emission peak at 1535.4 nm, which the others do not have. The most detailed 4I13/2 ? 4I15/2 fluorescence spectrum at 10 K is shown in Fig. 3. It clearly exhibits the five strong spectral components at 1535.4, 1544.2, 1568.1, 1573.6, 1578.3 nm and the three much weaker components at 1662.2, 1665.8 and 1672.2 nm, which are zoomed in and shown on the right inset in Fig. 3. The energy level diagram of the Stark-split ground state multiplet 4I15/2 of Er3+ ion in a-Al2O3 inferred from the measured fluorescence spectrum at 10 K is presented on the top inset in Fig. 3. For the three highest Stark levels of the energy level diagram, which were inferred from the mentioned three weakest fluorescence transitions, the energy (in cm1) is marked by an asterisk next to it. Due to the low strength and diffuse nature of these emission lines we could not identify these lines as related to the 4 I13/2 ? 4I15/2 transition of Er3+-doped Al2O3 with the 100% certainty. Therefore with the high degree of confidence we can state that we observe five out of eight theoretically expected spectral components corresponding to the inter-Stark transitions from the lowest Stark sublevel of the 4I13/2 multiplet to the first five lowest sublevels of the 4I15/2 ground state multiplet. It should be noted that the energy level diagram inferred from the fluorescence spectra in Fig. 3 substantially differs from that reported in [14]. As mentioned in [18], the energy level diagram derived from the 10 K fluorescence spectra for Er3+ 4I13/2 ? 4I15/2 in [14] can be partially related with the fluorescence of Er3+ in the residual YAlO3 compound which was likely contained in their Er3+:Al2O3 powders. From the measured fluorescence spectrum and the radiative lifetime, s of the 4I13/2 multiplet one can derive the stimulated emission cross section using the well known Fuchtbauer– Ladenburg relationship [19]:
remi ¼
gRk5 8pn2 cs kIðkÞdk
ð1Þ
where remi is the emission cross section, g is the fluorescence quantum efficiency, k is the fluorescence wavelength, n and c are the refractive index and the light velocity, respectively. It should
T. Sanamyan et al. / Optical Materials 35 (2013) 821–826
be noted that for our cross-section derivation the longest measured lifetime, 7.7 ms, was used. This is based on an assumption that in the samples exhibiting the longest fluorescence lifetime the nonradiative quenching processes are negligible and, thus, the observed 4I13/2 ? 4I15/2 transition lifetime is purely radiative. For the same reason, with this longest lifetime chosen for an estimate, the quantum efficiency g of this transition is assumed to be close to 1. The emission cross section spectrum of the 4I13/2 ? 4I15/2 transition obtained from the above derivation is presented in Fig. 4 for both RT and LNT. While on a relative scale, Fig. 4 indicates that with temperature reduction from 300 K to 77 K the maximum emission cross section is going up by a factor of 6. It can be seen that Er3+ ion in a-Al2O3 host has the cross section values quite adequate for an efficient laser. It should be suitable for high power laser operation based on the 4I13/2 ? 4I15/2 transition. The highest laser power scaling potential is expected at cryogenic temperatures where the observed peak cross section is practically on par with that of RT Nd3+-doped YAG. The excitation spectra corresponding to the 4I15/2 ? 4I13/2 intermultiplet transitions are shown in Fig. 5. Similarly to what was observed in the fluorescence spectra, we found four strong excitation lines at 1509.4, 1516.1, 1522.3 and 1535.4 nm and three much weaker components at 1455.8, 1457.1 and 1459.0 nm. The inferred energy level diagram of the 4I13/2 multiplet is shown on an inset of Fig. 5. For the reasons already mentioned above in relevance to the 4 I13/2 ? 4I15/2 fluorescence spectra, for these weak transitions the corresponding energy levels on the diagram have asterisks next to the level energy. Because at the moment we do not have any transparent samples with thicknesses and concentrations making direct absorption measurements possible it behooves us to compare the obtained excitation spectra for the 4I15/2 ? 4I13/2 transition with the absorption cross section for the same transition. The absorption cross section can be inferred from the Fuchtbauer–Ladenburg derived emission cross section by reciprocity from the following well known expression [20]:
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ration between the lowest Stark sublevels of the 4I13/2 and 4I15/2 multiplets, k, c and k are the Planck constant, light velocity, transition wavelength, T is the temperature of the host, respectively. The calculation results are presented in Fig. 6. As can be seen from Fig. 6, there is a satisfactory agreement between the relative excitation peak strength and the absorption cross sections.
3.3. 4I11/2 manifold – excitation spectra
where rabs and remi are absorption and emission cross sections, ZU and ZL are partition functions for the upper 4I13/2 and the lower 4 I15/2 states, EZL is the zero line photon energy, i.e. the energy sepa-
The energy structure of the 4I11/2 manifold is also of great practical importance since most of Er-based solid state lasers are currently pumped through this level. This manifold provides sufficiently strong absorption in the spectral range of emission of the most efficient commercially available InGaAs/AlGaAs 980-nm semiconductor lasers. Moreover, this level in most oxide laser materials is metastable, with the lifetime varying from the shortest (0.1 ms) in Er:YAG to the longest (3.6 ms) in Er:Y2O3 [21]. Lasers emitting around 3 lm have been built based on a number of Er3+-doped oxide hosts (Er:YAG, Er:YAlO3 Er:Y2O3, etc.), where laser action is initiated from this level, and exploits the nominally self-terminated 4I11/2 ? 4I13/2 laser transition [22]. The fluorescence from this level in our Er3+-doped Al2O3 samples was too weak in order to be used for any sensible interpretation. The high multiphonon decay rate of the 4I11/2 level in this material is due to the fact that the energy gap between this and the next level down (4I13/2) is about 3000 cm1 and the maximum phonon frequency of this host material is 750 cm1 (one of the highest among the oxide laser hosts) [23]. The excitation spectra of the Er3+-doped a-Al2O3, corresponding to the transitions from the ground state multiplet 4I15/2 to all Stark sublevels of the 4I11/2 state are shown in Fig. 7 for the range of temperatures 10–300 K. The excitation spectrum taken at 10 K corresponds to the transitions from the lowest Stark sublevel of the ground state multiplet 4I15/2 to all Stark sublevels of the 4I11/2 state. At this temperature we observed five out of six (expected theoretically) spectral components at 964.2, 965.0, 973.8, 975.2 and 979.5 nm. These 4I15/2 ? 4I11/2 transitions are indicated by arrows in Fig. 7. A very weak spectral component at 10,322 cm1 which can only be discerned after thorough data analysis and noise filtering in the 967–969 nm spectral range can actually be interpreted as a missing line of the 4I15/2 ? 4I11/2 transition at 968.8 nm. The energy level diagram of the 4I11/2 state is presented on the inset in Fig. 7.
Fig. 5. Excitation spectra corresponding to the Er3+ 4I15/2 ? 4I13/2 transitions in the temperature range of 10–300 K in Er3+ doped a-Al2O3. Indicated on the inset is the energy level scheme of the 4I13/2 multiplet inferred from these spectra.
Fig. 6. Emission (from Fuchtbaer–Ladenburgh) and absorption (from emission through reciprocity) cross sections along with excitation spectrum for the Er3+ 4I15/2 ? 4I13/2 and 4I13/2 ? 4I15/2 transitions at room temperature in Er3+ doped a-Al2O3.
rabs
ZU ¼ remi exp ZL
hc k
EZL kT
! ð2Þ
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sublevels of the 4I9/2 manifold. We have observed all five absorption lines (expected theoretically) in the 10-K excitation spectrum. These lines are marked by arrows in Fig. 8. The energy level diagram for the 4I9/2 derived from the 10-K excitation spectrum is shown on an inset in Fig. 8. 4. Conclusions
Fig. 7. Excitation spectra corresponding to Er3+ 4I15/2 ? 4I11/2 transition in the temperature range 10–300 K (from top 10 K to bottom 300 K).
3.4. 4I9/2 manifold – excitation spectra Normally, fluorescence from the 4I9/2 level of Er3+ ion is quenched by severe multiphonon relaxation (in almost all oxides) and only in a few hosts (the so called low-phonon hosts) is it observable. The fluorescence signal from this level in Er3+-doped Al2O3 was also too weak in order to be used for any meaningful interpretation. Nevertheless we have extensively explored the excitation spectra corresponding to the 4I15/2 ? 4I9/2 transition in a wide temperature range. Knowledge of the energy level diagram of this level is also important since the 4I15/2 ? 4I9/2 absorption band of Er3+ is very suitable for pumping with another efficient commercially available InGaAs/AlGaAs semiconductor laser emitting around 800 nm. The excitation spectra of Er3+-doped Al2O3 in the 788– 823 nm spectral range are shown in Fig. 8 for a range of temperatures of 10–300 K. The absorption peaks at 791.7, 792.5, 793.9, 808 and 808.7 nm are identified as corresponding to the 4I15/2 ? 4I9/2 inter-multiplet transitions. The arrows indicate the 10-K transitions from the lowest Stark sublevel of the ground state to all Stark
Fig. 8. Excitation spectra corresponding to Er3+ 4I15/2 ? 4I9/2 transition in the temperature range 10–300 K in Er:Al2O3 (from top 10 K to bottom 300 K).
This work presents spectroscopic evidence that the trivalent Er ions can be successfully doped into a-Al2O3 host despite the significant ionic radius mismatch between dopant and ligand. The observed Er3+ ion spectra in a-Al2O3 nanopowders and fully densified ceramics are a manifestation of purely crystalline environment and also indicate the nearly single-site nature of Er3+ activation. The emission cross section (derived via Fuchtbauer–Ladenburg routine) followed by absorption cross section derived from reciprocity are presented for both room and cryogenic temperatures. Presented detailed spectroscopic study of the ground state and the first three excited Er3+ multiplets (4I15/2, 4 I13/2, 4I11/2, and 4I9/2) in this material can be successfully used for future design and development of high power Er:Sapphire based lasers. The currently planned crystal field calculations for Er3+ ion in an octahedral Al3+ position will answer the question whether the observed spectral behavior supports the hypothesis of Er3+ ion incorporation into Al3+ site in a-Al2O3 single crystal structure. Acknowledgments The authors gratefully acknowledge financial support of this work by the High Energy Laser Joint Technology Office. The authors also thank T. Parker and J. Demaree of U.S. Army Research Laboratory for their assistance in conducting and interpreting the RBS results. References [1] T.Y. Fan, D.J. Ripin, R.L. Aggarwal, J.R. Ochoa, B. Chann, M. Tilleman, J. Spitzberg, IEEE J. Sel. Topics Quantum Electron. 13 (2007) 448–459. [2] S. Tokita, J. Kawanaka, M. Fujita, T. Kawashima, Y. Izawa, Appl. Phys. B80 (2005) 635–638. [3] M.G. Holland, J. Appl. Phys. 33 (1962) 2910–2911. [4] E. Cattaruzza, M. Back, G. Battaglin, P. Riello, E. Trave, Opt. Mater. 33 (2011) 1135–1138. [5] Q. Song, C. Li, J. Li, W. Ding, J. Xu, X. Deng, C. Song, Opt. Mater. 28 (2006) 1344– 1349. [6] R. Serna, S. Nunez-Sanchez, F. Xu, C.N. Afonso, Appl. Surf. Sci. 257 (2011) 5204– 5207. [7] L. Agazzi, J.D.B. Bradley, M. Dijkstra, F. Ay, G. Roelkens, R. Baets, K. Wörhoff, M. Pollnau, Opt. Express 18 (2010) 27703–27711. [8] R. Kumaran, T. Tiedje, S.E. Webster, S. Penson, W. Li, Opt. Lett. 35 (2010) 3793– 3795. [9] J.K. Krebs, U. Happek, J. Lumin. 94 (2011) 65. [10] P.A. Tanner, Z. Pan, N. Rakov, G.S. Maciel, J. Alloys Compd. 424 (2006) 347. [11] A. Ikesue, Y.L. Aung, T. Taira, T. Kamimura, K. Yoshida, G.L. Messing, Ann. Rev. Mater. Res. 36 (2006) 397–429. [12] T. Taira, Opt. Mater. Express 1–5 (2011) 1040–1050. [13] C. Scott, M. Kaliszewski, C. Greskovich, L. Levinson, J. Am. Ceram. Soc. 85 (2002) 1275–1280. [14] A.A. Kaplyanskii, A.B. Kulinkin, A.B. Kutsenko, S.P. Feofilov, R.I. Zakharchenya, T.N. Vasilevskaya, Phys. Solid State 40 (1998) 1310–1316. [15] N. Ter-Gabrielyan, V. Fromzel, M. Dubinskii, Opt. Mater. Express 1 (2011) 503– 513. [16] N. Ter-Gabrielyan, L.D. Merkle, G.A. Newburgh, M. Dubinskii, Laser Phys. 19 (2009) 867–869. [17] S.D. Setzler, M.P. Francis, Y.E. Young, J.R. Konves, E.P. Chiclis, J. Sel. Topics Quantum Electron. 11 (2005) 645–657. [18] P.A. Tanner, K.L. Wong, Y. Liang, Chem. Phys. Lett. 399 (2004) 15–19. [19] B.F. Aull, H.P. Jenssen, IEEE J. Quantum Electron. QE-18 (1982) 925–929. [20] S.A. Payne, L.L. Chase, L.K. Smith, W.L. Kway, W.F. Krupke, IEEE J. Quantum Electron. QE-28 (1992) 2619–2629. [21] T. Sanamyan, J. Simmons, M. Dubinskii, Laser Phys. Lett. 8 (2010) 569–572. [22] V.I. Zhekov, V.A. Lobachev, T.M. Murina, A.M. Prokhorov, Sov. J. Quantum Electron. 13 (1983) 1235–1237. [23] S.P.S. Porto, R.S. Krishnan, J. Chem. Phys. 47 (1967) 1009–1012.