Formation and Photoluminescence Characterization of Transparent

Institute for Solid State Physics, University of Tokyo, Kashiwanoha 5-1-5, ... silica particles during the solid-phase sintering process using infrare...
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J. Phys. Chem. C 2007, 111, 12973-12979

12973

Formation and Photoluminescence Characterization of Transparent Silica Glass Prepared by Solid-Phase Reaction of Nanometer-Sized Silica Particles Tomoko Yamada Department of Chemistry, Graduate School of Science and Technology, Kobe UniVersity, Nada-ku, Kobe 657-8501, Japan

Makoto Nakajima and Tohru Suemoto Institute for Solid State Physics, UniVersity of Tokyo, Kashiwanoha 5-1-5, Kashiwa-shi, Chiba 277-8581, Japan

Takashi Uchino* Department of Chemistry, Faculty of Science, Kobe UniVersity, Nada-ku, Kobe 657-8501, Japan, and SORST, Japan Science and Technology Agency, Kawaguchi, Saitama 332-0012, Japan ReceiVed: March 23, 2007; In Final Form: June 28, 2007

We investigate the structural transformations of nanometer-sized silica particles during the solid-phase sintering process using infrared and Raman spectroscopy and field emission scanning electron microscopy (FESEM). We also carry out detailed time-resolved photoluminescence (PL) measurements for the fully sintered transparent sample, which has recently been shown to exhibit a unique white PL emission under ultraviolet excitation (Uchino, T.; Yamada, T. Appl. Phys. Lett. 2004, 85, 1164). We show that the macroscopic enclosed-pore formation and elimination processes are occurring during sintering and are closely correlated to the condensation reaction of surface silanol groups. The structure and bulk density of the fully sintered samples are analogous to those of normal bulk silica glass although the structure of the non-heat-treated sample is rather different. FESEM measurements reveal a particulate morphology even in the thoroughly sintered and apparently transparent sample. As for the fully sintered sample, we observe three distinguishable PL bands with different decay kinetics ranging from nanosecond to millisecond time regions. These versatile PL characteristics probably result from the interfacial highly constrained structures created during the present solid-phase sintering process.

1. Introduction Because of its excellent optical, electrical, and thermal properties, silica (SiO2) glass or amorphous silica has a wide variety of technological applications, including ultraviolet (UV) optics, optical fibers, and insulating oxide films in metal-oxidesemiconductor (MOS) devices. A traditional method to prepare bulk silica glass is the melt-quenching technique. That is, silica glass is prepared by melting quartz in vacuum or in an inert medium at low pressure. In addition, thermally activated homogeneous oxidation of SiCl4 is also used to prepare highpurity silica glass. The oxidation reaction of SiCl4, namely, SiCl4 + O2 f SiO2 + 2Cl2, produces a finely divided particulate glass material commonly called “soot”.1 These glass soots are then again melted to obtain high-purity bulk silica glass or to fabricate preforms of optical fibers.2 However, such a melt-based technique often presents serious processing difficulties because melting of silica requires very high temperatures in excess of ∼2000 °C. The sol-gel method has been an alternative technique for preparing optical quality silica glass; it allows us to obtain sintered bodies at relatively low temperatures in the range ∼800 to ∼1200 °C.3 Sol-gel processing methods usually utilize the hydrolysis reaction of silicon alkoxides and the subsequent condensation, resulting in an amorphous gel. The gel can be sintered to clear glass by appropriate heat treatment. However, * To whom correspondence should be addressed. E-mail: uchino@ kobe-u.ac.jp.

one of the drawbacks of these sol-gel techniques is that the resulting sintered silica bodies usually contain a large amount of OH groups (>∼1000 ppm).4,5 In the sol-gel methods, the final OH content is directly related to the initial water content of the solution from which it is prepared, and the residual OH species can be difficult to remove during heating because of the very small pores in the gels. To reduce the residual OH content, further heat treatment under vacuum or in an atmosphere containing He and/or Cl2 must be required.6 Nanometer-sized silica particles can also be used as the starting material for making bulk silica glass. One example of such nanometer-sized silica particles is fumed silica, which is produced in a hydrothermal process by burning SiCl4 in an oxygen-hydrogen flame at 1200-1600 °C. Fumed silica is produced by the oxidized reaction of SiCl4 similar to glass soot mentioned above; however, fumed silica consists of aggregates of very small particles in the rage of 10-20 nm,7 whereas a silica soot is a low-density porous glass body made up of the flocks of small silica particles in the range of 100-200 nm.1,8 Although fumed silica has been widely used as an active filler for reinforcement of elastomers and a rheological additive in fluids for more than 20 years, it has attracted renewed attention in terms of its interesting structural, chemical, and optical properties.9-16 In the fabrication of high-purity silica glass from fumed silica, a colloidal suspension consisting of silica particles and water is normally used, followed by gelation, drying, and sintering at ∼1500 °C.17 However, we have recently found that even “dry” powder compacts composed only of fumed silica

10.1021/jp072312v CCC: $37.00 © 2007 American Chemical Society Published on Web 08/11/2007

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particles are sintered to optically transparency at ∼1000 °C, forming dense silica glass.18,19 This sintering process is accompanied by significant shrinkage, showing coalescence of the original particles. It has also been demonstrated that the resulting transparent silica glass exhibits unique white-light emission under ultraviolet (UV) excitation.18 To get a better knowledge about the “low-temperature” solid-phase sintering process of fumed silica, we here investigate the structural changes of fumed silica during the sintering process using vibrational spectroscopies and field-emission scanning electron microscopy. We also carry out further time-resolved PL measurements in the time range from ∼1 ns to ∼1 ms to get a more detailed picture of the PL processes from the thus prepared samples. 2. Experimental Section In this work, the solid-phase reaction of fumed silica was carried out according to the procedure reported previously.18 We used a nonporous amorphous fumed silica obtained from Sigma [specific surface area 390 ( 40 m2/g; the particle size 7 nm (product specification)]. As-received powders of fumed silica (0.3 g) were pressed in a uniaxial press at 530 MPa, forming a translucent self-supporting pellet of 19.3 mm in diameter and 0.9 mm in thickness. The pellets were then heated at 980 °C for 24-192 h in air to induce sintering between the constituent particles.18 This sintering process induced shrinkage, resulting in the apparently transparent sample (the percent transmittance T ) ∼80% at 600 nm) of 15.5 mm in diameter and 0.7 mm in thickness. Fourier transform infrared (FTIR) absorption spectra of the heated samples were recorded with 10 scans and a resolution of 1 cm-1 on a Perkin-Elmer Spectrum 1000 FTIR spectrometer. To measure the frequency region from 2500 to 4000 cm-1, which gives information about the stretching vibrations of both surface ≡Si-OH and adsorbed H2O, the sintered pellets were directly set in the sample holder. On the other hand, we used a conventional KBr disk technique to measure the FTIR spectra of the samples in the frequency region from 400 to 1600 cm-1, which are attributed to the Si-O-Si stretching and bending motions of the silica network. We also measured Raman spectra using a Perkin-Elmer System 2000R NIR Fourier-transform Raman spectrometer. The morphology of the unsintered and sintered samples was characterized by field emission scanning electron microscopy (FESEM) with a JEOL JSM-6700F microscope operating at 1.5 kV. We estimated the bulk density of the sintered pellets using Archimedes’ method. Time-resolved PL measurements of the sintered samples were carried out by using two different pulsed laser sources. One is the fourth-harmonic of a nanosecond Nd:YAG laser (266 nm) operated at a repetition rate of 10 Hz. The pulse width was 8 ns, and the average power density was 30 mW/cm2, namely, 3 mJ/cm2 per pulse. Using the pulsed Nd:YAG laser and a gated image intensifier CCD camera, we measured the PL decay on the time scale longer than submicroseconds. To monitor a much faster decay, we used the third-harmonic of a femtosecond Ti: sapphire laser (267 nm), which was amplified by a regenerative amplifier at a repetition rate of 1 kHz, in combination with a synchroscan streak camera. The pulse width of the Ti:sapphire laser was 150 fs, and the average power was 3 mW/cm2, namely 3 µJ/cm2 per pulse. During the time-resolved PL measurements, the sample temperature was controlled in a closed-cycle N2 cryostat (77-400 K) or a He gas flow cryostat (5-300 K). During the above PL measurements, we observed neither any damages of the samples nor changes in the PL characteristics caused by laser irradiation.

Figure 1. FESEM images of (a) as-received and (b) sintered (980 °C, 168 h) fumed silica. The inset shows a photograph of each sample in the form of a pellet.

3. Results 3.1. FESEM Images. Figure 1a shows an FESEM image of the compacted pellet before heat treatment. We can recognize primary particles of fumed silica with a nominal size of ∼7 nm. They are packed closely together, resulting in a translucent appearance (see the inset of Figure 1a). Thus, the starting pellet is composed of highly compacted and uniformly distributed amorphous silica fine particles. We next turn to an FESEM image of the pellet after sintering at 980 °C for 168 h (see Figure 1b). The sintered sample is apparently transparent as shown in the inset of Figure 1b, but in the FESEM image, one still sees a number of boundaries on the microscopic length scale. It is hence quite likely that even thoroughly sintered samples still retain the original particle-like features and the related interface regions. This is probably because the sintering temperature employed here (980 °C) is not high enough to induce viscous flow nor to allow full coalescence of the particles. 3.2. Infrared Spectra. Figure 2 shows FTIR spectra of the fumed silica samples before and after sintering in the wavenumber region from 400 to 1500 cm-1 along with the corresponding FTIR spectrum of normal bulk silica glass. As has been reported previously,15 the FTIR spectrum of unsintered fumed silica has several features different from those of normal bulk silica glass. One of the apparent differences between the FTIR spectrum of unsintered fumed silica and that of normal bulk silica can be seen in the Si-O-Si asymmetric stretching band at ∼1100 cm-1. That is, the ∼1100-cm-1 band of the unsintered fumed silica has rather a narrow feature (the fullwidth at half-maximum fwhm ) ∼150 cm-1) as compared with the corresponding band of bulk silica glass (fwhm ) ∼200 cm-1). This implies that that the random network structure of fumed silica is not identical with that of the bulk one.15 The above spectral features peculiar to unsintered fumed silica tend to become less distinctive as sintering proceeds. As shown

Transparent Silica Glass

Figure 2. Fourier transform infrared (FTIR) absorption spectra of fumed silica (a) before and after heat treatment at 980 °C for (b) 24 h, (c) 96 h, and (d) 168 h. FTIR spectrum of normal bulk silica glass is also shown in (e). A conventional KBr disk technique was used to measure these FTIR spectra.

Figure 3. FTIR absorption spectra of fumed silica after heat treatment at 980 °C for (a) 48 h, (b) 72 h, (c) 96 h, (d) 120 h, and (e) 168 h. To measure these FTIR absorption spectra, the sample in the form of a pellet was directly set in the sample holder.

in Figure 2, the Si-O-Si asymmetric stretching band at ∼1100 cm-1 becomes broader with increasing sintering time. Accordingly, the FTIR spectrum of the sample sintered at 980 °C for 168 h becomes almost comparable to that of normal bulk silica glass. This suggests that the average structure of well-sintered fumed silica is very similar to that of the bulk one although the former sample still retains a particle-like morphology as mentioned in section 3.1. We also notice from Figure 2 that the FTIR spectrum of unsintered fumed silica has a weak shoulder at ∼980 cm-1, which is attributed to the Si-(OH) stretching mode of isolated silanol groups.20 This shoulder becomes less pronounced and appears to be merged into the broad Si-O-Si asymmetric stretching band at ∼1100 cm-1 with increasing sintering time. To examine the spectral changes associated with the OH groups in more detail, we then analyze the FTIR absorption spectra in the frequency region from 2600 to 4000 cm-1 (see Figure 3). We see from Figure 3 that the FTIR spectrum of the sample heated at 980 °C for 48 h, which is translucent and dose not show noticeable sintering-induced shrinkage, is characterized by two distinct spectral features: one is the sharp peak located at 3745 cm-1 attributed to isolated noninteracting surface silanol groups and the other is the very broad band peaking at ∼3400 cm-1 associated with hydrogen-bonded OH groups and adsorbed water molecules.21 In this spectrum, one also notices a weak shoulder near 3660 cm-1. According to previous infrared studies of fumed silica, the vibrational mode near 3660 cm-1 is attributed to the stretching motion of internal silanols,21-23 which are more or less inaccessible to most reactants such as atmospheric water molecules. With increasing heating time, the broad absorption band at ∼3400 cm-1 shows a dramatic decrease in intensity. Consequently, the shoulder near 3660 cm-1 tends to be highlighted

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Figure 4. Raman spectra of fumed silica (a) before and after heat treatment at 980 °C for (b) 24 h, (c) 96 h, and (d) 168 h. The Raman spectrum of normal bulk silica glass is also shown in (e).

and emerges as an isolated peak. The sharp peak at 3745 cm-1 then completely disappears when the samples are sintered for more than 120 h. It is interesting to note that the samples becomes apparently transparent after sintering for more than 120 h.18 Thus, we suggest that the samples are sintered to transparency when the condensation reaction of both hydrogenbonded and isolated silanol groups at the surface is completed, and only internal silanols remain to be removed in the transparent samples. It is known that such internal silanols also exist in bulk silica glass and give rise to the infrared absorption band at ∼3660 cm-1 as well.24,25 As for the 3660 cm-1 band in bulk silica glass, its molar absorptivity  is reported to be 86 L mol-1 cm-1.25 If we assume that the value of  is also applied to the present transparent sample, we can estimate the concentration of water (as H2O) from the observed values of the absorbance, thickness, and density of the sample. The concentration of water in the transparent sample prepared by sintering at 980 °C for 168 h is estimated to be ∼440 ppm. We also found that the water concentration is almost unchanged after prolonged sintering as long as the heat treatment is carried our in air atmosphere. 3.3. Raman Spectra. Changes in the FT-Raman spectra of the samples with sintering time are shown in Figure 4. Figure 4 also includes the FT-Raman spectrum of bulk silica glass for comparison. As reported in our previous paper,15 the unsintered sample has rather strong Raman peaks at 495 and 606 cm-1, which are assigned to the breathing modes of regular four- and three-membered silica rings, respectively.26-28 This indicates that such small-membered rings are more frequent in fumed silica than in the bulk, also showing that the network connectivity of fumed silica is different from that of the bulk one. In the FT-Raman spectrum of the unsintered sample, we also see a peak at ∼950 cm-1 associated with the Si-(OH) stretching mode, similar to the case of the FTIR measurements. We see from Figure 4 that the intensities of the sharp Raman bands at 495, 606, and 950 cm-1 observed in the FT-Raman spectrum of the unsintered sample tend to decrease as sintering proceeds. Finally, the FT-Raman spectrum of the sample sintered for 168 h becomes almost analogous to that of bulk silica glass. This indicates that the network structure of the fully sintered fumed silica sample is, on average, comparable to that of the bulk one, in accordance with the results of FTIR spectra. 3.4. Photoluminescence Measurements. Figure 5 shows a typical example of time-integrated and time-resolved PL measurements of the fumed silica pellet sintered at 980 °C for 168 h. From the time-integrated PL spectrum shown in Figure 5a, we see that the observed emission consists apparently of two broad overlapping bands peaking at ∼360 and ∼510 nm. The broad PL features indicate a broad distribution of the energy levels associated with the relevant emission centers. Such a broad distribution probably results from a structural inhomog-

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Figure 6. PL decays of fumed silica heated at 980 °C for 168 h at 360 nm measured at different temperatures. The fourth-harmonic (266 nm) of a pulsed Nd:YAG laser was used for excitation. The solid lines are least-square fits with a pure exponential function. For the decay data measured at 77 K, one sees a slight deviation from the fit at delay times longer than ∼550 µs. This deviation results from the effect of the slower PL component peaking at ∼510 nm.

nm by employing 3 µs and 20 ms for the gate delay and gate width time, respectively. On the other hand, if we employ a shorter gate delay and gate width of, for example, 0.2 and 0.2 µs, respectively, only a visible PL band peaking at ∼360 nm is observed (see Figure 5c). These results indicate that the PL bands at ∼360 and ∼510 nm are characterized by decay times of submicrosecond and several microseconds, respectively, at room temperature. In addition to these two bands, another PL emission band peaking at ∼410 nm was newly resolved when we employ a gate delay and gate width of 0 and 0.1 µs, respectively (see Figure 5d). This demonstrates that a much faster PL component on a time scale of nanoseconds is also involved in the whole PL emission process. As for the ∼510 nm band, we have already reported that the decay has a fast component of a single-exponential form, followed by a stretched exponential decay.19 This can be represented by the following equation:

I(t) ) A exp(-t/τf) + B exp[-(t/τs)β]

Figure 5. (a) Time-integrated and (b)-(d) time-resolved photoluminescence (PL) spectra of fumed silica heated at 980 °C for 168 h under excitation of the fourth-harmonic (266 nm) of a pulsed Nd:YAG laser. The gate delay and gate width used to measure the respective timeresolved PL spectra are shown. The sharp peak in (d) is the second harmonic (532 nm) of a pulsed Nd:YAG laser.

eniety of the emission centers embedded in the disordered silica network. Figure 5a also shows the time-integrated PL spectrum of the unpressed fumed silica powders that are heat treated at 980 °C for 168 h; however, this powdered sample did not show any appreciable PL emission irrespective of the same heat treatment employed for the pellet samples. This indicates that the PL from the pellet samples results not from any impurities or carbon-related contaminations in fumed silica but from some intrinsic defects that are created during the solid-phase sintering process peculiar to the pellet samples. Figure 5b-d shows the results of time-resolved measurements on the fumed silica pellet sintered at 980 °C for 168 h. As shown in Figure 5b, we observe an isolated PL band peaking at ∼510

(1)

Here τ is an effective decay time, β is a dispersion factor between 0 and 1, and subscripts f and s are the values for the fast and slow components, respectively. Since the decay dynamics of the ∼510 nm band was already discussed in refs 18 and 19, we will mainly investigate the time decay of the newly resolved peaks at ∼360 and ∼410 nm. The time decay of the PL band at ∼360 nm is shown in Figure 6. In contrast to the case of the ∼510 nm band, the decay profiles of the ∼360 nm band are well fitted by a pure exponential function, namely,

I(t) ) A exp(-t/τ)

(2)

Figure 7 shows the temperature dependence of the decay time of the ∼360 nm PL band. We see from Figure 7 that τ(T) is described by the following Mott-type equation:29

τ(T) )

τR 1 + τRνNR exp(-(Ea/kT))

(3)

Here τR is the radiative lifetime and νNR is the transition rate of a competing nonradiative recombination process with an activation energy Ea. The fitted values of τR, Ea, and νNR are 225 µs, 0.033 eV, and 7.2 × 106 s-1, respectively. It should be noted that the τR value of the ∼360 nm PL band is even longer than those obtained for the slow and fast components (∼60-70 µs)

Transparent Silica Glass

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Figure 7. Temperature dependence of the decay time of the PL at 360 nm. The decay time at each temperature was obtained by fitting the corresponding data shown in Figure 6 with eq 2. The solid line is a least-squares fit with eq 3.

Figure 9. Changes in the bulk density (solid circles) and the absorption intensity of the infrared band at 3745 cm-1 shown in Figure 3 (empty boxes) as a function of heating time. The lines are guide to the eye.

4. Discussion

Figure 8. PL decay profiles of fumed silica heated at 980 °C for 168 h at 410 nm measured at 5 K (empty circles) and 250 K (solid triangles). The third-harmonic (267 nm) of a pulsed Ti:sapphire laser was used for excitation. The solid line is a least-squares fit with eq 4 for the data obtained at 5 K.

of the ∼510 nm PL band reported previously.19 We have previously reported that the temperature dependence of the decay time of the slow and fast components of the of the ∼510 nm PL band is also described by the Mott-type equation; the obtained activation energies of the nonradiative processes for the fast and slow components are 0.061 and 0.28 eV, respectively. Thus, the apparent submicrosecond decay time of the ∼360 nm PL band near and above room temperature can be ascribed to a rather small activation energy of the competing nonradiative recombination process. We should also note that the fitted value of νNR (7.2 × 106 s-1) is substantially smaller than a typical thermal attempt frequency of ∼1013 s-1. Although its physical origin is unknown, the reduced attempt frequency has also been found in the system in which phonon-assisted tunneling is likely to occur.30 Further work needs to be done to clarify the origin of the anomalous decrease in the attempt frequency. Next, we move on to the decay dynamics of the ∼410 nm PL band. The decay profiles detected at temperatures of 5 and 250 K are shown in Figure 8. As shown in Figure 8, the temporal decay of this emission was found to be temperature independent within experimental error. We also found that the decay curve can be well described by a single stretched exponential function,

I(t) ) A exp[-(t/τ)β]

(4)

The least-square fitting gave almost the same values of τ (0.03 ( 0.01 ns) and β (0.3 ( 0.02) for all the decay curves detected at temperatures from 5 to 300 K.

4.1. Microscopic and Macroscopic Mechanisms of SolidPhase Sintering. Changes in the OH absorption bands with heating time shown in Figure 3 suggest that the coalescence of the nanometer-sized silica particles, from a microscopic point of view, is closely related to the dehydroxylation reactions of the surface silanol groups. From a macroscopic view point, on the other hand, the pore elimination process will play a vital role in the entire sintering process. Since the apparent density of the bulk sample, i.e., bulk density, is influenced by the formation and elimination of pores within the sample, it is interesting to investigate the changes in the bulk density with heating time (see Figure 9) and to discuss how the bulk density is related to the dehydroxylation process that is occurring on the microscopic length scale. From Figure 9, we see that the bulk density of the non-heattreated pellet is 2.20 g/cm3, which is identical with the true density of the fumed silica particle and normal silica glass. This indicates that the pellet processing does not create enclosed pores in the sample. The bulk density then shows a gradual decrease with increasing heating time up to 72 h at 980 °C, illustrating that the open pores tend to be enclosed as sintering proceeds. From 72 to 96 h, however, the bulk density of the sample is almost unchanged. This suggest that the additional formation of the enclosed pores does not occur during this period of heating time. It is worth mentioning that when the heating time amounts to 120 h or over, the bulk density of the sample shows a substantial increase and reaches the value of the starting nonheat-treated pellet (∼2.20 g/cm3). This result demonstrates that the shrinkage and elimination of the enclosed pores occurs by heating the sample for more than 120 h. This phenomenon is consistent with the fact that the samples are sintered to transparency after heating the sample for 120 h.18 In Figure 9, we also show the changes in the peak intensity of the infrared band at 3745 cm-1, which is attributed to the stretching vibration of noninteracting surface silanol groups (see section 3.2). We see from Figure 9 that there is a correlation between the bulk density and the intensity of the infrared band at 3745 cm-1. That is, the change in the bulk density up to ∼98 h is accompanied by the concomitant change in the infrared absorption at 3745 cm-1. This implies that the formation of enclosed pores results from the interparticle dehydroxylation reaction of the surface silanol groups, namely, 2 ≡Si-OH f ≡Si-O-Si≡ (see Figure 10a,b). When the sample becomes transparent by heating the sample for more than 120 h, the infrared band related to the surface silanol groups almost disappears. This result strongly suggests that the elimination of the enclosed pores that leads to transparency is accomplished by the completion of the dehydroxylation reaction of surface

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Figure 10. Schematic atomic arrangements of (a) a non-heat-treated sample with open pores, (b) a heat-treated sample with macroscopic enclosed pores, and (c) a heat-treated sample with no macroscopic enclosed pores. Broken lines represent hydrogen bonds. The Si atoms involved in and derived from the edge-sharing tetrahedral units are circled in (b) and (c), respectively. The shaded part in (c) indicates the region in which the interfacial atomic arrangements are not fully relaxed and are hence highly deformed.

silanol groups. It can hence safely be said that the formation and elimination processes of the enclosed pores on the macroscopic length scale are closely related to the microscopic dehydroxylation reactions of surface silanol groups. Bunker et al.31 have previously reported that edge-shared surface defects are created when fumed silica powders are dehydroxylated at 900 °C according to the following reaction:

The resulting edge-sharing tetrahedral units have a relatively large strained energy (∼1.8 eV32,33) and are destroyed instantaneously via dissociative chemisorption reactions with, for example, ammonia and methanol.31,34 Thus, it is quite likely that these edge-sharing tetrahedral units are formed in the present dehydroxylation process of fumed silica as well and that the interparticle siloxane bonds are then created by way of the subsequent dissociative reactions involved in the interfacial edge-sharing units, as shown schematically in Figure 10b,c. Thus, we propose that the observed low-temperature sintering behavior of fumed silica results not only from the large inherent surface energy of small silica particles6,7 but also from the highly reactive nature of the edge-sharing tetrahedral units that are expected to be created during dehydroxylation reaction of fumed silica. 4.2. Photoluminescence Characteristics. We have shown that there exist at least three different PL mechanisms in terms not only of the PL wavelengths but also of the decay kinetics. Table 1 summarizes the PL characteristics of the three PL bands observed from the sample prepared by sintering at 980 °C for 168 h. The PL band at ∼360 nm shows an almost pure exponential decay. This implies that the energy states associated with the emission are interpreted in terms of simple two-level systems. It is hence quite likely that the energy states related to the ∼360nm PL band are highly localized at the germane emission center. That is, the ∼360-nm PL band probably results from a defective oxide structure whose ground and excited electronic states are highly localized.

TABLE 1: PL Characteristics of the PL Bands Observed from the Consolidated Fumed Silica Sample Prepared by Sintering at 980 °C for 168 h PL peaks (nm) decay time region at room temp decay profile temp dependence of the decay profile

360

410

510

0.1-1 µs eq 2 yes

1-1000 µs eq 1 yes

1-10 ns eq 4 no

On the other hand, the PL bands at ∼410 and ∼510 nm exhibit nonexponential decays described by eqs 1 and 4, respectively. We have already discussed the decay kinetics of the ∼510-nm PL band in detail in our previous paper.19 From the temperature dependence of the decay profiles, we have proposed that the slow PL component of eq 1 results from trapping-controlled diffusion of photoexcited electrons from localized to extended states, whereas the fast PL component is ascribed to the radiative process of the photoexcited electrons before experiencing the diffusion. The ∼410 nm PL exhibits a stretched exponential decay, as in the case of the slow PL component of the ∼510 nm band, suggesting that the dispersive transport of the relevant photoexcited electrons is responsible for the PL phenomena. We should note, however, that the decay profiles shown in Figure 8 do not vary with temperature, implying that the dispersive diffusion does not result from “trapping”, which involves thermal activation from the site to conduction states, but from “hopping”, which is governed by tunneling directly between localized sites. Provided that the diffusion is based on the hopping mechanism, theoretical considerations predict that the parameter β in eq 4 is temperature independent and reflects the spatial distribution of localized states.35,36 When the diffusion event occurs in a fractal with dimension d˜ , it has been shown that the relaxation function φ(t) is given by36

φ(t) ) exp(-ptd˜ /2)

(6)

where p is the parameter related to the density of the hopping sites. If we assume that eq 6 can be applied to the decay process of the ∼410-nm PL band shown in Figure 8, we can estimate

Transparent Silica Glass a fractal dimension d˜ = 0.6 from the fitted value of β (∼0.3). Since the fractal dimension obtained is less than 1, it is likely that the hopping sites are highly scattered throughout the system. The apparently textured morphology related to the interface shown in Figure 1b may be responsible for the scattered distribution of the hopping sites expected from the estimated fractal dimension derived from eq 6. Unfortunately, we have not yet established the microscopic structural models of these PL emissions. As shown in Figure 1b, the FESEM image of the transparent sample retains some characteristics derived from the original nanometer-sized particles. This allows us to suggest that the interparticle bondings created by the dissociative reaction of edge-sharing units proposed in section 4.1 are not fully relaxed and will be characterized partly by deformed structures because of the lowtemperature sintering. In other words, the interparticle siloxane bondings created by the interparticle reaction will have highly strained configurations (see the shaded region in Figure 10c). We propose that these highly strained atomic configurations at the interface will create some midgap states and are responsible for the unique PL characteristics peculiar to the present transparent samples. It should also be noted that as shown in Figure 5a, the PL emissions cannot be found from the powdered samples but from the pellet even when the same heating process is applied to these two types of samples. Furthermore, the PL emission from the sintered pellet is very weak when the sintering time is too short (∼24 h) to induce shrinkage.18 These results strongly suggest that the interparticle reactions and the related interfacial atomic rearrangements play a major role in creating the germane emission centers. However, additional experiments and theoretical considerations using, for example, the density functional theory37 will be required to get more information about the structural origins of the emission centers. 5. Summary We have investigated the structural transformation of compacted fumed silica pellets during the sold-phase sintering at 980 °C that eventually leads to the apparently transparent silica glass. We have also carried out detailed time-resolved PL measurements to characterized the entire PL phenomena observed from the fully sintered samples. We have shown that the solid-phase reaction proceeds in concomitant with the dehydroxylation reaction of surface silanol groups. This hydroxylation will partly result in the formation of highly reactive edge-sharing tetrahedral units, followed by the formation of the interparticle siloxane bonds. Formation and elimination of the enclosed pores are also accompanied by the dehydroxylation reaction of surface silanol groups. Thus, there is a good correlation between the macroscopic and microscopic structural evolutions during the present solid-phase sintering process. The infrared and FT-Raman spectra of the fully consolidated samples are very similar to those of bulk silica glass, indicating that the structure of the former resembles that of the latter on average. However, the FESEM image of the fully sintered transparent sample still retains a particulate morphology. We suggest that some of the atomic arrangements at the interface of the particles are characterized by highly deformed structures and are

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