J. Am. Ceram. Soc., 94 [5] 1380–1387 (2011) DOI: 10.1111/j.1551-2916.2010.04260.x r 2010 The American Ceramic Society Journal Effect of SiO2 on Densification and Microstructure Development in Nd:YAG Transparent Ceramics

Adam J. Stevenson,w,z Xin Li,z Miguel A. Martinez,z Julie M. Anderson,z Daniel L. Suchy,y Elizabeth R. Kupp,z Elizabeth C. Dickey,z Karl T. Mueller,y and Gary L. Messingz zDepartment of Materials Science and Engineering and Materials Research Institute, The Pennsylvania State University, University Park, Pennsylvania 16802

yDepartment of Chemistry, The Pennsylvania State University, University Park, Pennsylvania 16802

This paper examines the influence of SiO2 dopingondensification There is substantial interest in transparent Nd:YAG and other and microstructure evolution in Nd3xY33xAl5O12 (Nd:YAG) ceramic laser hosts (Sc2O3,Y2O3,Y3A5O12) doped with optically ceramics. Nd:YAG powders were doped with 0.035–0.28 wt% active ions (Nd31,Er31,Yb31,Tm21,Ho41).6–11 Where direct 29 SiO2 and vacuum sintered between 14841 and 17501C. Si comparisons are possible, ceramic laser hosts have similar optical magic-angle spinning nuclear magnetic resonance showed that properties, lasing threshold, and lasing efficiency as melt-grown Si41 substitutes onto tetrahedrally coordinated Al31 sites. High- single crystals.6 At high laser powers, Nd:YAG ceramics may resolution transmission electron microscopy showed no grain show higher total output power than Czochralski grown single boundary second phases for all silica levels in samples sintered crystals because ceramics do not contain certain defects inherent at 16001–17501C. Coarsening was limited by a solute drag mech- to the Czochralski process.12,13 anism as suggested by cubic grain growth kinetics and transmis- Transparent Nd:YAG ceramics are fabricated by sintering sion electron microscopy energy-dispersive X-ray spectroscopy Y2O3,Al2O3,andNd2O3 powders (solid-state reactionr (SSR)) observations of increased Nd31 concentration near grain bound- or solution-derived Nd:YAG powders. Ikesue et al.6 showed aries. Increasing SiO2 content increased both densification and that doping an SSR system with 0.14 wt% SiO2 was critical to grain growth rate and led to increasingly coarsening-dominated achieving transparency and enabled laser slope efficiencies sim- sintering trajectories. Fine-grained (o3 lm), highly transparent ilar to Czochralski grown crystals. All subsequent literature re- (482% real in-line transmission) ceramics were produced by ports about sintering transparent Nd:YAG ceramics (both SSR sintering 0.035 wt% SiO2-doped ceramics at 17501Cfor8h. and YAG powder approaches) use SiO2 as a sintering aid. Coarse-grained (18 lm), transparent samples were obtained with Fine-grained (1–5 mm) ceramic laser host materials are 0.28 wt% SiO2-doped Nd:YAG when sintered at 16001Cfor8h. desirable for high power applications because they have higher strength and thus can perform at higher maximum output powers than coarse-grained (410 mm) ceramics. Finer grain sizes also improve laser performance by mitigating thermally I. Introduction induced optical birefringence effects between grains, thus 14 OBLE first demonstrated that polycrystalline ceramics could enabling temporally stable laser output. Cbe sintered to a visually translucent/transparent state.1 Grain sizes between 10 and 50 mm are common to SSR syn- Peelan and Metselaar2 showed that the primary scattering sites thesized Nd:YAG as a result of the 417251C sintering temper- 6,10,11 15 in translucent alumina were pores and highly transparent ce- ature required to achieve full density. In 2002, Lu et al. ramics could only be achieved by eliminating residual porosity. reported highly transparent, laser quality Nd:YAG by sintering Since then, a variety of transparent materials including Yttralox coprecipitated powders. These ceramics (available through 3 4 5 (90%Y2O3–10%ThO2), MgAl2O4, and AlON have been suc- Konoshima Chemical Inc., Osaka, Japan) are 2 mm grain cessfully produced by sintering to exceptionally high densities size and demonstrate lasing efficiency and output power similar 8 16 (i.e., 499.9%) and thus reducing/eliminating light scattering to single crystals. Recently, Lee et al. demonstrated fine- from residual porosity. In many cases, pressure assisted densifi- grained (2–3 mm), transparent Nd:YAG densified by a sinter- cation (hot pressing (HP) and/or hot isostatic pressing (HIP)) is HIP process between 16001 and 17501C. By the HIP approach, used to eliminate the final few pores in transparent ceramics, but Lee and colleagues showed that silica content could be reduced contamination and cost concerns inherent to HP and HIP make to 0.04 wt% while still producing ceramics with in-line trans- pressureless sintering the preferred process when possible. While mission 482%at1064nm(thelasingwavelengthinNd:YAG). there are many studies about final stage densification and mi- Thermodynamic modeling of the SiO2–Y2O3–Al2O3 system crostructure evolution in ceramics, few studies have concen- shows that SiO2 forms a liquid phase with YAG at temperatures 17 trated on sintering and microstructure development at the 414001C. However, SiO2 solid solution in the YAG lattice was 17 exceptionally high densities required for transparent ceramics not included in the thermodynamic model. Boulesteix 18 (499.9%). In this paper, we examine these issues for et al. used densification kinetics and scanning electron micros- Nd3xY33xAl5O12 (Nd:YAG) ceramics, a material system of copy (SEM) observations to determine that a liquid phase does current interest for advanced solid-state lasers. develop in 0.3 wt% SiO2-doped YAG ceramics sintered between 14001 and 15501C. However, high-resolution transmission elec-

A. Krell—contributing editor tron microscopy (HRTEM) images of grain boundaries in trans- parent YAG and Nd:YAG samples sintered between 16751 and 17501C and cooled at 201C/min revealed no evidence of an amorphous phase at grain boundaries or triple lines.16,19 Manuscript No. 28190. Received June 16, 2010; approved October 17, 2010. To understand how SiO2 interacts with the YAG lattice, first This work was supported by the United States National Science Foundation grant number DMR749391. consider the following crystal features. Y3Al5O12 has a cubic w 10 Author to whom correspondence should be addressed. e-mail: [email protected] crystal structure (space group Ia3d or Oh )with160atoms(eight 1380 May 2011 Effect of SiO2 on Nd:YAG Transparent Ceramics 1381 formula units) per unit cell and lattice constant 5 12.000 A˚.Y31 (0.1019 nm) occupies dodecahedrally coordinated 24(c)sites. Al31 occupies two crystallographically distinct sites: octahed- rally coordinated 16(a) sites (0.0535nm) and tetrahedrally coordinated 24(d) sites (0.039 nm). O2 occupies the 96(h)sites that make up the vertices between the cation polyhedra. Each O2 is shared between two Y31 dodecahedra, one Al31 tetrahedron, and one Al31 octahedron. Nd31 (0.1109 nm) substitutes for Y31 onto the dodecahedrally coordinated 24(c) sites resulting in a lattice strain due to its larger ionic radius. Several defect mechanisms have been proposed for SiO2 solubility in YAG. Based on atomistic modeling, Kuklja predicted that the most likely defect mechanism for Si41 (0.026 nm) solubility in YAG is

5 1 1 1 SiO þ Alx þ Y x ! Si0 þ V00 þ Y Al O (1) 2 6 Al 2 Y Al 3 Y 6 3 5 12 where the preferred substitutional site is the tetrahedrally coordinated 24(d )site.20 This seems logical due to the similar Fig. 1. FESEM micrograph of 1 at.% Nd:YAG powder used in this study calcined at 11501C for 2 h and 12501Cfor10h. ionic radii of tetrahedrally coordinated Si41 (0.026 nm) and the 24(d ) tetrahedral Al31 sites (0.039nm) in YAG. In this article, we determine the local chemical environment cold isostatic pressing (CIP) at 200 MPa. Samples denoted as 41 ‘‘tape cast’’ were tape cast according to the process detailed surrounding Si in SiO2-doped Nd:YAG ceramics and discuss elsewhere.16 Organics were removed by heating at 0.21C/min to how SiO2 doping affects mass transport mechanisms, densifi- cation, and grain growth between 16001 and 17501C. Further, 6001C and held 12 h in air. After burnout, samples were CIPed we demonstrate how sintering trajectory is affected by changing at 200 MPa. SiO2 content and how this information can be used to obtain either coarse-grained (410 mm) or fine-grained (o3 mm) highly (2) Sintering transparent Nd:YAG ceramics. We also discuss the influence of Samples were sintered in a pure YAG embedding powder that green forming on sintering trajectory and the ability to produce hadbeencalcinedinairat16001C for 40 h. The embedding fine-grained, highly transparent Nd:YAG ceramics. Coprecipi- powder contained no SiO2. Sintering occurred between 14841 tated Nd:YAG powders are used in this study, however the and 17501C for up to 60 h in a tungsten mesh heated vacuum results are applicable to SSR derived Nd:YAG ceramics as well furnace (M60, Centorr Vacuum Industries, Nashua, NH) under because phase formation in SSR powders is complete by 5 103 Torr. The heating and cooling rates were 101 and 201C/ 1 1500 C. min, respectively.

II. Experimental Procedure (3) Microstructural and Optical Characterization (1) Powder Synthesis and Green Forming Sintered density up to 99.9% was measured using the Archime- des method. Sintered densities over 99.9% were obtained by Nd:YAG powders were synthesized by reverse strike coprecip- image analysis. For SEM imaging (XL20, Philips, Eindhoven, itation of 1.5M aqueous solutions of Al(NO3)3 9H2O(981%, the Netherlands), the samples were ground with a diamond Sigma-Aldrich, St. Louis, MO), YCl3,andNd(NO3)3 nH2O wheel and then polished sequentially on silk pads with 6, 1, and (99.9%, Alfa Aesar, Ward Hill, MA). 1.5M YCl3 was prepared 0.25 mm diamond slurries. Sample thicknesses after polishing by dissolving Y2O3 powder (99.99%, Yarmouth Inc., Danbury, were nominally 1 mm. The polished specimens were thermally CT) in HCl (J.T. Baker, Phillipsburg, NJ). The salt solution was etched between 13001 and 15501C for 0.25–10 h. Grain sizes and added dropwise to a rapidly stirring aqueous solution of 0.45M pore sizes of the sintered samples were obtained by the linear (NH4)2SO4. The precipitation pH was maintained at 9.2 with intercept method (at least 200 counts) where the average inter- NH4OH. After aging for 1 h, the precipitate was repeatedly cept length was multiplied by 1.56 to calculate the average grain washed and centrifuged with 0.05M (NH4)2SO4 and a final size. In-line optical transmission on sintered samples annealed in wash was performed using ethanol. The washed precipitate air for 10 h at 16001C was measured with a spectrophotometer was dried for 16 h at 801C and crushed in a mortar and pestle. with integrating sphere (Lambda 950, Perkin-Elmer, Wellesley, The coprecipitated powder was calcined at 11501Cfor2hto MA). form the garnet phase and subsequently calcined at 12501Cfor 12 h to increase the average particle size to 200 nm (Fig. 1). Each powder particle is composed of two to three 80 nm crystallites. (4) Nuclear Magnetic Resonance (NMR) All powders used in this study were 1at.% Nd:YAG, and will be NMR was used to determine the coordination and bonding referred to simply as Nd:YAG. environment of Si41 in sintered Nd:YAG ceramics. Coprecipi- The Nd:YAG powders were mixed with different amounts of tated Nd:YAG powders were mixed with 0.14 wt% and 29 29 tetraethoxysilane (TEOS, 99.9999%, Alfa Aesar) to obtain 0.28 wt% Si-enriched SiO2 (SiO2,96.74% Si, Cambridge 0.035, 0.07, 0.14, and 0.28 wt% SiO2 doped 1 at.% Nd:YAG. Isotobe Laboratories, Andover, MA). Samples were milled, dry The Nd:YAG powder for each composition was taken from a pressed, CIPed, and sintered at 16001C for 2 h according to the single powder batch, thus the stoichiometry is identical for each procedure described above. The Nd:YAG samples were probed composition. Powder batches were ball milled in anhydrous eth- using 29Si magic angle spinning (MAS) NMR. These experi- anol (reagent grade, J.T. Baker) for 16 h using high purity Al2O3 ments were carried out in a 7.0 T magnetic field with a Tecmag balls (99.9%, 5 mm diameter, Nikkato Corp., Sakai, Japan). Discovery NMR spectrometer (Tecmag Inc., Houston, TX) op- The volume ratio of powder:milling media:alcohol was 1:4:8, erating at a 29Si irradiation/detection frequency of 59.081 MHz. and the milling media had no measureable weight loss after The spectrometer was interfaced to a 7.5 mm double resonance milling. The milled suspensions were dried in an oven and NMR probe that was run with MAS rates of 3 kHz. A spectral ground in an alumina mortar. Samples denoted as ‘‘dry pressed’’ width of 40 kHz was covered with the acquisition of 2048 com- were uniaxially pressed in a 4 mm die at 20 MPa followed by plex data points. For each spectrum, a total of 4220 scans was 1382 Journal of the American Ceramic Society—Stevenson et al. Vol. 94, No. 5 acquired. The pulse length for irradiation of 29Si was 7 ms, providing approximately 901 of spin nutation, and the recycle delay between pulses was 60 s.

(5) TEM To observe grain-boundary structure and the spatial distribu- tions of Nd31 and Si in the microstructure, HRTEM was per- formed on tape cast Nd:YAG ceramics doped with 0.14 and 0.28 wt% SiO2 sintered at 17001 and 17501C for 8 h. A JEOL EM-2010F Field-Emission TEM/STEM (JEOL Ltd., Tokyo, Japan) equipped with an energy-dispersive X-ray spectroscopy (EDS) detector (Oxford Instruments, Abingdon, U.K.) and Ga- tan Enfina electron energy loss spectrometer (EELS) (Gatan Inc., Pleasanton, CA) was used for TEM characterization. The samples were polished to 0.5 mm surface roughness using dia- mond impregnated polymer films followed by argon ion milling on a liquid nitrogen cooling stage. HRTEM phase contrast im- ages were taken along the [111] direction for one of two adjacent grains with the other grain showing clear lattice fringes. Fig. 3. HRTEM image of a grain boundary in tape cast 0.28 wt% SiO2-doped Nd:YAG sintered at 17001C for 8 h. No amorphous phase is observed at the grain boundary.

III. Results and Discussion

(1) SiO2 in Nd:YAG Ceramics Figure 3 is an HRTEM micrograph of a grain boundary in In this section, we investigate the coordination and distribution 0.28 wt% SiO2-doped Nd:YAG sintered at 17001Cfor8h.The 41 of Si in r0.28 wt% SiO2-doped Nd:YAG ceramics between grain boundary is o1 nm thick with lattice fringes extending to 16001 and 17501C to inform our discussions about the effects of the boundary of each grain. Grain boundaries and triple lines SiO2 on sintering and microstructure evolution. show no signs of crystalline second phases or amorphous phases Figure 2 shows the NMR spectrum obtained from 0.28 wt% in this sample or any other sample in this study. EELS was also SiO2-doped Nd:YAG sintered at 16001C for 2 h. Two overlap- used to detect Si. The k ionization edge for Si at 1839 eV was ping peaks with full-width at half-maxima (FWHM) of 5 and detected at grain centers, grain boundaries, and triple lines in- 3ppmareobserved.AFWHMof3–5ppmisconsistentwith dicating that Si41 substitutes into the YAG lattice and is also crystalline order (amorphous phases frequently have present at grain boundaries and triple points. This observation 41 FWHM420 ppm), indicating that SiO2 is primarily coordi- provides further evidence that Si enters into a solid solution nated in highly crystalline environments. The 80 to 85 ppm with the Nd:YAG lattice during sintering. chemical shift is consistent with tetrahedrally coordinated 0.28 wt% SiO2 doping corresponds to 0.5% of the available 41 21 41 31 Si . Octahedrally coordinated Si results in chemical shifts Al sites in the YAG structure. Because SiO2 was added to between 180 and 220 ppm, but no peaks were observed in stoichiometric Nd:YAG powder, 0.28 wt% SiO2 doping leads to this range. It is therefore clear that the majority of Si41 occupies 0.5 at.% Al31 cation excess over the stoichiometric composition. the tetrahedrally coordinated Al31 24(d) site as opposed to oct- The precise limit of stoichiometric variation in YAG ceramics is ahedrally coordinated Al31 16(a) sites in the YAG structure. not known, but Patel et al.22 found that Al31 excess may be This result confirms predictions from atomistic modeling that limited to 0.2 at.% by comparing lattice constant measurements Si41 preferentially occupies tetrahedral sites in YAG. from X-ray data with atomistic simulations of defect mecha- The two distinct 29Si41 peaks observed in Fig. 2 may be ex- nisms. However, no evidence of second phases was found up to plained by cationic substitutions in the second coordination 0.62 at.% excess Al31 in that study. Further, the YAG ceramics 29 41 21 41 shell of the Si .Fyfeet al. observed similar chemical shifts in Patel and colleagues were doped with 0.14 wt% SiO2,butSi in tetrahedrally coordinated 29Si41 when ionic substitutions oc- content was not considered when determining sample stoic- curred in the second coordination shell surrounding the 29Si41 hiometry. After accounting for Si41 in the sample stoichiome- 31 31 tetrahedra. The second coordination shell of each Al 24(d)site try, the Al excess limit in SiO2-doped YAG ceramics is at least in YAG consists of eight Y31 24(c) sites and four octahedrally 0.45 at.%. Clearly, further work is required to determine the 31 coordinated Al 16(a) sites. In Nd:YAG, second coordination precise limits of nonstoichiometry in SiO2-doped YAG ceram- shell substitutions could occur by Nd31 occupying a 24(c)site, ics. In the present study, no second phases were observed by or by antisite defects (Y31 occupying 16(a) and 24(d)sitesor SEM or TEM in any sample, thus it is likely that Al31 excess Al31 occupying 24(c) sites) which are known to occur in YAG caused by Si41 doping does not result in second phase formation 20,22 ceramics and single crystals. up to at least 0.28 wt% SiO2. Given that the chemical shift and relatively narrow FWHM observed in Fig. 2, the fact that Si is detected at grain interiors by EELS, and that grain boundaries observed by HRTEM (Fig. 3) show no evidence of an amorphous phase, we conclude that Si41 primarily exists as solute ions located at tetrahedrally coordinated 24(d) sites in the YAG crystal structure for r0.28 wt% SiO2-doped Nd:YAG ceramics sintered between 16001 and 17501C.

(2) Densification Figure 4 shows relative density during final stage densification as a function of temperature and time for tape cast Nd:YAG ce- ramics doped with 0.035–0.28 wt% SiO2. Increasing silica Fig. 2. NMR spectrum of dry pressed Nd:YAG doped with 0.28 content increases density at all temperatures and times. The 29 wt% Si -enriched SiO2 sintered at 16001C for 2 h. 0.28 wt% SiO2-doped samples densify to over 99.9% relative May 2011 Effect of SiO2 on Nd:YAG Transparent Ceramics 1383

factor for both grain boundary and lattice diffusion in YAG ceramics.23,24 This implies that increasing Y31 and Nd31 diffu- sion rates should also lead to faster densification kinetics. Ionic diffusion in ceramics typically occurs by a vacancy mechanism and the self diffusion coefficient can be written25: Ej D ¼ D0 exp (2) kBT

where D is the diffusion constant, Ej is the activation energy to jump to an adjacent vacancy, kB is Boltzmann’s constant, T is 25 temperature, and D0 is :

2 D0 ¼ fa vXv (3)

where f is the correlation factor, a is the nearest neighbor jump distance, n is the jump frequency, and Xv is the site fraction of vacancies on the appropriate sublattice. According to Eq. (1), Xv 31 31 should increase on the Y /Nd sublattice with increasing SiO2 content. According to Eqs. (2) and (3), increasing Xv will in turn cause both D 3þ and D 3þ to increase with increasing SiO Fig. 4. (a) Relative density versus temperature for tape cast Nd:YAG Y Nd 2 content. Diffusion experiments on Nd2O3 coated YAG ceramics samples doped with 0.035, 0.07 0.14, and 0.28 wt% SiO2. Holding time 31 was 2 h. (b) Relative density versus sintering time at 16001C for tape cast confirm that Nd lattice and grain-boundary diffusion con- stants increase with increasing silica content.24 Thus, by creating Nd:YAG samples doped with 0.035, 0.07, 0.14, and 0.28 wt% SiO2. vacancies on the Y31/Nd31 sublattice (Eq. 1), and thereby in- 31 31 creasing diffusion rates for both Y and Nd ,SiO2 doping increases densification kinetics in Nd:YAG ceramics. density at short times (2 h for all temperatures 416001C) and Figure 5 shows representative microstructures of 0.035 wt% low temperatures (16001C for 8 h) while samples doped with less SiO2-doped Nd:YAG sintered for 2 h (Fig. 5a), 4 h (Fig. 5b), SiO2 achieve 499.9% density at longer times (8–20 h) and/or and 20 h (Figs. 5c and d) at 16001C. The isolated pores in the higher temperatures (17501C). Figure 4(b) shows that all sam- images are typical of final stage densification in SiO2-doped ples sinter to over 99.9% density after 20 h at 16001C. While Nd:YAG ceramics. Pores are similar in size to grains and are all samples had measured densities over 99.9% after sintering separated by large defect-free regions, i.e., regions where there at 16001C for 20 h, only the 0.28 wt% SiO2-doped Nd:YAG are no pores or second phases present at grain boundaries and samples were highly transparent (480% in-line transmission at triple points. In samples with densities 499.9%, these defect- 1064 nm) at this sintering condition. free regions (shown in Fig. 5(d)) were observed to be 4900 mm2. To explain the increasing densification rate as a result of in- In order to sinter to the exceptionally high densities 41 creasing SiO2 content, we examine the effect of Si substitution (499.995%) required for high efficiency laser materials, we on Y31 and Nd31 diffusion. Y31 and rare earth ion (Nd31, must consider final stage sintering characterized by a small num- Er31,Yb31, etc.) diffusion is reported to be the rate limiting ber of pores, similar in size to the average grain size, separated

Fig. 5. SEM micrographs of tape cast Nd:YAG ceramics doped with 0.035 wt% SiO2 sintered at 16001C for 2 h (a) and 8 h (b) and 20 h (c and d). 1384 Journal of the American Ceramic Society—Stevenson et al. Vol. 94, No. 5

Fig. 6. Sintering trajectory for tape cast and dry pressed 0.14 wt% Fig. 8. G3–G3 vs sintering time at 16001C for tape cast Nd:YAG 1 1 0 SiO2-doped Nd:YAG ceramics sintered between 1550 Cand1750Cfor ceramics containing 0.035, and 0.14wt% SiO . 2–8 h. 2 by multigrain distances, rather than small pores distributed content samples maintain grain sizes below 3 mm even at sinte- along grain boundaries and triple lines. ring conditions up to 17501C for 8 h. To determine the grain Because of large pores like those observed in Fig. 5 usually growth exponent (n) and rate constant (k), kinetic data was result from green forming defects, we investigated the role of plotted according to the following equation: green forming on microstructural development and densificat- ion. Dry pressed and tape cast samples doped with 0.14 wt% n n G G0 ¼ kt (4) SiO2 were sintered between 14841 and 17501Cfor1–8h.The average green densities of tape cast and dry pressed samples where G is the average grain size, G0 is taken to be the grain size were 48.3% and 44.7%, respectively, and the sample thicknesses at 2-h hold time for each temperature, and t is time. Grain after sintering were 2 mm for both tape cast and dry pressed growth data was fit to n values of 2–4. n 5 3 provided the best fit samples. for all temperatures and compositions (linear regression R2 pa- Figure 6 shows that tape cast samples have smaller average rameters ranged from 0.93 to 0.99). Boulesteix et al.11 and grain size at the same density as dry pressed samples. Because Kochawattana et al.24 also found that grain growth kinetics fit dry pressed samples have a small population of larger pores that 1 1 16 a cubic rate law for temperatures between 1525 and 1850 C. are not found in tape cast samples, dry pressed materials re- Figure 8 shows grain growth kinetics for tape cast 0.035 and quire longer sintering times and higher sintering temperatures to 0.14 wt% SiO2-doped Nd:YAG sintered at 16001Cfor2–60h achieve comparable densities to tape cast samples. The longer plotted according to Eq. (4) with n 5 3. sintering times and higher sintering temperatures required in dry As shown in Fig. 5, the vast majority of grain boundaries pressed samples lead to larger average grain sizes at identical do not intersect pores. Therefore, it is likely that grain growth sintered densities compared with tape cast samples. Clearly, the for samples sintered at 416001C is controlled by grain-bound- sintering trajectory is affected by the green forming process, and ary mobility rather than pore drag. In boundary-controlled as demonstrated in Fig. 6, colloidal methods like tape casting are systems, grain growth exponents equal to 3 are explained by better suited than dry pressing to achieving fine-grained trans- either diffusion-controlled growth in liquid phase containing parent ceramics by pressureless densification. This data coin- 27 26 systems or by solute drag mechanisms in solid-state systems. cides well with previous studies that found that colloidal As shown above, no liquid phase is observed in o0.28 wt% processing methods lead to increased particle packing homoge- SiO2-doped Nd:YAG ceramics sintered between 16001 neity and, therefore, decreased sintering temperatures than dry and 17501C, and thus we conclude that the cubic grain growth pressing. Also, because all samples in this study sintered to kinetics observed in Fig. 8 are due to solute drag. 499.9% density at 16001C and densification was only limited As noted above, Y31 and rare-earth ion diffusion limit mass by a small number of large pores, refining green forming pro- transport in YAG ceramics. Because Nd31 has both a larger cedures to increase green density and particle packing homoge- ionic radius and a greater atomic mass than Y31, it is expected neity would eliminate these large pores thus enabling sintering of to have a slower diffusion rate than Y31. Therefore, the ion in transparent Nd:YAG in low silica content (i.e., o0.14 wt%) Nd:YAG most likely to result in solute drag during grain ceramics at 16001C. growth is Nd31. During solute drag limited grain growth, the concentration (3) Grain Growth of the slow-diffusing solute atoms is higher at or near grain boundaries than the grain interiors. Therefore, if Nd31 causes Figure 7 shows the effects of temperature on grain growth for solute drag in the Nd:YAG system, we should find higher Nd31 0.035–0.28 wt% SiO2-doped Nd:YAG. Increasing silica content concentrations near grain boundaries. Figure 9 shows the increases grain size under all conditions. The 0.035 wt% silica Gibbsian interfacial excess (g)ofNd31 atgrainboundariesas determined by TEM EDS analysis of relative Nd31 content

Fig. 7. Grain size as a function of temperature between 14841Cand 17501C for tape cast Nd:YAG ceramics doped with 0.035, 0.07, 0.14, Fig. 9. Interfacial excess of Nd31 at grain boundaries in 0.14 wt% and 0.28 wt% SiO2. Sintering time was 2 h at each temperature. SiO2-doped Nd:YAG sintered at 17001Cfor8h. May 2011 Effect of SiO2 on Nd:YAG Transparent Ceramics 1385

Fig. 10. Grain growth rate constant, k,versusSiO2 content for tape Fig. 12. Sintering trajectory for tape cast Nd:YAG ceramics containing cast Nd:YAG ceramics doped with 0.035, 0.07, 0.14, and 0.28 wt% SiO2 0.035 and 0.14 wt% SiO2 sintered at 16001C for 8–60 h. sintered at 17001C for 2–8 h.

between grain boundaries and grain interiors where G is calcu- where o is molar volume, dgb is the grain-boundary thickness, 28 lated according to the following equation : Da is the diffusivity of the solvent ions, Nv is the solvent atom concentration, Q is the partition coefficient, and CN is the solute "# GB GI atom concentration in the grain center. As shown above, in- CNd CNd creasing SiO content increases both Y31 and Nd31 diffusivity G ¼ NY w (5) 2 CY CY in YAG which increases both Da and Db in Eq. (6). Because increasing Mb increases k as well, this analysis provides a plau- 31 31 sible explanation that is consistent with a solute drag hypothesis where NY is the number density of Y lattice sites, CNd is Nd 31 for the increasing grain growth rate as a function of SiO2 concentration, CY is Y concentration, w is the width of the analysis area (33 nm) and GB and GI refer to measurements content. 1 taken at the grain boundary and grain interior, respectively. The The activation energy for grain growth between 1600 and 1 average interfacial excess for 10 boundaries was 0.789 Nd31 1750 C was calculated using the Arrhenius equation: nm2. This correlates to an absolute concentration of 1.25 at.% 31 E Nd at grain boundaries which is 25% more than the as- ln k a C (7) 31 ð Þ¼ þ batched composition. Optical techniques have also shown Nd kBT segregation at grain boundaries. Ramirez et al.29 used confocal scanning optical microscopy (CSOM) and also showed that where k is the rate constant in Eq. (4), Ea is the activation en- Nd31 concentration increased at grain boundaries by 10% rel- ergy, and C is a constant. Figure 11 shows the Arrhenius plots ative to grain interiors in 0.14 wt% SiO2-doped Nd:YAG. The for each silica content. The calculated activation energies for lower spatial resolution of CSOM (o0.5 mm 0.5 mm grain growth of 0.035-, 0.07-, 0.14-, and 0.28 wt% SiO2-doped 2–3 mm) relative to TEM may account for the differences in samples were 960, 1210, 1130, and 1160 kJ/mol, respectively. grain-boundary concentration observed by the two techniques. These values are consistent with the activation energy for grain 41 1 1 A similar analysis was not performed for Si because the Y La growth in SSR ceramics sintered between 1700 and 1850 Cby peak overlaps with the SiKa peak. Kochawattana et al.11 Grain growth constants (k) were calculated according to Eq. (4) for 0.035, 0.07, 0.14, and 0.28 wt% SiO2-doped Nd:YAG (4) Microstructural Control and Sintering Trajectory 1 1 1 1 samples sintered at 1600 ,1650,1700, and 1750 Cfor2–8h. Along with green forming, understanding the links between Figures 10 and 11 show that k increases as a function of SiO2 sintering temperature, silica content, densification rate, and doping for samples sintered for 2–8 h at 17001C and all sintering 30 grain growth rate is key to controlling microstructure evolution temperatures used in this study. Cahn showed that boundary in transparent Nd:YAG. Figure 12 shows grain size as a func- mobility (Mb) in a solute drag limited system should vary with tion of density for tape cast 0.035 and 0.14 wt% SiO2-doped grain-boundary diffusivity of the solute atoms (Db) according to Nd:YAG sintered at 16001C for 8–60 h. Figures 4 and 7 show the following eqaution30: that increasing SiO2 content increases both densification and coarsening, and Fig. 12 shows that the sintering trajectory be- 1 O 1 4NVOQC1 gins to follow a more coarsening path in samples with increasing Mb ¼ þ (6) kBTdgb Da Db SiO2 content. While increasing SiO2 content leads to higher

Fig. 13. In-line transmission for tape cast 0.0.35 wt% SiO2 Nd:YAG Fig. 11. Arrhenius plot of ln k vs 1/T for tape cast 0.035, 0.07, 0.14 and sintered at 17501C for 8 h and 0.28 wt% SiO2 Nd:YAG sintered at 0.28 wt% SiO2-doped Nd:YAG sintered between 16001 and 17501Cfor 16001C for 8 h. Grain sizes were 2.8 and 18 mm and sample thicknesses 2–8 h. were 1.3 and 1.1 mm, respectively. 1386 Journal of the American Ceramic Society—Stevenson et al. Vol. 94, No. 5 densities at lower sintering temperatures and shorter times, com- leads to lower sintering times and temperatures as well as de- parable densities in samples with lower SiO2 content have creased grain size at the same density as dry pressed samples. As smaller average grain size. silica content increases, the sintering trajectory in Nd:YAG ce- Figure 13 shows the in-line transmission for tape cast ramics occurs along an increasingly coarsening path. Higher sil- Nd:YAG samples doped with 0.035 and 0.28 wt% SiO2.The ica content samples densify at lower temperatures and times polished samples were 1.3 and 1.1 mm thick, respectively. At than lower silica content samples. Lower silica content samples 0.28 wt% silica content, highly transparent ceramics with an may achieve full density by sintering at relatively high temper- 18 mm average grain size were achieved by sintering at 16001C atures, but the more densifying sintering trajectory leads to for 8 h. In the case of 0.035 wt% SiO2, highly transparent smaller average grain sizes than higher silica content samples ceramics with 2.8 mm average grain size were achieved after sintered at lower temperatures. sintering at 17501C for 8 h. These samples clearly illustrate the relationships between SiO2 content, sintering kinetics, and sinte- ring trajectory. 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