Effect of iron impurity on structural development in ball-milled ZrO2–3 mol% Y2O3

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CERAMICS INTERNATIONAL

Ceramics International 41 (2015) 1121–1128 www.elsevier.com/locate/ceramint

Effect of iron impurity on structural development in ball-milled ZrO2–3 mol% Y2O3 Ali Saria,b,n, Mourad Keddamb, Abdurahim Guittoumc a Nuclear Research Center of Birine, Box 180, Ain Oussera, Algeria Laboratoire de Technologie des Matériaux, Département de Sciences des Matériaux, Faculté de Génie Mécanique et Génie des Procédés, USTHB, BP No 32, 16111 El-Alia, Bab-Ezzouar, Algiers, Algeria c Nuclear Research Center of Algiers, 2 Bd Frantz Fanon, Box 399, Alger-Gare, Algiers, Algeria

b

Received 15 June 2014; received in revised form 1 September 2014; accepted 8 September 2014 Available online 19 September 2014

Abstract Powders Mixture of monoclinic ZrO2 and 3 mol% Y2O3 were prepared by high-energy ball milling. Ball-milled products were investigated by means of X-ray diffraction, Rietveld refinement, spectrophotometry and Mössbauer spectroscopy. The tetragonal Zr1–xYxO1.982 (x ¼0.037) was formed within 30 min of ball milling in which the crystallite size decreases up to 13 nm and appears almost unchanged even after 30 h of milling. Peak intensities of the monoclinic phase decrease and disappear after 10 h of milling. The unit cell-volume of tetragonal phase decreases with increasing iron content incorporated from milling media. In the calcined product ball-milled for 30 h the monoclinic phase takes place in the range 700–900 1C and disappears at 1100 1C due to the presence of Y3 þ cations. It was concluded that at room temperature, the stabilization of tetragonal phase at the beginning of milling process was attributed to the substitutional Fe3 þ cations, while the hematite Fe2O3 was formed at the end of ball milling. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: A. Milling; Rietveld method; Iron impurity; Mössbauer spectroscopy; Zirconia

1. Introduction Pure ZrO2 exhibits three polymorphous transformations: monoclinic (m-ZrO2) below 1100 1C, tetragonal (t-ZrO2) between 1200 1C and 2360 1C and cubic (c-ZrO2) up to the melting point [1]. High temperature polymorphs of ZrO2 can be stabilized at room temperature by the addition of aliovalent oversized dopant notably Ca2 þ , Mg2 þ and Y3 þ [2]. The enhancement of strength and toughness through polymorphous transformation in partially stabilized zirconia, led to important applications in oxygen sensors, fuel cells, jewellery etc., thanks to the high oxygen ion conduction in yttria-doped zirconia [3–5]. Depending on the doping concentration, the oversized aliovalent such as Y3 þ causes a stabilization of t-ZrO2 (smaller amount of oxygen vacancies) or n Corresponding author at: Nuclear Research Center of Birine, Box 180, Ain Oussera, Algeria. Tel./fax: þ213 5 50 85 77 43. E-mail address: [email protected] (A. Sari).

http://dx.doi.org/10.1016/j.ceramint.2014.09.038 0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

c-ZrO2 (higher amount of oxygen vacancies) polymorphs of ZrO2 by reducing the coordination number of Zr4 þ cations due to the introduction of oxygen vacancies [6–8]. In yttria-doped zirconia, trivalent dopant cations Y3 þ substitute for some of the Zr4 þ cations and, in order to keep the charge in balance, one O vacancy must be created for each pair of dopant cation. The usual way of introducing the oxygen vacancies into the ZrO2 lattice is to form a solid solution with aliovalent oxides [6]. However, as shown in the work of Li [9], this will only occur when the dopant cations have larger ionic sizes than the Zr4 þ cations. Nevertheless, t-ZrO2 and c-ZrO2 phases can also be stabilized at room temperature by the incorporation of the undersized Fe2 þ and/or Fe3 þ cations in the ZrO2 lattice [10], while this stabilization has not been efficient due to their smaller ionic radius in comparison with that of the Zr4 þ cation [8]. Effect of iron impurity on the polymorphous transformation in zirconia has rarely been reported, however, in most cases Fe2O3 was added purposely in small amounts (up to 3 mol%) to the Y-TZP in order to lower the sintering temperature

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required for a relative density 495% or, to regulate the grain size and grain growth [11,12] and also, to probe the change in electrical and ionic conductivity of the material such as oxygen sensor [13]. Most of the later investigations neglected the effect of impurities and explained the observed polymorphous transformation of zirconia simply in terms of surface energy changes [8]. Among several methods of preparations, high-energy ball milling is a very easy solid-state processing technique for addition of Fe2 þ and/or Fe3 þ cations in the zirconia lattice for exhibiting new and unusual properties [14–16]. Murase and Kato [17] examined the conversion of t-ZrO2 to m-ZrO2 during the ball milling in different atmospheres and found out that the absence of oxygen or water has significantly increased the stability of a metastable t-ZrO2. Gajovic et al. [18] shown that the ball milling has little or no influence on the stability of t-ZrO2 when the zirconia balls and vials were used. Stefanic et al. [19] observed that the presence of impurities in the ball-milled m-ZrO2 powder significantly influenced the stability of the t-ZrO2 product. It was found that addition of iron oxide facilitates the activation of diffusion processes at grain boundaries, and significantly reduces the sintering temperature in the materials of the system ZrO2–Y2O3–Fe2O3 [20]. Verkerk et al. [21] studied the effect of impurities on sintering and conductivity of yttria-stabilized zirconia. They observed a decrease in the unit-cell volume of cubic phase 8YSZ by doping with Fe2O3. However, the substitution of Fe3 þ cations is similar with the Y3 þ cations. So, the substitution of Zr4 þ cations by the undersized Fe3 þ cations led to a decrease in the unit-cell volume [22]. It was found that undersized dopant Fe3 þ has a lower solubility (less than 3 mol%) in zirconia, than the oversized one Y3 þ ; their supersaturated tetragonal phase can be maintained only at relatively low calcination temperatures [23]. Belous et al. [24] studied the structure features of zirconia stabilized by a complex dopant (Y2O3 and Fe2O3) as a function of the chemical composition, precipitation, heat treatment, and ageing conditions. They found that the solubility of iron in zirconia increases with yttrium content. Raming et al. [25] found that in fully crystalline tetragonal zirconia (Y-TZP), only small amounts of Fe3 þ (o1 mol%) could be dissolved into the zirconia lattice. Stefanic et al.[19] found that the solubility limit of the Fe3 þ cations into the zirconia decreases with increasing calcinations temperatures, inducing Fe3 þ cations to segregate out of zirconia and form a α-Fe2O3 phase containing probably some Zr4 þ cations [14]. The aim of the present work was to study the effect of the iron impurity on structural and microstructural changes in the ball-milled ZrO2–3 mol% Y2O3 powders mixture using steel ball milling assembly. Ball-milled products were investigated by means of X-ray diffraction, Rietveld refinement, spectrophotometry and Mössbauer spectroscopy. 2. Materials and methods

NYCO, France. The average diameter of particles was about 44 mm. Ball milling was performed up to 30 h in a planetary ball mill (Model PM200, Retsch, GmbH, Germany) with stainless steel balls and vials. The ball-to-powder weight ratio and the rotation speed of the vials (125 ml) were fixed to 20:1 and 500 rpm, respectively. To avoid excessive heating during milling, each 1 h of milling was followed by a pause of 30 min. After ball milling for given times, batches of 300 mg of product were taken out from the vials for crystal structure analysis using Panalytical X’Pert Pro MPD (Cu Kα) with a step size of 0.021 and 2θ range from 151 to 901. Ball-milled powders from 30 min to 30 h were analyzed with AAnalyste 400 spectrophotometer (Perkin Elmer) to estimate amounts of iron and chromium impurities. For each ball milling time, 0.5 g was taken from ball-milled products and diluted in an acid solution (HNO3, HF, HClO4 and H2SO4) and then heated at 150 1C. The resulting precipitate was removed by filtration and washed repeatedly in doubly distilled water. The products ball-milled for 5 and 30 h were selected for Mössbauer spectroscopy analysis to determine the valence state of iron impurity and solid solution that takes place during the milling process. For this purpose, Mössbauer spectroscopy was carried out at room temperature with a Wissel instrument for a constant acceleration mode, using a radioactive 57Co source diffused into a Rhodium matrix. Metallic iron was used for energy calibration and also as a reference for isomer shift. Mössbauer spectra were recorded by means of Recoil software [26]. For thermal stability study, batches of 300 mg taken from the selected product milled for 30 h were calcined at the Table 1 Structural and microstructural parameters deduced from the Rietveld refinement of the XRD patterns of unmilled and ball-milled ZrO2–3 mol% Y2O3 (X’Pert HighScore plus). Milling Phase Phase time present fraction (%)

Crystallite Unit-cell volume (A3) size (nm)

Microstrains GoF (%)

0h

– – – 67.25 – – 67.21 – – 67.15 – 66.97 – 66.86 – 66.81 66.63 – 66.50 – 66.35 –

0.156 0.033 0.255 0.540 0.253 0.420 0.713 0.261 0.558 0.686 0.608 0.673 0.554 0.673 0.486 0.651 0.673 0.265 0.904 0.228 0.870 0.178

0.5 h

1h

2h 3h 4h 5h 10 h 20 h

2.1. Material 30 h

The starting powders used were monoclinic ZrO2 (purity 99.995%) and Y2O3 (purity 99.99%) which were supplied by

M c-Y2O3 M T c-Y2O3 M T c-Y2O3 M T M T M T M T T α-Fe T α-Fe T α-Fe

95 5 80 16 4 70 27 3 49 51 38 62 27 73 19 81 95 5 92 8 90 10

49.6 95.2 48.5 19.6 48.0 33.8 27.7 31.8 21.6 29.6 17.1 33.5 14.7 36.0 14.2 40.6 13.0 8.6 11.9 10.3 11.1 10.9

M¼monoclinic (m-ZrO2); T ¼tetragonal (t-Zr0.963Y0.037O1.982).

1.98 1.56

1.76

1.74 1.65 1.51 1.54 1.52 1.33 1.10

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temperatures of 500, 700, 900 and 1100 1C for 2 h in air at standard pressure. The calcination was performed in the chamber furnace Anton Paar HTK12. A heating and cooling rate of 30 1C/min was used in this heat treatment. 2.2. Method of structure refinement Rietveld method [27] was adopted to refine structural and microstructural parameters of unmilled, ball-milled and calcined powders mixture. The procedure of calculations consists in modeling the diffraction profiles by analytical functions using the X’Pert Highscore plus software v. 2.1.0 [28,29]. The structural parameters were described through the measurement of phase fraction, unit-cell volume, while the refinement of microstructural parameters led to measure crystallite sizes and lattice microstrains [30]. In principle, the Rietveld method is based on Eq. (1): k p2

yic ¼ ∑ ∑ Gpik I k þ yib p k ¼ kp 1

ð1Þ

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where ‘yic ’ is the net intensity calculated at point ‘i’ in the pattern, ‘yib ’ is a polynomial function for reproducing the background intensity, ‘Gik ’ is a normalized peak profile function modeled by pseudo-Voigt function (Eq. (2)), ‘I k ’ is the intensity of the ‘kth’ Bragg reflection, ‘k1… k2’ are the reflections contributing intensity to point ‘i’, and the superscript ‘p’ corresponds to the possible phases present in the sample. i1 h i C0 h 1 þ C0 X 2jk þ ð1 γ Þexp  C 1 X 2jk Hk π 1=2

Gik ¼ γ

ð2Þ

where C0 ¼ 4, C 1 ¼ 4ln2, γ is a refinable “mixing” parameter, X jk ¼ ð2θi  2θk Þ=H k where H k is the half-width at half-maxima of the kth Bragg reflection defined by the function described as [31]  1=2 H k ¼ U tan2 θ þ V tan θ þ W ð3Þ The experimental profile was built according to Eq. (1) and adjusted with least-squares process until the best fit to the experimental diffraction pattern is obtained. However, size and

Fig. 1. Refined patterns (Rietveld software HighScore plus) of unmilled and ball-milled products as a function of milling time of: (a) 0 h, (b) 5 h, (c) 10 h and (d) 30 h. Small symbol markers at the top of each diffraction line indicate the position of all the phases identified before and after ball milling process. Experimental data points are shown as dot (Io), calculated patterns are shown as continuous lines (Ic). Residual of fitting (Io–Ic) between observed (Io) and calculated (Ic) intensities of each fitting is plotted as a continuous line in the lower field.

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strain broadened experimental profiles were fitted by a model based on a pseudo-Voigt approach (Eq. (2)) [32]. Before any attempt to refine XRD patterns of un-milled and ball-milled products, a Silicon SRM 640c certified for line position and line profile for powder diffraction was used as an instrument standard [33]. XRD pattern of Si was recorded on the same instrument using the same settings as used later to measure the ball-milled product. Refining XRD pattern of Si and its atomic coordinates (x, y, z), allows to collect the refined profile and shape parameters which are assumed to be a pseudo-Voigt function with asymmetry (γ a1) [34]. The approximated background ‘yib ’ (Eq. (1)) of each pattern is fitted by a polynomial function of fourth degree. The obtained values of profile parameters of SRM 640c were used as input data in each phase of un-milled and ballmilled product for Rietveld analysis. To ovoid divergences in the least-squares during fitting process, the coefficient of Cagliotti V and the atomic coordinates of special Wyckoff positions were not refined. Structure refinement continues until the goodness of fit (GoF) approaches the value of 1 [35–37]. During fitting process, quantitative phase analysis was performed for each phase of the unmilled and ball-milled products according to the CHUNG method (program HighScore plus) [38]. Phase identification was achieved by means of X’Pert Highscore plus software, supported by data from PDF-2 database (ICDD-International Centre for Diffraction Data, Newtown Square, PA). In order to perform the Rietveld refinement we need structure data for all identified phases: (i) for m-ZrO2 (monoclinic, ICDD PDF # 01-083-0938), the space group was taken as P21/c (14), with respectively, Zr, OI and OII atoms in general Wyckoff positions (4e) having different x, y and z values for different atoms. (ii) for t-Zr0.963Y0.037O1.982 (tetragonal, ICDD PDF # 01-0830113), the space group was taken as P42/nmc (1 3 7), with O atom in general Wyckoff positions (4d) and both Zr and Y atoms in special Wyckoff positions (2b). (iii) for α-Fe (cubic, ICDD PDF # 01-087-0722), the space group was taken as Im-3m (2 2 9), with Fe atom in special Wyckoff positions (2a). (iv) for α-Fe2O3 (rhombohedral, ICDD PDF # 01-073-2234), the space group was taken as R-3c (1 6 7), with both Fe and O atoms in general Wyckoff positions (12c) and (18e), respectively. (v) for c-Y2O3 (cubic, ICDD PDF # 01-083-0927), the space group was taken as la-3 (2 0 6), with Y(1) atom in special Wyckoff positions (8a), Y(2) and O atoms in general Wyckoff positions (24d) and (48e), respectively.

content and other impurities are below the detection limit of X-ray diffraction (Fig. 1a and b). According to the results from the analysis by spectrophotometry, we have neglected the effect of chromium in term of impurity because of its small amount compared to that of the iron. Before milling, the ZrO2–3 mol% Y2O3 powders mixture consists substantially of 95% m-ZrO2 and 5% c-Y2O3 with crystallite sizes equal to 50 nm and 95 nm, respectively (Table 1). In the course of ball milling, the peak intensities of the c-Y2O3 reflections decrease rapidly and completely disappear after 2 h of milling (Table 1). This indicates the diffusion of yttrium atoms into the m-ZrO2 matrix in accordance with Ref. [6]. It is well known that Y2O3 has a large solid solubility range in ZrO2 and can be used to stabilize the tetragonal phase of (Y2O3)x(ZrO2)1  x over the composition range 0.02 o xo 0.09 and the cubic phase with 0.04o x o 0.4 [6]. However, peak intensities of the m-ZrO2 reflections have been reduced by the ball milling process, which results in a decrease of crystallite size up to 14 nm and an increase of microstrains up to 0.486% (Table 1). t is further noticed that a decrease in the phase fraction of m-ZrO2 was attributed to the

Fig. 2. Effect of milling media on the phase transformation in the ZrO2–3 mol% Y2O3 ball-milled for 30 min.

3. Results and discussions The results from Rietveld refinement of the XRD patterns recorded for different milling times are summarized in Table 1. The values of the goodness of fit of all the fitted patterns lie within 1.10–1.98. Some of the selected refined patterns corresponding to 0 h (a), 5 h (b), 10 h (c) and 30 h (d) of milling, are reported in Fig. 1. At early stage of ball milling, peaks of the α-Fe phase are not observed because the iron

Fig. 3. Variation of molar fraction of iron impurity and the unit-cell volume of t-Zr0.963Y0.037O1.982 phase with increasing milling time.

A. Sari et al. / Ceramics International 41 (2015) 1121–1128

onset of the m-ZrO2 to tetragonal Zr1  xYxO1.982 (x¼ 0.037) transition [39] which occurred after 30 min of milling (Table 1). Stafanic et al. [19] investigated the phase development in pure zirconia and found that the onset of m-ZrO2 to t-ZrO2 transition occurred between 10 and 15 h of ball milling and that the Fe3 þ cations stabilized the t-ZrO2 phase at room temperature. On the other hand, it has been shown that the presence of undersized Fe3 þ in yttria-doped zirconia structure promote the mobility of Y3 þ cations which diffuse into the tetragonal crystal lattice and that the substitution of Fe3 þ cation is similar with the Y3 þ cation [24]. Ongoing of our preliminary results obtained from ball-milled ZrO2–3 mol% Y2O3 with zirconia balls and vials (Fig. 2), no tetragonal reflections were detected after 30 min of milling. The results as given in Figs. 1 and 2 indicate that the iron impurity facilitates the dissolution of yttrium in the zirconia matrix under the formation of a tetragonal phase in which also some iron is dissolved. As can be seen from Fig. 3, the molar fraction of iron impurity increases slightly from 0.01 to 1 mol %, which is in good agreement with the previous works [16,40]. Such evolution is due to the smaller ionic radius of Fe3 þ (0.78 A) compared to that of Zr4 þ (0.84 A) [41]. The same effect was also observed by Stefanic et al. [42]. They have explained the decrease of the unit- cell volume of the t-ZrO2 through a substitution of the Zr4 þ cations by the undersized Fe3 þ ions. However, there are two types of sites Table 2 Mössbauer parameters at room temperature of products ball-milled for 5 and 30 h. Time

Phase (ions)

IS (mm/s)

QS (mm/s)

Hhf (T)

SP (%)

5h

α-Fe (Fe3 þ )paramag α-Fe (Fe3 þ )paramag Hematite

0.003 0.377 0.00 0.386 0.745

0.002 0.896 0.002 0.850 2.600

32.882 — 32.747 — —

58.9 41.1 36.2 51.6 12.2

30 h

(10) (10) (32) (31) (23)

IS ¼isomer shift; QS¼ quadrupole splitting; Hhf ¼ hyperfine field; SP¼site populations of component; (Fe3 þ )paramag ¼Fe3 þ in paramagnetic state.

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for Fe3 þ ion occupation in structure lattice of yttria-doped zirconia, which are substitutional sites [43] and interstitial sites [44]. Verkerk et al. [21] suggested a decrease of the unit-cell volume, when the Fe3 þ cations substituted for the Zr4 þ cations, but the Fe3 þ interstitials give rise to an increase of the unit-cell volume [21]. It becomes clear that in our study, the Fe3 þ cations occupied a substitutional site in the t-Zr0.963Y0.037O1.982 lattice. These results agree with those obtained with Rietveld structure refinement. It is interesting to determine the valence of iron impurity and identify the solid solution which could be formed inside the ballmilled product. In this respect, ball-milled ZrO2–3 mol% Y2O3 for 5 and 30 h were selected for Mössbauer spectroscopy analysis. The Mössbauer parameters recorded at room temperature are given in Table 2. The spectrum of the ball-milled product for 5 h (Fig. 4a) is characterized by a doublet with an Isomeric Shift (IS) equal to 0.377 mm/s and the quadrupole splitting (QS) equal to 0.896 mm/s (Table 2), indicating the presence of Fe3 þ component [45]. The (QS) value of Fe3 þ component can also be assigned to the formation of solid solution of t-Zr0.963(Y3 þ ,Fe3 þ )0.037O1.982 inside a m-ZrO2 matrix [24]. The sextet with a magnetic hyperfine field, oHhf 4, equal to 32.882 T is characteristic of α-Fe and this value agrees with the reference works [46–49]. Nevertheless, the spectrum of the ball-milled product for 30 h (Fig. 4b) shows a superposition of two doublets, one typical of Fe3 þ component with (IS) and (QS) values equal to 0.39 mm/s and 0.85 mm/s, respectively (Table 2) and the other typical of hematite (Fe2O3) with (IS) and (QS) values equal to 0.75 mm/s and 2.6 mm/s, respectively, whose parameters agree with those published for hematite [50–52]. On the basis of analysis of the spectra parameters, the peak areas of Mössbauer spectra show a superposition of α-Fe (59%) and Fe3 þ (41%) components recorded after 5 h of milling and a superposition of α-Fe (36%), Fe3 þ (52%) and hematite (12%) components recorded after 30 h of milling (Table 2). From these results, it is clearly seen that the population of Fe3 þ component increases with increasing milling time, indicating an increase of the paramagnetic state characteristic of Fe3 þ component at the expense of ferromagnetic state characteristic of iron (α-Fe) [50].

Fig. 4. Iron-57 Mössbauer spectrum recorded from ZrO2–3 mol% Y2O3 ball-milled for 5 h (a) and 30 h (b) collected at room temperature.

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Fig. 5. Refined patterns (Rietveld software HighScore plus) of ball-milled product for 30 h after calcination at different temperatures for 2 h: (a) 500 1C, (b) 700 1C, (c) 900 1C and (d) 1100 1C. Small symbol markers at the top of each diffraction line indicate the position of all phases identified before and after ball milling process. Experimental data points are shown as dot (Io), calculated patterns are shown as continuous lines (Ic). Residual of fitting (Io–Ic) between observed (Io) and calculated (Ic) intensities of each fitting is plotted as a continuous line in the lower field.

Table 3 Structural and microstructural parameters deduced from the Rietveld refinement of XRD patterns of the product ball-milled for 30 h calcined for 2 h at 500, 700, 900 and 1100 1C (X’Pert HighScore plus). Milling time

Phase Phase present fraction (%)

Unit-cell volume (A3)

Crystallite size (nm)

Microstrains GoF (%)

30 h (500 1C)

T

43

66.16

13.0

0.448

1.10

H M

57 16

– –

66.5 48.3

0.194 0.125

1.84

T H M

38 46 8

66.42 – –

23.2 35.8 51.8

0.232 0.242 0.024

1.28

T H T

33 59 43

66.94 – 67.05

41.7 134.7 75.9

0.135 0.052 0.101

1.14

H

57

–

537.2

0.025

30 h (700 1C)

30 h (900 1C)

30 h (1100 1C)

M¼ monoclinic (m-ZrO2); T ¼tetragonal (t-Zr0.963Y0.037O1.982); H ¼hematite (α-Fe2O3).

For structural stability study, the product ball-milled for 30 h was subjected to calcination for 2 h between 500 and 1100 1C and the XRD patterns are reported in Fig. 5. The results of structure and microstructure refinement are given in Table 3. It is worth to note that in the calcined product at 500 1C (Fig. 5a), the tetragonal phase (43%) contains a certain proportion of metastable t-(Zr4 þ ,Fe3 þ )O2 phase and a large proportion of stable t-Zr0.963(Y3 þ ,Fe3 þ )0.037O1.982 phase (Table 3). Interestingly, the annealing treatment above 700 1C (Fig. 5b) causes an impoverishment of iron impurity explained by the diminution of tetragonal fraction up to 38% due to the conversion of metastable t-(Zr4 þ ,Fe3 þ )O2 into thermodynamically stable m-ZrO2 polymorph (Table 3). Calcination at 900 1C causes almost a complete release of Fe3 þ ions, while the m-ZrO2 phase decreases in favor of t-Zr0.963Y0.037O1.982 phase (Fig. 5c). As shown in Fig. 5d, the monoclinic m-ZrO2 disappears at 1100 1C due to the diffusion of Y3 þ cations within the m-ZrO2 lattice [6]. Indeed, in order to maintain the charge in balance between the Zr4 þ and Y3 þ cations, oxygen vacancy must be created as a result of calcination for each pair of Y3 þ cations into the ZrO2 lattice [8]. In addition, the introduction of oxygen vacancies would be

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probably associated with the Fe3 þ cations giving rise to the hematite coalescence. This is in agreement with the results of phase fraction of hematite varying from 46 to 59%, when the calcination temperature goes up (Table 3). The obtained results from calcination confirm the formation of both t-(Zr4 þ ,Fe3 þ )O2 and t-Zr0.963(Y3 þ ,Fe3 þ )0.037O1.982 solid solution. 4. Conclusion Effect of iron impurity on structural and microstructural changes in the ball-milled ZrO2–3 mol% Y2O3 powders mixture was investigated. The tetragonal phase was formed after 30 min of milling, which corresponds to the Zr1 xYxO1.982 (x¼ 0.037) form. It is found that the decrease in the unit-cell volume of the tetragonal phase is due to the Fe3 þ ions dissolved into the tetragonal lattice. During calcination, the monoclinic phase forms at approximately 700 1C and decreases up to 900 1C due to the large solubility of Y3 þ in monoclinic m-ZrO2. At 1100 1C, the calcined product contains 40% of tetragonal and 60% of α-Fe2O3. It was concluded that at room temperature, the stabilization of tetragonal phase at the beginning of milling process was attributed to the substitutional Fe3 þ cations, while the hematite Fe2O3 was formed at the end of ball milling. Acknowledgments The authors wish especially to thank (a) Prof. M. Azaz and Dr. S. Triaa (from USTHB, Algiers, Algeria), (b) Dr N. Souami (Nuclear Research Center of Algiers) and (c) Drs. F. Mernache and S. Ladjouzi (Nuclear Research Center of Draria) for their kind help during the experiments and for their valuable suggestions during the preparation of this work. References [1] A. Fujio, M. Seiichi, Y. Koichi, Tetragonal to monoclinic transformation and microstructural evolution in ZrO2–9.7 mol% MgO during cyclic heating and cooling, J. Mater. Sci. 32 (1997) 513–522. [2] G. Herrera, N. Montoya, A. Doménech-Carbo, J. Alarcon, Phys. Chem. Phys. 15 (2013) 19312–19321. [3] E.C. Subbarao, Zirconia– An overview, In advances in Ceramics, in: A. H. Heuer, L.W. Hobbs (Eds.), Science and Technology of Zirconia, American Ceramic Society, vol. 3, Columbus, OH, 1981, pp. 1–24. [4] D. Simwonis, H. Thülen, F.J. Dias, A. Naoumidis, D. Stöver, Properties of Ni/YSZ porous cermets for SOFC anode substrates prepared by tape casting and coat-mix process, J. Mater. Process. Technol. 1 (1993) 107–111. [5] M.J. Verkerk, B.J. Middelhuis, A.J. Burggraaf, Effect of grain boundaries on the conductivity of high-purity ZrO2–Y2O3 ceramics, Solid State Ionics 6 (1982) 159–170. [6] S.A. Ostanin, E.I. Salamatov, Microscopic mechanism of stability in yttria-doped zirconia, JETP Lett. 74 (2001) 552–555. [7] H. Shih-Ming, On the structural chemistry of zirconium oxide, Mater. Sci. Eng. 54 (1982) 23–29. [8] G. Stefanic, B. Grzeta, S. Music, Influence of oxygen on the thermal behaviour of the ZrO2–Fe2O3 system, Mater. Chem. Phys. 65 (2000) 216–221. [9] P. Li, I.W. Chen, E. James, Penner-Hahn, Effect of dopants on zirconia stabilization—an X-ray absorption study: I, trivalent dopants, J. Am. Ceram. Soc. 77 (1994) 118–128.

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