Optical Materials 33 (2010) 14–18

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Optical Materials

journal homepage: www.elsevier .com/locate /optmat

An investigation on the dependence of photoluminescence in Bi2O3-dopedGeO2 glasses on controlled atmospheres during melting

Xin Jiang, Animesh Jha *

The Institute for Materials Research, Houldsworth Building, University of Leeds, LS2 9JT, United Kingdom

a r t i c l e i n f o a b s t r a c t

Article history:Received 30 December 2009Received in revised form 9 July 2010Accepted 13 July 2010Available online 16 August 2010

Keywords:Bi2O3-doped germanate glassesPhotoluminescenceOxidizing atmosphere

0925-3467/$ - see front matter Crown Copyright � 2doi:10.1016/j.optmat.2010.07.011

* Corresponding author.E-mail address: a.jha@leeds.ac.uk (A. Jha).

We have investigated the influence of a range of melting atmospheres on the control of the photo-activeBi-ion states in a Bi2O3-doped germanate glass, having a molar composition of 70GeO2–18Li2O–11Ga2O3–1Bi2O3. The atmosphere during melting was controlled by purging with pure oxygen, air, and nitrogen.The melting temperatures investigated were in the range of 1200 to 1450 �C. The UV–visible absorbanceand the photoluminescence (PL) spectroscopic characterizations in the visible (600–800 nm) and near-IR(1000–1500 nm) regions were carried out for determining the contributions of potential Bi-ion states onthe absorbance and PL. The PL properties of undoped GeO2 glass were also compared for the analysis ofpotential defects states present in the structure.

Crown Copyright � 2010 Published by Elsevier B.V. All rights reserved.

1. Introduction

The fluorescence in Bi3+-doped borate, phosphate, and germa-nate glasses were first reported at and below room temperaturein the visible range (300–575 nm) by Webb and Parke in 1973,followed by Reisfeld and Boehm, and Reisfeld, Boehm and Barnett[1–3]. However, it was not until 2001 when the photolumines-cence (PL) in the longer wavelength parts of visible (600–800 nm) and near IR (1000–1500 nm) spectra were first reportedby Fujimoto and Nakatsuka, who also demonstrated near-IR broad-band PL and signal amplification at 1300 with 800 nm excitation[4,5]. Since then there have been intense investigations in this areaby several groups using different hosts and processing techniquesin bulk glass, glass–ceramic and fibre forms [6–9]. The research hasalso demonstrated modest to high slope efficiency laser devicesand broadband ASE in fibres and waveguides, using semiconductordiode laser pumping schemes at 800 and 980 nm wavelengths, andNd-doped fibre lasers at 1064 nm [9–11]. The strong emissionband in near-IR has also allowed the generation of frequency-dou-bled lasers in the yellow region, demonstrated by Dianov and co-workers [9,11]. The PL in the 1000–1500 nm region overlaps withthe zero-dispersion wavelength of silica fibre at 1317 nm, at whichcurrently there are no efficient commercial optical fibre amplifiers,which is why there is much interest in the bismuth ion doped fi-bres for amplifier design. It is anticipated that a Bi-doped fibreamplifier device will surpass the quantum efficiency (QE) of 5%in the Pr3+-doped commercial fibre amplifiers, developed in the

010 Published by Elsevier B.V. All

early 1990s by NTT in Japan and British Telecom in UK. In the lab-oratory other types of Pr3+, Pr3+–Yb3+, Nd3+ and Dy3+ doped fibreswere also tested successfully for pumping at 1040, 980, and800 nm, respectively [12–16], but no commercial devices wererealized. Although the ultra-high speed network for long-haulcommunication is growing as a result of the technological ad-vances in the erbium-doped fibre (EDF) and Raman amplifiers,availability of an amplifier with comparable performance in the1250–1500 nm range with 980 nm pumping scheme will be ableto reduce the cost of future systems significantly [17]. The mainfocus for developing the Bi-doped fibre laser and amplifier is todesign an amplification system which can be pumped with asuitable 980 nm diode laser source. It is for this reason, there hasalso been studies on co-doping of Bi-glasses with Yb3+-ions [18].

In order to analyze the role of Bi-ions as active colour centres ina glass for amplifying signals, it is fundamental to understand thedependence and control of the valence states of ions under meltingconditions. Such an analysis is necessary since any variation in thefabrication parameters, especially the temperature and atmo-sphere, appears to contribute to a dramatic variation in PL, whichis apparently known to be linked with melting condition and com-position of the glass hosts [19]. In an attempt to quantify the ef-fects of processing on the valence states of Bi-ions in the glass,we have chosen a germanium oxide glass, which can be meltedat much lower temperatures than the silicates. The lower meltingtemperature of GeO2 glasses than that in silicate will exponentiallyreduce the thermodynamic driving force for chemical decomposi-tion of Bi2O3 to lower valence state oxides, as shown below. Suchan approach might allow us to study the contributions from notonly the Bi2O3 but also the lower valence states on the PL in these

rights reserved.

X. Jiang, A. Jha / Optical Materials 33 (2010) 14–18 15

glasses. Another important reason for choosing the GeO2 glassesfor Bi-ion doping is the multiplicity of (Ge4+) cations in tetrahedronand octahedron sites in germanium oxide glasses [20], which is un-likely to happen in silicate and phosphate hosts. Since the Bi3+-ionshave a large ionic radius (0.098 nm), when doped in silicate glassesat smaller concentrations due to the size-limiting geometrical con-strain, the ion is energetically more favourable to occupy limitednon-bridging sites which provide energetically favourable environ-ments for oxygen co-ordination. By comparison in a GeO2 glass,due to the presence of both the tetrahedral and octahedral sites[20], the Bi-ions may be able to distribute around the non-bridgingoxygen sites near the 4-fold and 6-fold co-ordinated germaniumsites. Besides the structural contribution to the overall emissionand absorption spectra for Bi-ions, the radiative transition proba-bilities of Bi-ions in a metastable state also increases with thesquare of refractive index of a dielectric medium, from which weanticipate that the radiative rate for Bi-ions in a GeO2-rich environ-ment will increase in comparison with when Bi-ions are present insilicates [21]. Since the Bi3+ ions are not known to have any PL atroom temperature between 800 and 1500 nm, in this paper we at-tempt to characterize the probable lower valence states of Bi-ions,which are more likely to form at higher temperatures and loweroxygen partial pressures, due to favourable thermodynamic equi-librium conditions in reactions:

Bi2O3 ¼ Bi2OðgÞ þ O2 ð1aÞ

and

Bi2O3 ¼ 2BiOðgÞ þ 0:5O2ðgÞ: ð1bÞ

The thermodynamic conditions for the decomposition reactionsin (1a) and (1b) are also an area for discussion for relating the 1+and 2+ valence states, responsible for near-IR emission in Bi-iondoped glasses. From reactions (1a) and (1b), it is apparent that asthe oxygen partial pressure at the reaction interface is reduced,which may be in equilibrium with the ambient gas (e.g. nitrogen,air, and oxygen), the equilibrium points in reactions (1a) and(1b) will shift in accordance with the Le Chatelier’s principles inthe forward direction. In the corresponding Eqs. (2a) and (2b) be-low, the values of K1a and K1b are constant at a given temperaturebecause the values of DGo depend linearly on temperature within aphase boundary condition. The DGo values can be found out fromthe thermodynamic database and software [22,23].

K1a ¼ exp �DGo1a

RT

� �¼

pBi2O � pO2

pBi2O3

ð2aÞ

K1b ¼ exp �DGo1b

RT

� �¼

pBiO � ðpO2Þ0:5

pBi2O3

: ð2bÞ

Below in Section 3, we have described the complex equilibriumconditions for trivalent, divalent and monovalent bismuth states.Since the equilibrium constant, Ki is constant at a given tempera-ture, the chemical potential, li ¼ RT‘nðpiÞ ¼ RT‘nðciXiÞ, of dissolvedbismuth states in a glass is then proportional to their molar concen-trations from the dilute solution theory, since the value of ci may beassumed to be constant in the dilute solution range of 1 wt.% ormol% concentration. In view of the assumption on dilute solution,it is apparent from Eqs. (2a) and (2b) that the values of partial pres-sures of Bi2O and BiO are inversely proportional to the p(O2) and(pO2)0.5, respectively. From this condition the concentrations ofthe Bi+ and Bi2+ states are directly dependent on the values of theequilibrium partial pressures of Bi2O and BiO in the melting gaseousatmosphere. In other words, we can infer that as the melting tem-perature increases and as more oxygen is lost (i.e. lower values ofp(O2)) from the glass-forming liquid, consequently the partial pres-sures of lower valence states of bismuth (in oxide form) will in-

crease, in accordance with Eqs. (2a) and (2b) at a constanttemperature. The resulting glass-forming liquid is, therefore, ex-pected to have higher concentrations of the lower valence statesthan the starting 3+ states of bismuth ions. By assuming the aboveequilibrium conditions, we have chosen three different atmo-spheres of gaseous phase in the present investigation: pure oxygen,air and nitrogen for melting and homogenization of glass-formingliquids. The temperature range for melting was between 1275 and1500 �C. The samples obtained were then analyzed for the absorp-tion and photoluminescence properties of germanium oxide glassesdoped with Bi2O3 oxide.

The influence of melting conditions on the resulting colour ofglass-forming liquid was reported for a lead oxide –aluminiumoxide containing GeO2 glass by Hughes et al. [24], however it isnot clear from this investigation whether the colour changes inglass samples were due to PbO, or Bi2O3, or both. This is evidentfrom the observations made by Sharonov et al. [25], who have ar-gued a point defect based colour centre model in such glassesincluding those with Te, Sb, In and Sn. No specific description ofpoint defects, however, was given in either of the two referencesin [24,25]. Besides the dependence of melting and constituent ele-ments with 5p (Sn, Sb) and 6p (Bi, Pb) electrons, colour centres arealso known to be affected by the composition of the glass hosts,especially with phosphate and alumina [2,3,6,11,26]. A contrastingmodel based on Bi5+ states, proposed by Fujimoto and Nakatsuka[4,5], is unlikely to arise in a oxygen deficient atmosphere of airand nitrogen used in our investigation.

Our aim in this paper is to characterize the dependence of Bi-ionrelated colour centres as functions of temperature and oxygen par-tial pressure in the context of equilibrium conditions, describedabove in Eqs. (1a) and (1b). We have specially chosen GeO2 basedglasses without the presence of any other 5p or 6p electronic struc-ture elements, except Bi. The reason for choosing GeO2 glasses isalso apparent from their ease of preform fabrication [27,28] usingrod and tube technique [29]. The data for the proposed colour cen-tre are also compared with the results obtained from Denker et al.[19,30] which is helpful for developing a much more generalizedmodel for Bi-ion colour centres.

2. Experimental

The molar percentage concentrations of bismuth-doped(1 mol% Bi2O3), germanium oxide (GeO2), lithium oxide (Li2O), gal-lium oxide (Ga2O3) were 70, 18, 11 mol%, respectively. Theseglasses are designated with an acronym GLGB, and were meltedin air (P(O2) = 0.209 atm), oxygen (P(O2) = 1 atm), nitrogen gasatmospheres. The residual concentrations of CO2, H2O and O2 incommercial nitrogen gas cylinders approximate to P(O2) � 10�5

atm. For each melting atmosphere and temperature, the derivedglass samples are compared in Fig. 1. In order to compare the ef-fects of atmosphere, we also analyzed the photoluminescence inundoped Bi2O3 glasses, which were also melted in the temperaturerange of 1275–1475 �C in air, as shown in Fig. 1, and these we des-ignate as GLG glasses. Note that the starting materials used forglass melting were analytical grades of chemical reagents with99.999% certified purity. Each melt contained a 10 g batch of chem-icals in a platinum crucible, which was melted inside a tube fur-nace, purged with 1 l per minute of a chosen gas (air or oxygenor nitrogen). The overall melting and homogenization time at anisotherm was 3 h after which the liquid was cast into a preheatedmould at 360 �C and then annealed at 380 �C for 3 h before it wascooled down to room temperature at a rate of 0.5 �C min�1.

Once the cast glasses were annealed, each glass was cut and finepolished for spectroscopic analysis using a Perkin Elmer Lambda 19UV–visible spectrophotometer in the range of 250–850 nm. The PL

Fig. 1. A pictorial comparison of germanate glasses melted under oxygen (series 1),air (series 2), and nitrogen (series 3) atmospheres from 1275 to 1500 �C. UndopedBi-doped glasses fall into series 4 compositions.

300 400 500 600 700 800Wavelength [nm]

0.0

0.2

0.4

0.6

0.8

Tran

smis

sion

[%]

Fig. 2. A comparison of the UV/visible transmission spectra of the GLGB glassesmelted in air. Sample thickness = 3 mm.

16 X. Jiang, A. Jha / Optical Materials 33 (2010) 14–18

spectra of glass samples, melted in air, oxygen, and nitrogen, weremeasured with an Edinburgh Instrument PL spectrometer at roomtemperature by using the visible and near infrared (NIR) photo-multiplier tube (PMT) detectors. The pump source for excitationwas an Argon ion laser line at 514 nm. The filters were used tominimize the scattering of stray light inside the spectrometer,which especially may affect the analysis of emission spectrum.

300 400 500 600 700

0.0

0.2

0.4

0.6

0.8

Tran

smis

sion

[%]

Wavelength [nm]

GLG-Air GLGB-O2

GLGB-Air GLGB-N2

Fig. 3. A comparison of UV/visible transmission spectra of the GLGB and GLGglasses melted at 1350 �C. Sample thickness = 3 mm.

3. Results and discussion

From Fig. 1, it is clear that with increasing melting temperature,the glass colour changes from transparent to red brown and finallyto an opaque dark. We observe that the lack of oxidizing atmo-sphere is also critical in enhancing the colour in bismuth-dopedglasses. For undoped GeO2 (GLG) glasses, the temperature-depen-dent of colorization phenomenon is also apparent in Fig. 1 and, itis unrelated with the presence of Bi-ions. The intensity of absorp-tion due to colour centres in GLG glasses, however, is much weakerthan that observed for the GLGB family of glasses, melted underdifferent atmospheric conditions. The decrease in the partial pres-sure of oxygen enhances the colour, as it is apparent from the melt-ing trials at various isotherms in air, oxygen and nitrogen. Withincreasing temperature, under a given atmospheric condition, theabsorption generally increases, which suggests that the strongestcolours are due to higher concentrations of lower valence statesof bismuth ions, in accordance with the equilibrium conditions de-scribed in Eqs. (1a) and (1b).

Fig. 2 shows the UV/visible transmission spectra of the glasses,melted in air at various temperatures between 1300 and 1400 �C.In general, the spectra show three absorption bands centred at370, 460, and 500 nm. The UV cut-off edge shifted to longer wave-lengths and the overall transmission reduced for all glasses meltedat higher temperatures. Fujimoto and Nakatsuka [4,5] reportedthree absorption peaks at 500, 700, and 800 nm in bismuth-dopedsilica glasses, which by comparison differs with germanate glasses,in which the three peaks at 370, 500, and 700 nm and were re-ported by Peng et al. [6] and Meng et al. [26] and Ren et al. [31,32].

In Fig. 3 we compare the UV/visible transmission spectra ofthree GLGB and a GLG samples, melted at 1350 �C in the atmo-spheres of air, oxygen and nitrogen gases, as shown in Fig. 1. In thisfigure, the UV-cut-off edge of a GLG glass melted in air is at around268 nm, which is blue-shifted with respect to the UV cut-off edgesin Bi-doped glasses, melted in air, oxygen, and nitrogen. Althoughin GLG glasses there were no Bi-ions present, the melting in air at1350 �C exhibits an absorption band just below 330 nm on the ris-ing edge of the bandgap absorption, which suggests that there maybe sub-gap absorption centres present in this glass, and thesemight be due to the Ge–O related defect centres [33]. In addition

we observe that the slopes of bandgap absorption curves in GLGBreduce dramatically when comparing oxygen melted glass with anitrogen melted, indicating that the shape and the population ofthe density of state curves are strongly dependent on the meltingcondition. In Fig. 3, the apparent differences in the slopes of theshort wavelength absorption edge, absence of absorption peaksabove 330 nm, and the blue shift of the absorption edge in GLGglass at 268 nm with respect to the edge at 323 nm in GLGB glassesmay point out that the GLG glasses may not have been affected bythe presence of Bi-ionic states. When the oxygen partial pressure is1 atmosphere in the melting chamber, as shown in Fig. 3, theabsorption edge at 323 nm rises sharply in GLGB glass, almost ina similar manner as in GLG glass melted in air.

In Fig. 3 for all GLGB glasses melted at 1350 �C, the meltingatmosphere also affects the strength of absorption bands at460 nm and 500 nm, which lie below the electronic gap of the cho-sen germanium oxide glass composition. Similar comparabletrends were also observed in Bi-ion doped GLGB glasses at othermelting temperatures. It should be noted, by comparing the Figs. 2and 3, that the absorption band at 370 nm is more sensitive tomelting temperature than the oxygen present in the atmosphere.For example, the strength of absorption is more dominant at higher

6.0 6.1 6.2 6.3 6.4 6.5

1

10

5.6 5.8 6.0 6.2 6.4 6.6 6.81E-8

1E-7

1E-6

1E-5

1E-4

1E-3

0.01

Equ

ilibr

ium

con

stan

ts, K

2a a

nd K

2b

Reciprocal of Absolute Temperature (10000/T), K -1

K2a

K2b

Abso

rban

ce, (

cm-1) a

t 460

nm

and

500

nm

Reciprocal of absolute temperature, (104)/T

N2 (460 nm) air (460 nm) O2 (460 nm) N2 (500 nm) air (500 nm) O2 (500 nm)

Fig. 5. A comparison of the values absorbance on log scale against the reciprocal ofabsolute temperatures. The inset is a comparison of the plots of log K2a and log K2b

as a function of the reciprocal absolute temperature (T) for reactions (1a) and (1b).

X. Jiang, A. Jha / Optical Materials 33 (2010) 14–18 17

temperatures, as shown in Fig. 2, than at lower melting tempera-tures in air. In summary, it is the combination of these threeabsorption bands and their shapes at 370, 460, and 500 nm wave-lengths which affect the strength of overall absorption in GLGBglasses.

Fig. 4 shows the photoluminescence properties of selectedglasses in 700–1400 nm range using a 514 nm excitation wave-length. In general, all spectra consist of two broad PL bands peak-ing at 800 and 1200 nm. The relative intensities of the two peakschange depending upon the melting atmosphere for each glass.For glass samples melted in O2, the intensity of peak at 800 nm isstronger than that at 1200 nm. By comparison, the PL for glassesmelted in N2 and air exhibited much stronger IR emission peakat 1200 nm. From the PL analysis it was found that irrespectiveof oxidizing atmosphere over a narrow temperature range of50 �C for a given melting atmosphere, the intensity ratio of PL at1200 nm:800 nm peaks increases with the temperature. Further-more, with increasing melting temperature, the peak position inthe 1100–1300 nm region also shifts to longer wavelengths. Thisshift is the largest in the absence of oxygen atmosphere. No signif-icant infrared emission was observed in GLG glasses when excitedwith a UV–visible tungsten flash lamp source at 300 nm. Strongbroadband visible emission was observed in these glasses, whichare shown in the inset of Fig. 4.

From the Lambert–Beer’s law, the intensity of absorption isdependent on the molar concentrations of the colour centre andits extinction coefficient. Since the molar concentrations of colourcentres are dependent on the equilibrium conditions, explained viaEqs. (2a) and (2b), a plot of the absorbance values at 460 and500 nm, for example, against the reciprocal of absolute tempera-tures should yield a linear relationship for a given type of colourcentre (e.g. Bi1+ or Bi2+). This relationship is consistent with theVan’t Hoff equation, discussed by Denker et al. [30], and is shownin Fig. 5 for each type of melting atmosphere. In contrast with theprediction from the equilibrium model in Eqs. (2a) and (2b), theabsorbance versus reciprocal temperature relationship in Fig. 5 isnot linear across the temperature range of our investigation, whichsuggests that the GLGB glass hosts has definitely more than one Bi-ion states present. The graphical relationship in Fig. 5 shows morenonlinearity for oxygen and air melted samples than for the sam-ples melted in nitrogen, suggesting that the presence of nitrogen

700 800 900 1000 1100 1200 1300 14000.0

0.4

0.8

1.2

1.6

2.0

300 400 500 600 700 800 900

0.0

0.2

0.4

0.6

0.8

1.0

Nor

mal

ised

inte

nsity

[a.u

.]

Wavelength [µm]

PL spectra of GLG glass excited at

280 nm 300 nm 320 nm 340 nm 360 nm

Nor

mal

ised

PL

Inte

nsity

Wavelength [nm]

GLGB1400-O2

GLGB1375-O2

GLGB1400-Air GLGB1375-Air GLGB1350-N2

GLGB1325-N2

Fig. 4. A comparison of the room temperature photoluminescence spectra inselected Bi doped glass samples. The excitation wavelength was 514 nm. The insetis the PL spectra of a GLG glass with various excitation wavelengths in UV (280–360 nm).

in the melting atmosphere is favouring the dominance of Bi1+ state,based on the condition for Gibbs invariance discussed above. Forcomparing the computed values of equilibrium constants, log10K1a

and log10K1b with the data in Fig. 5, we have plotted the valuesagainst the reciprocal of absolute temperature in the inset of thisfigure. From the comparison of small and large slopes in Fig. 5 in-set, these two slopes may be related with the temperature depen-dence of the logarithms of the partial pressures of BiO and Bi2Ogaseous species, respectively, in equilibrium with the glass-form-ing liquid and, therefore, the concentrations, Xi, of the two colourcentres.

Another method of comparison of data for the defect state anal-ysis was also shown by Denker et al. [19,30]. From the oxygen par-tial pressure dependence of the defect states, a linear relationshipfor the absorbance in logarithmic scale against the partial pressureof oxygen is expected for a specific defect or a colour centre. Thebasis for this type of plot is the defects in oxygen deficient oxides.This is because in bismuth oxide as each oxygen (Oo) leaves the an-ion ion sub-lattice site, the concentrations of vacancies €Vo are inequilibrium with the partial pressure of oxygen gas via Eq. (3a)[34]:

Oo ¼12

O2ð Þgas þ €Vo þ 2e ð3aÞ

€Vo

h i� A PO2

� ��x � exp � DGo

ajBT

� �ð3bÞ

Here €Vo, defined in a generalized form for all oxygen deficientoxides, is the number of vacant anion sites in the sub-lattice of bis-muth oxide and 2e represents the carriers required to balance thecharge in the absence of O2� at sub-lattice. In Eq. (3b) the pre-exponent A and a within exponent are constants, and depend onthe stoichiometry of the oxide. The exponent x outside the expo-nential is dependent on the stoichiometry of the oxygen deficientoxide under consideration. The DGo is the Gibbs energy for the for-mation of defects which is represented in the atomic or ionic unitwith the Boltzmann constant, jB. As a consequence of the equilib-rium, we expect from Eq. (3b) the relationship to be linear over atemperature range, when log €Vo is plotted against the reciprocalvalue of log P(O2). As it is evident from Ref. [34], the linear regimeis explained by a constant activation enthalpy, which is likely tochange in a complex oxide such as the bismuth oxide, in accor-dance with the equilibria in (1a) and (1b). Fig. 6 characterizes the

1x10-5 1x10-4 1x10-3 1x10-2 1x10-1 1x1000.01

0.1

1

This work in glass (mol%)70 GeO2-18 Li2O-11Ga2O3-1Bi2O3

Denker et al [19,30] in glass (mol%)30 MgO-11 Al2O3-(59-x)SiO2-x Bi2O3

a(460 nm) a(500 nm) Denker et al

Abso

rban

ce a

t 500

nm

, cm

-1

log P(O2), atmospheres

Fig. 6. A plot of the absorbance in log scale against the log P(O2) for the dataobtained in this investigation for a GeO2 glass which is compared with a silicateglass from Denker et al. [19]. The circled region is where there appears to be atransition in the absorbance from linear relationship, as expected from Eq. (3b). (Seeabove-mentioned references for further information.)

18 X. Jiang, A. Jha / Optical Materials 33 (2010) 14–18

features, explained in Eqs. (3a) and (3b) and, therefore, appears tobe consistent with the equilibrium model discussed by Denkeret al. [19,30]. A close examination of data from both Denker et al.and in this work in Fig. 6 show a slope change in the vicinity ofP(O2) � 0.2–0.4 atm. As there are fewer number of data points forcarrying out activation enthalpy analysis, it is difficult to concludeunambiguously which of the two valence states of bismuth iondominate the colour centres over a temperature range. FromFig. 4, where the emission data for various melting atmospheresare compared, it may be concluded that the shorter wavelengthemission may be due to the dominance of Bi2+ states in combina-tion with the Bi1+ state, and the longer IR emissions may be dueto the preponderance of Bi1+ in combination with the Bi2+ states.This unambiguous conclusion is possible by verifying the emissiondata in Fig. 4 against the increasing concentrations of oxygen in themelting atmosphere, where we expect more Bi2+ states to be pres-ent than the Bi1+ states, as explained in Eqs. (2a) and (2b). Thisfinding is also consistent with the model proposed in Eqs. (3a)and (3b). It should also be noted that our data is not expected toagree completely due to the apparent difference in the glass com-position between the present work and that of Denker et al. [19].The difficulty in ascertaining the regime of colour centre stabilityas a function of temperature and oxygen pressure has beenencountered as a result of a lack of precise control of P(O2) forachieving the condition for Gibbs invariance at a given tempera-ture and pressure.

4. Conclusions

In summary, we have observed NIR broadband emissions in bis-muth-doped germanate (GLGB) glasses. The influence of gaseousatmosphere and temperature during glass melting on PL at near-IR has been studied in detail and analyzed in the context of suitableanion defect model, explained in Eqs. (3a) and (3b). From the re-sults, discussed above, it is the lower valence states of Bi-ions,Bi1+/Bi2+, which are largely responsible for the near-IR PL at1200–1500 nm. By comparison under more oxidizing conditions,

the photoluminescence is more dominant at 800 nm, due to thedominance of equilibrium condition for the stability of Bi2+ states,with respect to the Bi1+ states. This implies that the origin of the800 nm PL might be attributable to the dominance of Bi2+ defectsstates in the GLGB glass. The GLG glass, on the other hand, doesnot exhibit any PL at IR, instead a strong visible PL is observeddue to the presence of absorption band at around 370 nm, asshown in Fig. 4 inset.

Acknowledgements

The authors acknowledge the financial support from the Glaxo-SmithKline plc and EPSRC Multi-core fibre (EP/C515234/1) project.We also sincerely thank Dr. Denker for sending data for makingcomparisons in Figs. 5 and 6.

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