Hydrothermal Growth of Single Crystals of Lu3Al5O12 (LuAG) and Its

Apr 23, 2013 - Zhuravleva , M.; Yang , K.; Spurrier-Koschan , M.; Szupryczynski , P.; Yoshikawa , A.; Melcher , C. L. J. Cryst. Growth 2010, 312, 1244...
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Hydrothermal Growth of Single Crystals of Lu3Al5O12 (LuAG) and Its Doped Analogues Cheryl A. Moore, Colin D. McMillen, and Joseph W. Kolis* Department of Chemistry and Center for Optical Materials Science and Engineering Technologies (COMSET), Clemson University, Clemson, South Carolina 29634, United States S Supporting Information *

ABSTRACT: The growth of Lu3Al5O12 (LuAG) as high quality single crystals using the hydrothermal method is described. The growth protocol is similar to that of Y3Al5O12 (YAG) using binary oxides (Lu2O3 and sapphire) as feedstocks and 2 M K2CO3 as mineralizer. The growth conditions are similar to those of YAG with a dissolution zone temperature of 640 °C and a growth zone temperature of 610 °C. To grow large crystals, commercially available YAG single crystals with (100) faces were used as surrogate seed crystals. The close lattice match allows for high quality growth of LuAG single crystals on the YAG seeds. Preparation of Lu2O3 doped with a number of lanthanide ions (Nd, Tm, Er, Yb, Ho) in various concentrations is described, and use of these materials as feedstocks leads to formation of doped LuAG. Similarly, a process for the growth of single crystals of Lu3AlxGa(5−x)O12 with Ga doped in the Al site is also provided. Preliminary absorption spectroscopic data are provided for the various lanthanide doped LuAG crystals. emission to 2.0 μm into the H2O- and CO2-free transmission window for LIDAR applications.9 Lutetium garnet is also a useful host for applications other than lasers. For example, it is an excellent potential host of fast scintillator ions such as Pr3+ and Ce3+, especially for gamma ray detection.10 In addition, LuAG has received substantial interest as a potential high numerical aperture window material for deep UV photolithography. It is one the few materials that is transparent below 190 nm while still maintaining a high refractive index at those wavelengths (2.14), as well as a moderately low intrinsic birefringence (30.1 nm/cm), making it a promising material for 193 nm immersion lithography systems.11 Given the many interesting possible applications for LuAG, we felt it was of interest to explore new routes to single crystals of the material. It is normally grown, similar to YAG, by the classical Czochralski route,12,13 as well as micropulling-down (μ-PD)14,15 and vertical Bridgman methods.16 However, the melting point of LuAG is over 100 °C higher than YAG (2060 °C vs 1930 °C), making melt methods somewhat more problematic.12 In addition, there is evidence of antisite defect issues at these high growth temperatures, where the Lu host ion or a dopant ion occupies the six-coordinate site to a measurable degree in the lattice.17 Finally, the cost of lutetia feedstock is considerably greater than that of yttria, so it is advantageous to

1. INTRODUCTION The use of the lutetium ion as the primary building element in oxide crystal hosts is not especially common, but it does have a number of intriguing potential advantages. Recently, we began to explore some of those applications with the hydrothermal growth of single crystals of the base sesquioxide Lu2O3.1 Lutetia as a host for lasing ions does display some very attractive physical properties,2,3 but its full scale development has been limited by the lack of availability of single crystals due to its exceptionally high melting point (ca. 2500 °C). However, the hydrothermal method provides a potentially useful alternative growth method.4 A more common lutetium-based oxide host is Lu3Al5O12 (LuAG), which is of course the lutetium analogue to Y3Al5O12 (YAG). This material displays many of the positive physical properties of YAG (cubic structure, excellent physical stability, etc.) as well as a number of improved optical properties, making it of interest as host material. In particular, the Lu3+ host ion has nearly the same size and mass as the important heavy lanthanide lasing ions (Yb3+, Er3+, Ho3+, and Tm3+), so there is virtually no lattice strain or defect introduced at any dopant concentration. Also, the similarity in mass of the Lu3+ host to these heavy lanthanide ions means that the decrease in thermal conductivity upon doping is much less than that of similarly doped YAG.5 Subtle differences in optical properties also distinguish LuAG from YAG. For example, LuAG displays increased absorption cross-section for Yb3+ doped crystals, making it an ideal host for high power 1 μm lasing.6−8 Another subtle but important optical property is the red shift of Tm © 2013 American Chemical Society

Received: November 29, 2012 Revised: April 10, 2013 Published: April 23, 2013 2298

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uniformly doped lutetia source material for Nd3+, Yb3+, Er3+, Ho3+, and Tm3+ doping. The growth of lanthanide-doped LuAG then proceeded as described above using this Ln:Lu2O3 starting material. Since LuAG is a much less common laser host material, initial growth experiments were performed using YAG seed crystals because of their somewhat similar lattice parameters (11.91 Å vs. 12.01 Å). Substrate seeds were rectangular slabs (typically on the order of 10 × 5 × 2 mm3 or larger) cut from Czochralski-pulled YAG boules (Scientific Materials Corporation (Bozeman, MT) and MTI Corporation (Richmond, CA)) with top and bottom faces of {100} orientation. Seeds with both {110} and {100} edge orientations were used for hydrothermal growth experiments. Once a reliable growth protocol was determined, the new LuAG growth was wire-sawed off the initial YAG substrate and used as a substrate itself in later experiments. Samples of hydrothermally grown material suitable for absorbance spectrophotometry were cut from the surrogate and given an inspection polish. Absorption measurements were performed on a PE Lambda 900 spectrophotometer scanned from 2200 to 200 nm at a scan speed of 3.7−80 nm/min and a sampling interval of 0.01−0.1 nm. The absence of hydroxide in the hydrothermally grown crystals was observed using infrared spectroscopy. Spectra were collected on asgrown crystals mixed in a KBr matrix using a Nicolet Magna 550 FTIR spectrometer. Data were collected from 400 to 4000 cm−1 under flowing nitrogen. Energy dispersive X-ray analysis (EDX) was used to estimate dopant concentrations on an Oxford INCA EDX analyzer attached to a Hitachi SU6600 scanning electron microscope against a Cu standard. This instrument was also used for elemental mapping (accel. voltage of 30.0 kV and exposure time on the order of 10 min) and linescans. Electron imaging (typically 200−1000× magnification) was performed using a backscatter detector equipped on the Hitachi TM3000 and 3400N instruments. PXRD data was obtained using a Rigaku Ultima IV diffractometer with CuKα radiation (λ = 1.5415 Å) at 0.02° intervals in 2θ at a rate of 1°/min from 5 to 65°. The dopant concentration of Nd:LuAG was determined by neutron activation analysis by Elemental Analysis Inc., Lexington, KY.

develop a growth method that is less profligate of the starting feedstock than is the Czochralski method. We previously demonstrated that the hydrothermal method can be an excellent route to the growth of high quality crystals of refractory oxides.4 Hydrothermal techniques have been successfully used for the growth of a number of garnet phases through the years. Several groups first studied the hydrothermal growth of YIG (Y3Fe5O12) and YGG (Y3Ga5O12) at relatively low temperatures, around 400 °C, in the early 1960s.18−21 Two subsequent studies by Nielsen and co-workers22 and Kolb and Laudise23 demonstrated in concept that YAG could be grown hydrothermally as bulk single crystals using 8 M K2CO3 at 500 °C and that YAG was congruently saturating in hydrothermal solutions. From that point, the hydrothermal growth of YAG was essentially ignored, and no attempts were made to optimize the hydrothermal process for 35 years. We recently however found that the hydrothermal growth method could be applied to grow thick epitaxial layers of doped YAG on oriented YAG seeds.24 By carefully matching the lattice size and growth rate, multiple layers of differentially doped material could be grown. This enables the preparation of multifunctional single crystals with several optical functionalities in a single crystal. Given our success with YAG and the interest in LuAG as described above, we felt it appropriate to investigate the hydrothermal growth of LuAG single crystals as well as some doped congeners. Some successes have been reported recently by Kucera25 and Zorenko26 on the order of a few tens of micrometers using liquid phase epitaxy. The results of our preliminary crystal growth investigations as well as the spectroscopy of a number of doped materials are described herein.

2. EXPERIMENTAL SECTION Undoped LuAG crystals were grown by hydrothermal reactions using a stoichiometric mixture of the oxide components (3 Lu2O3 + 5 Al2O3), whereas, in the doped crystals, the lanthanide dopant was introduced as a predoped oxide component (Ln:Lu2O3) prepared as described later. In a typical reaction Al2O3 (1 g of crushed sapphire windows) and Lu2O3 (2.34 g, HEFA rare earth, 99.997%) were loaded into silver reaction ampules along with 5 mL of aqueous 2 M K2CO3 that serves as a mineralizer, corresponding to a 50−70% fill of the reaction vessel. Hole-drilled target substrate seeds were tied securely using thin silver wire (Alfa Aesar, 0.1 mm dia., 99.997%) to the upper part of a ladder fashioned from thicker silver wire (Alfa Aesar, 1 mm dia., 99.99%) serving to physically stabilize the seed as well as suspend it in the cooler growth zone for the entirety of the reaction. Loaded reaction tubes were weld-sealed and themselves loaded into autoclaves equipped with two exterior ceramic band heaters capable of maintaining a 30 °C thermal gradient between the hotter, dissolution zone (640 °C) at the bottom of the autoclave, and the cooler growth zone (610 °C) at the top of the autoclave. DI water added to autoclaves gives the necessary counter pressure to prevent the ampule from bursting. These temperatures tend to result in internal autoclave pressures between 20 and 35 kpsi. These conditions were maintained for 2−4 weeks followed by either discontinued heat and natural cooling over 12 h or quick quenching by an air stream. For growth of the lanthanide-doped LuAG crystals, predoped lutetia feedstock was prepared by a coprecipitation method from component oxides. In the case of Nd-doping, Lu2O3 and Nd2O3 (HEFA rare earth, 99.997%) were mixed to reach the desired doping level and completely dissolved with hot nitric acid. The solution was cooled and subsequently precipitated in a 50/50 (v/v) solution of 28% NH4OH and water. After being filtered and rinsed with water or a repeated process of settling and decanting, this precipitate was dried at 150 °C then calcined in air at 1000 °C for one day. The resulting crystalline feedstock was confirmed to be Lu2O3 by powder X-ray diffraction (PXRD). This process proved suitable for the preparation of a

3. RESULTS AND DISCUSSION 3.1. Feedstock Development and Initial Growth. The growth of the host material LuAG can be done straightforwardly using Lu2O3 and alumina feedstock with 2 M K2CO3 as a mineralizer at 630 °C and 1−2 kbar. Under these conditions, the phase space seems very stable, and LuAG is the only observed phase under the growth conditions described here. The growth temperature must be maintained above 600 °C because a lower reaction temperature can lead to formation of hydrogarnets or garnet phases that contain OH− in the lattice. Maintaining the growth conditions above 600 °C eliminates this issue, and only anhydrous garnet with no detectable OH− is produced as evidenced by the included (Supporting Information) IR spectrum taken using spontaneously nucleated LuAG material in a pressed KBr pellet. The presence of excess alumina does not seem to have any deleterious effect on the growth or produce other phases. The reaction product is generally phase-pure, and the LuAlO3 perovskite phase was only occasionally observed as a minor product, and only when a very large excess of Lu2O3 was present in the feedstock. This is comparable to the results observed in our previous YAG study.24 Powdered LuAG feedstock can also be mineralized by carbonate mineralizer to recrystallize LuAG in the autoclave, supporting the conclusion that LuAG is congruently saturating in hydrothermal fluids. However, this does not lead to the formation of especially large crystals, even over long reaction periods, and the overall growth of large LuAG crystals is faster under conditions using the binary oxides as feedstock rather 2299

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Figure 1. SEM image of spontaneously nucleated doped LuAG. Smaller faces become more prevalent as nucleations get larger.

of simply using a stoichiometric ratio of dopant oxide powder (e.g., Nd2O3 or Yb2O3) added to undoped Lu2O3 feedstock is much less attractive since it neglects any kinetic or thermodynamic solubility difference of the various metal oxides (or any other useful dopant source including metal chlorides or hydroxides). This creates the risk that the homogeneity of the dopant in the resultant host lattice is reduced. Thus we developed a suitable procedure to form appropriately doped lutetia for use as feedstock. Lutetium oxide power is dissolved in nitric acid along with the appropriate stoichiometry of the dopant oxide. The solution is then precipitated with ammonium hydroxide. PXRD analysis at various steps during the preparation of these doped precursors is shown in Figure 2. As-precipitated powder dried overnight in a 150 °C oven appears amorphous. Continued

than the parent LuAG powder. Work is underway to optimize the transport of true LuAG feedstock to a LuAG seed. The crystals that result from hydrothermal treatment of the various feedstocks at 630 °C in the presence of 2 M carbonate are generally well formed with clearly defined facets as seen by SEM images (Figure 1). The rhombic dodecahedral morphology typical of most garnets is most commonly observed. Spontaneously nucleated crystals are bound by both {110} and {211} crystal forms, with {110} being the more prevalent forms. We do, however, note that the {211} forms tend to become more prevalent in the larger nucleated crystals, and in some cases (Figure 1, bottom right) can be almost as large as the {110} forms. The resulting lattice parameter from single crystal X-ray diffraction of the hydrothermally grown LuAG crystal was found to be 11.909(3) Å. This is essentially identical to the lattice parameters (within the e.s.d.) of single crystals grown by traditional methods (11.906(4) Å) and those derived from powder XRD data (11.912 Å) indicating no significant structural differences between these and the hydrothermal crystals.27,28 More macroscopic crystals up to about 1 mm in size have been obtained by hydrothermal spontaneous nucleation from the component oxides. Since we ultimately want to grow a fairly extensive series of LuAG crystals doped with a variety of heteroatoms and dopant concentrations, we undertook to develop a variety of doped oxide feedstocks. Thus, we report a new predoping protocol to prepare doped lutetia precursors similar to those reported by Wang29 and Li30 but using NH4OH as a precipitating agent instead of carbonate solutions. Preparing predoped lutetia feedstock ensures that in the reactive growth process of LuAG, the host and dopant ions will dissolve at an identical rate. A feedstock that is predoped with a suitable concentration of the desired dopant ion will generate a steady state concentration of dopant in solution throughout any growth process. This should improve the homogeneity of dopant ion distribution in the lattice. The alternative approach

Figure 2. Progression of 6% Nd:Lu2O3 precursor powder pattern against database Lu2O3 (PDF 012-0728) as calcined from 300−1000 °C. Samples were taken after at least 2 h at each given temperature. Crystallization begins between 300−400 °C and fully converts to highly crystalline powder at 1000 °C. 2300

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heating to 1000 °C begins to show crystallization between 300 and 400 °C. Peaks get sharper as the calcination temperature increases due to the larger crystal particle formation. As expected, when the smaller Lu3+ host ion is substituted by the larger Nd3+ dopant ion for example, the resultant expansion in the unit cell is accompanied by a shift toward smaller angles in the pattern, especially as 2θ increases. According to Vegard’s law, this shift should be proportionate to the concentration of doping ions in the host lattice. We do note small shifts in the peak positions based on the resultant expansion of the lutetia unit cell upon doping. It is also possible to prepare a powdered form of doped feedstock of the actual LuAG material itself to be used as powdered feedstock. Thus we can dissolve the appropriate (3 Lu: 5 Al) ratio of components as Lu2O3 dissolved in HNO3 and Al(NO3)3·9H2O in water. Precipitation in NH4OH as described above followed by the same calcination protocol gives a powder pattern consistent with phase-pure, crystalline LuAG powder. It should be noted that the feedstock can be readily doped with the usual trivalent lanthanide ions that are of interest as lasing ions (e.g., Er, Tm, Yb, Nd, Ho as below). It would be useful for a number of applications to be able to prepare a host that can systematically increase the lattice size of the host and/or vary the index of refraction of the host for subsequent preparation of various optical devices.31 Thus we also developed a route to a feedstock that is partially doped with gallium at the aluminum site. We have so far prepared material that is nominally 5−30% Ga substituted at the Al site in LuAG powder. Vegard’s Law suggests that ∼30% Ga doping for Al in LuAG should be sufficient to provide a lattice match to YAG. The powder patterns in Figure 3 clearly demonstrate

our technique a precursor composition for the coprecipitated LuAG feedstock that is roughly 60% Ga would correspond to 30% Ga doping in the resultant powder, offering an exact lattice match to YAG. Gallium doped LuAG was also prepared by the hydrothermal reaction of Lu2O3, Al2O3, and Ga2O3, but for reasons similar to those described above it would be preferable to be able to combine the dopant and host ions in a precursor. The ability to substitute gallium is significant for several reasons. It allows us to match the lattice size of the garnet, which is important for epitaxial growth of various doped garnets on substrates. It is also especially important for tuning of the refractive index of material. This will allow us to grow a variety of epitaxial layers and create waveguides by increasing the refractive index of the gallium-doped layers. Alternatively, the refractive index of a cladding layer can be matched to be similar to the core layer so that the cladding layer can perform various optical functions such as suppression of amplified spontaneous emission.31 The ability to systematically substitute at both the lanthanide and aluminum site of the garnet now enables us to prepare a nearly unlimited combination of lattice sizes, refractive indices, and optical functions. 3.2. Seeded Growth of LuAG on YAG Substrates. In order to form large single crystals of LuAG and doped LuAG for optical applications, transport to a suitable seed is required. This process recently worked well for the growth of doped YAG on YAG seeds.24 Because LuAG offers the hope of possessing better properties than YAG, such as a higher thermal conductivity (especially when doped with laser-active lanthanide ions32), higher density, greater hardness, and possibly a higher damage threshold, it is desirable to be able to grow large single crystals of doped LuAG as we did with YAG. The development of LuAG is in its infancy compared to YAG however, and suitable single crystal seeds are much less available than they are for YAG. One starting point would be to use readily available large, high quality YAG single crystals as surrogate seeds, with the possibility of cleaving any new LuAG growth for further study and use as subsequent seeds. Indeed, we found this approach to be successful, as LuAG can be transported to YAG seeds and grown as high quality epitaxial layers. It appears that the lattice differential between the two garnets (12.01 Å for YAG and 11.91 Å for LuAG) is not sufficient to cause substantial growth problems. We used a thermal gradient of approximately 30° (630−600 °C or 640− 610 °C) over about 10 cm to achieve satisfactory transport from feedstock to seed. This provided a growth rate of roughly 0.3 mm per side per week, similar to that of YAG transport. The growth appears to be of high quality with no visible defect interface between the seed and new growth material shown in the electron micrograph in Figure 4. Element mapping across this interface indicates a clear division between the yttriumcontaining substrate and the lutetium-containing new growth. Oxygen and aluminum levels appear consistent across the interface. As a result of the high quality growth and minimal substrate interface, the fully faceted as-grown crystals (with sizes of 5−15 mm) often have a gem-like appearance. This is accentuated by the presence of {211} and {210} crystal forms in addition to {110} forms adding a number of degrees of internal reflection. In these bulk hydrothermally grown LuAG crystals, morphology termination often occurs by way of adjoining opposing {210} faces, which is typically a relatively minor face of importance in garnets. The {210} faces were also observed in the YAG study, though they are much more prominent in the

Figure 3. X-ray diffraction powder patterns of coprecipitated YAG, “30%Ga:LuAG”, and LuAG. Gallium doping shifts peaks from LuAG closer to YAG. Inset shows high angle regions where the shift is most apparent.

some degree of substitution of Ga for Al by the contraction of the 2θ values due to the increase in lattice parameter from those of LuAG, although not to the point where an exact lattice match with YAG was obtained. EDX analysis of the “30%” gallium doped LuAG feedstock indicates that actual resultant doping is closer to 15%, which would explain why we observe the peak positions of “30% Ga:LuAG” to still be intermediate of YAG and LuAG, but more closely positioned to those of YAG. However, it is clear we can systematically vary the garnet lattice size by the degree of Ga substitution. We estimate that using 2301

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Figure 4. Elemental map of the interface of LuAG grown on YAG shows segregation of yttrium and lutetium in their appropriate regions.

{100} orientation. We were also able to observe growth of a wide variety of dopants in LuAG on YAG seeds (vida infra). One interesting result is the clear demonstration of LuAG with high gallium substitution for aluminum. An element mapping pattern across one interface of the cross section of 30% Ga:LuAG grown on YAG clearly displays the substitution and oriented growth (Figure 6). Gallium and lutetium are confined to the outer regions of the cross section (the new Ga:LuAG growth, on the right-hand side of each image), and this is accompanied by an aluminum deficiency in this same region relative to the central region (the YAG seed crystal, on the lefthand side of each image). However, we also notice that the diameter of the seed crystal has decreased to about 1.4 mm compared to its original thickness of 2 mm, indicating a significant amount of back-dissolve has taken place during the warm-up process. This accounts for the roughness visible at the edges and surface of the crystalline interface. Also, a small, but detectable amount of yttrium is present in the new growth such that the actual composition resembles Lu2.59Y0.41Al3.68Ga1.32O12. This is much more significant back-dissolve than we have observed when LuAG or Ln:LuAG is grown on the YAG substrate and is likely attributable to the addition of Ga2O3 to the feedstock in this experiment. The addition of Ga2O3 (which itself is highly soluble in the hydrothermal fluid) must change the solubility behavior of YAG to make it much more soluble in the carbonate mineralizer so that nearly 30% of the seed crystal dissolved during the warm-up process. This increased solubility also affected the growth rate of the crystal, as this particular crystal grew at a rate of 1.37 mm/side/week, nearly 4.5 times faster than that of undoped LuAG grown under the same conditions. The faster growth rate seems to have resulted in some inclusions and cracks, and the process will require further optimization. Thus, an important step in future growth studies to reduce the seed crystal dissolution and control the growth rate and quality will be developing a technique where the Ga:LuAG powders described in the previous section can be mineralized and transported directly, thus mitigating the mineralizing effects of Ga2O3 in a reactive transport scheme. 3.3. Growth of Doped Nd3+, Yb3+, Er3+, Ho3+, and Tm3+:LuAG Single Crystals. One purpose of this investigation was to develop a convenient route to LuAG crystals doped with a wide variety of useful lasing ions in a number of different concentrations. We began our study of doped LuAG with Nd3+, which has several features of interest. It is certainly a very common lasing ion, but it is also one of the lightest and hence largest lanthanide ions. It is significantly larger than the

LuAG system. The ability to grow crystals of this size and quality is useful because the new growth is large enough to be easily cut away from the original seed and polished to create a new [100] oriented seed of sufficient size and quality to be used as a seed for follow-on growth reactions or used for subsequent optical evaluation. The optical clarity of such cleaved and polished LuAG slabs is apparent in Figure 5. Absolute

Figure 5. (a) Fully faceted LuAG single crystals, in the top case, with the silver mounting wire still clearly visible and (b) LuAG crystals harvested from YAG surrogate seeds.

temperature and temperature gradient both play a role in growth rate and quality. In general, a slower growth rate (smaller ΔT) leads to higher quality material, particularly in the early stages of the growth process. We have not yet performed a systematic investigation of ΔT or absolute growth temperatures. Because the temperature and heating protocol used for growth thus far was adopted directly from our accepted protocol for YAG, it is also possible that this might not be the ideal heating protocol for LuAG systems and further refinement might lead to faster growth and still achieve the optimum clarity and quality of growth. The best quality of growth material seemed to occur on welloriented (100) faces, particularly if the edges are also cut to be 2302

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Figure 6. Element map of 30%Ga:LuAG grown on a YAG seed. Map shows substitution of Ga into the Al sites.

host Lu3+ ion and is also larger than the other most common lasing ions Yb3+, Tm3+, Er3+, and Ho3+, all of which have similar ionic radii to the host Lu3+. Thus, we felt that if we could grow high quality single crystals of LuAG doped with significant quantities of Nd3+, we would anticipate not having problems with all of the remaining smaller ions. Growths were performed using several Nd doping levels ranging from 1 to 4% as initial concentration of Nd:Lu2O3 powder as feedstock. This led to Nd doped LuAG single crystal layers grown as described above on YAG. Resultant Nd doped LuAG was cut away from the surrogate YAG seed and given an inspection polish. The resultant products were high quality single crystals with no visible defects and provided sharp Nd3+ absorption spectra as shown in Figure 7. Crystals grown by other methods give

3% Nd in the smaller LuAG lattice is significant. The implication of this large difference in maximum possible dopant is not yet known in terms of lasing behavior but may lead to the ability to prepare smaller laser cavities that provide the same power levels. Additional tests are underway to maximize the amount of Nd3+ capable of being substituted into the LuAG lattice. Detailed spectroscopic examination of this system is underway, and the results will be reported shortly.37 Doping with Yb3+ presents some particular challenges. Once again we can grow single crystals with a wide range of dopant ion concentrations from 3 to 10% Yb. In this case, the dopant ion concentration can be very high because of the identical size of the dopant and host ions. However, Yb is unique because of the well-documented38−42 tendency of colorless Yb3+ to become reduced to blue Yb2+ in situ during melt based growth reactions. We observe similar behavior during hydrothermal growth with visible layers of bluish Yb doped LuAG. The absorption spectrum of the blue material indicates that not all of the Yb has been reduced to divalent state. Solid state absorbance spectrophotometry of the 4% Yb3+:LuAG slab shows the typical peaks for Yb3+ in the NIR around 917, 940, 968, and 1031 nm. The absorption spectrum also shows the presence of strong, broad peaks at ∼375 and 600 nm in addition to the standard Yb3+ peaks (900−1050 nm), and this can be attributed to Yb2+.39−42 Annealing in a ceramic crucible in an open-air 1000 °C furnace for 1−2 h is sufficient to oxidize all the Yb2+ back to its colorless Yb3+ state (Figure 8). We are not clear as to the source of the reducing agent in our closed growth system, but we do observe that most of the blue color occurs at the initial interface and final surface of the growth. This might suggest that the reduction in our case occurs during the thermal transition at the beginning and end of the growth period. One important observation that can be made is the relative intensity of the peak near 970 nm versus that of the manifold around 940. In the more well-known case of Yb:YAG, the intensity of the peaks at 940 are substantially greater (ca. 2:1) than that at 970. However, in the Yb:LuAG case, the relative intensities are approximately equal. This may have important ramifications for minimizing the quantum defects in high power lasers by improved pumping efficiency at 970 rather than the more common 940 nm. The other common lasing ions can be introduced in LuAG host lattices as well using our methods. In the case of thulium doping, high quality single crystals of 5%Tm:LuAG were grown. Absorption spectroscopy done on polished parallel flats

Figure 7. Absorption spectrum of 2%Nd:LuAG from 300 to 900 nm. Inset shows detail of the region of 780−830 nm.

similar absorption spectra14,33−35 with four clusters of peaks at 525, 590, 740, and 810 nm. Peaks are slightly sharper in comparison to Nd:YAG,36 while the intensity of the cluster centered near 590 nm is significantly increased in relation to the others. In the case of the highest nominal Nd ion concentration we attempted (4%), detailed spectroscopic calculations imply that the actual doping levels are closer to 2.6%.37 We corroborated that estimated concentration by data obtained by neutron activation analysis (2.7 ± 0.3%.) This is interesting because it is substantially greater than the maximum doping concentration attainable in Nd:YAG crystals (∼1%) grown by the Czochralski method. Given that the lattice size of LuAG is smaller than that of YAG, the ability to insert nearly 2303

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Figure 10. Absorption spectrum of 5%Ho:LuAG. Right inset details the pumping region between 1700 and 2200 nm of interest for 2.1 μm lasing. Left inset shows the other near-IR transitions.

Figure 8. Absorption spectrum comparison between Yb2+/3+ vs Yb3+:LuAG. Top trace: Before annealing. Bottom trace: After airannealing crystal becomes colorless as seen from the absence of the intense peaks in the visible region. Inset shows detail of the region of 875−1050 nm.

Another improvement of this system over Ho:YAG is a result of the greater energy level splitting in LuAG due to its shorter metal oxide bond distances due to the smaller lattice size. Because the next Stark level above the ground state is slightly elevated in energy relative to YAG, it is slightly less thermally populated and gives a more efficient emission. Walsh et al. present a thorough treatment of the complex energy level determination of Ho3+ in LuAG,45 which includes at least 49 different manifolds, and offer some improvements of those previously reported46,47 in YAG.48,49 Our absorption spectra from hydrothermal crystals agree with others reporting Ho:LuAG.45,50,51 Three sets of similarly spaced low-intensity peaks occur between 280 and 425 nm and another more intense set of three groups between 440 and 650 shows the sharper transitions common to f-electron states. The pumping region between 1700 and 1950 nm can be seen to consist of a number of closely spaced, fairly intense peaks. The absorption spectrum of 30% Er:LuAG is presented below in Figure 11. Erbium has a complex absorption and

of 5%Tm:LuAG gives the spectrum shown in Figure 9 with the characteristic sharp peak at ∼690 nm, clusters at 790 nm

Figure 9. Absorption spectrum of 5%Tm:LuAG. Inset shows detail of the region of 1600−2000 nm.

(pumping region for LIDAR applications), 1200 nm, 1600− 1800 nm, and a low-intensity broad group between 1800 and 2000 nm. Spectroscopic analyses of similarly doped crystals in both LuAG and YAG grown by Czochralski method by Kmetec et al.9 and later by Kalaycioglu43 and Wu44 give comparable spectra, and the transitions can be assigned as in the literature. The sharp, intense line at 682 nm is four times as intense as any other absorption peak and attributed to the transition of the ground state 3H6 of Tm3+ to the 4F2/4F3 excited states. Much less intense (about 20%) are the multiple peaks around 775 and 1200 nm due to the 3H4 and 3H5 transitions, respectively. The broad manifolds from 1750 to 1950 are the transitions from 3F4 and are the transitions of interest in the 2-μm lasing behavior of Tm3+. The absorption spectrum of 5%Ho:LuAG is presented below in Figure 10. Ho doped garnets are also used for eye-safe applications utilizing their emission in the 2.1 μm range. Pumping into the three-peak cluster at 1.9 μm gives lasing at ∼2.1 μm corresponding to the 5I7 excited state to ground state, 5 I8.

Figure 11. Absorption spectrum 30%Er:LuAG. Inset shows detail of the 1450−1650 nm region.

emission behavior, and typically for lasing applications the desire is for relatively small doping levels (0.25−0.4%) for 1.5μm emission ( 4I 13/2 to 4 I15/2), or very high (≥30%) concentrations for 3-μm lasing resulting from the 4I11/2 to 4 I13/2 transition. In our case, we grew a highly doped sample with the understanding that a smaller concentration is also easily obtainable. In the case of the 30% doped sample below, 2304

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the spectrum is quite well resolved, and the significant peaks near 1500 nm are strong and resolved. Er3+ absorption spectroscopy is characterized by clusters surrounding a set of five sharp, regularly spaced peaks of decreasing intensity and a silent region stretching from ∼1000 to 1425 nm. Setzler et al. give a good analysis of Er-doping in YAG in regard to its eye-safe applications,44 and our spectra produced by hydrothermally grown crystals are similar to theirs and others reporting Er:LuAG absorbance spectra.52

ASSOCIATED CONTENT

S Supporting Information *

Infrared spectrum of hydrothermally grown LuAG (Figure S1) is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

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4. CONCLUSIONS We demonstrate for the first time that high quality single crystals of LuAG can be grown using hydrothermal methods similar to those for YAG growth. The LuAG crystals can be grown directly under conditions similar to those for YAG, using reactions of the parent oxides between 600 and 650 °C in water with 2−8 M K 2 CO 3 as mineralizer. In addition to spontaneously nucleated small single crystals, the feedstock can be transported to oriented single crystals of either LuAG or YAG seeds. With a slow growth process, the lattice mismatch with YAG is easily overcome, and high quality LuAG crystals of several millimeters in thickness can be prepared this way. In general, the quality of growth is high with clean faceting in the products. The growth rate is similar to that of YAG, and the relative rate of face growth is also similar to that of YAG under comparable hydrothermal conditions. In addition, a variety of dopants can be introduced into the lattice using new coprecipitation techniques followed by hydrothermal crystallization, or by direct hydrothermal reaction of oxide components. These include the usual selection of lanthanide trivalent ions such as Yb3+, Tm3+, Er3+, and Ho3+. Since these have a size similar to the host Lu3+ ions, they can be substituted at very high concentration with little or no strain. In the case of Yb, we sometimes observe partial reduction to Yb2+ but find that it can easily be converted to Yb3+ upon simple postgrowth heating. We have also been able to introduce the larger Nd3+ ion in fairly high concentrations (2.7%) and obtain good absorption spectra. We also find that the Al3+ ion can be substituted by Ga3+ in the lattice in a rational fashion by including small amounts of Ga as Ga2O3 in the feedstock for hydrothermal growth, as well as utilizing the above-mentioned coprecipitation techniques. This provides a mechanism to tune both the lattice size as well the index of refraction of the lattice, which opens the door for future epitaxial growth and waveguide production.

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Article

AUTHOR INFORMATION

Corresponding Author

*Phone: 864-656-4739. Fax: 864-656-6613. E-mail: kjoseph@ clemson.edu. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We are grateful to the National Science Foundation for financial support of this work (Grant #DMR-0907395). 2305

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