Influence of zirconia addition on the properties of magnesia refractories

Ceramics International 42 (2016) 11373–11386

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Influence of zirconia addition on the properties of magnesia refractories Robert Kusiorowski n, Józef Wojsa, Bronisław Psiuk, Teresa Wala Institute of Ceramics and Building Materials, Refractory Materials Division in Gliwice, ul. Toszecka 99, 44-100 Gliwice, Poland

art ic l e i nf o

a b s t r a c t

Article history: Received 23 March 2016 Received in revised form 12 April 2016 Accepted 13 April 2016 Available online 14 April 2016

This paper concerns basic refractories used, among others, in the cement industry. Zirconium oxide was added to modify the properties of an MgO refractory. These MgO–ZrO2 materials can replace magnesiachromite refractory ceramics in accordance with current environmental policy guidelines and can improve the properties, such as thermal shock resistance, of a typical magnesia refractory. In the experiments, commercially available magnesia clinker and monoclinic zirconia were used. The effect of particle size and amount of zirconium oxide content were determined in the prepared ceramic masses. The addition of zirconia was at a level of from 0 to 8 wt% in green compacts. After the firing cycle at 1660 °C the ceramic properties were measured, such as linear shrinkage, open porosity, apparent density, cold crushing strength, gas permeability and thermal shock resistance (water cycles). For selected materials, resistance to corrosion by two different corrosive agents was determined. The research results show that the addition of zirconia improved the refractories' properties. The cold crushing strength of samples with 2 wt% of added ZrO2 reached maximum values. The obtained ceramics were also characterised by good dimensional stability (firing shrinkage less than 1%), high cold crushing strength (50–75 MPa) and uncommonly high values of thermal shock resistance (8–10 water cycles). The addition of ZrO2 also clearly increased the materials’ corrosion resistance against a cement-like and glass-like agent. & 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: C. Corrosion D. MgO D. ZrO2 E. Refractories

1. Introduction Magnesium oxide (MgO) is the main component of basic refractory materials, which are widely used in thermal devices in steel or non-ferrous metals metallurgy as well as in other industrial sectors where magnesia-based refractories are used, i.e. as lining in the rotary kilns of the cement industry or as a part of the glass furnace regenerator [1,2]. Due to the basic chemical nature of MgO, magnesia refractory products have very good chemical resistance to alkaline agents and operating conditions, thus these products are resistant to basic slag, lime and cement clinker. At the same time, they quickly corrode in an acidic environment. The disadvantage of magnesia bricks is their sensitivity to stress mainly caused by thermal shock, which is explained by the high thermal expansion coefficient of periclase – the main mineral phase in this refractory material. For most magnesia products, thermal shock resistance is very small and often does not exceed several air changes [3]. The traditional solution to this problem, which has allowed to n

Corresponding author. E-mail addresses: [email protected], [email protected] (R. Kusiorowski). http://dx.doi.org/10.1016/j.ceramint.2016.04.065 0272-8842/& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

improve the resistance of material to sudden temperature changes, was based on the introduction of modifying additives. Chrome ore was often introduced in the initial period of modifying magnesia refractories. Magnesia-chromite bricks have been the standard lining for hot cementitious kiln zones since about 1940 [4]. These developed magnesia-chromite refractories show excellent thermal shock resistance, satisfactory corrosion resistance in many processes and good hot strength [1]. Despite the advantageous properties of magnesia-chromite refractories, over the years a significant environmental problem has appeared. Oxidising conditions may occur during the operation of refractory materials containing chromium at high temperatures and, consequently, a toxic hexavalent form of chromium Cr(VI) is formed. According to Lee and Nassaralla [5], the formation of hexavalent chromium is favoured by a high concentration of CaO in the reaction system and by cyclical temperature changes (heating and cooling) of the refractory lining. Cr(VI) is formed at the phase boundary of Ca-rich slag and chromite grains during cooling of the furnace lining below
1230 °C. According to the International Agency for Research on Cancer (IARC), compounds containing the hexavalent form of chromium are classified as carcinogenic and mutagenic substances. The toxic effect of Cr(VI) is due to its strong oxidising properties and free radical formation during its reduction in living cells. Furthermore,

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the compounds of Cr(VI) are easily water soluble and bio-available, which enhances their toxicity [6,7]. To avoid the problem with the presence of hexavalent chromium in refractory materials, another way to modify the MgOCr2O3 refractory was developed. This solution focused on alumina addition, so the Mg–Al spinel (MgO  Al2O3) was created during the firing cycle of refractory ceramics. This compound is characterised by a lower thermal expansion coefficient in comparison to periclase grains (8–9  10  6 °C  1 vs. 13–15  10  6 °C  1, respectively [8]). The difference in this parameter makes the magnesia-spinel material structure discontinuous, thus resulting in good flexibility and a faculty for the prevention of crack propagation [9]. On the other hand, the aluminium form of spinel can relatively easily react with calcium compounds from cement raw materials and products, and calcium aluminates with a low melting point (below 1500 °C) can be created [10]. These new phases may contribute to premature wear of the kiln lining. With the drastic reduction in the use of MgO–Cr2O3 materials for both medical and ecological reasons, further magnesite modification research studies have been conducted. One of these modifying agents could be zirconium oxide (ZrO2) as an additive to the ceramic mass. Zirconia is characterised by a very slight tendency to enter compounds with other oxides, so it has very high resistance to slag corrosion. It also exhibits a very high melting point (above 2700 °C). Moreover, it can react with calcium oxide (CaO) and, in consequence, high-refractory calcium zirconate (CaZrO3) can be formed. There have not been many reports which have focused on MgO–ZrO2 materials, and these are practically related with determining corrosion resistance against different agents [11–13]. Some information in this area is provided in the paper by Ceylantekin and Aksel [14], where among the magnesia-spinel composites with the addition of zircon (ZrSiO4), a MgO–ZrO2 refractory containing up to 30 wt% of zirconia was also described. The mechanical strength of the material decreased after significant addition of zirconia to the mass. The aim of the study was to determine the influence of the amount and grain size of zirconium oxide in ceramic masses on the properties of modified magnesia refractory material. The development of MgO–ZrO2 refractories, which are characterised by high thermal shock resistance and high corrosion resistance, was studied under cement kiln as well as glass-furnace regenerator conditions, which could be a replacement for currently used magnesia-chromite products.

2. Experimental

Table 1 Chemical analysis of the raw materials (in wt%). Materials

Magnesia clinker

Zirconia (medium–large)

Zirconia (fine)

Symbol MgO CaO SiO2 Al2O3 Fe2O3 ZrO2 HfO2 L.O.I. Total

N 97.47 0.97 0.56 0.21 0.45 – – 0.29 99.95

Zr1 0.21 0.53 0.37 0.20 0.12 96.29 1.62 0.20 99.54

Zr2 0.22 0.53 0.34 0.05 0.12 96.43 1.62 0.30 99.61

Table 2 Compositions of the samples. Sample designation

0 2Zr1 4Zr1 6Zr1 8Zr1 2Zr2 4Zr2 6Zr2 8Zr2

Raw material content; wt% N

Zr1

Zr2

100 98 96 94 92 98 96 94 92

– 2 4 6 8 – – – –

– – – – – 2 4 6 8

2.2. Ceramic sample preparation and sintering Refractory samples were prepared by adding different amounts of zirconia (0–8 wt%) to the MgO clinker. The materials were weighed according to the compositions shown in Table 2. The grain size distribution of all types of ceramic masses was fitted according to the Dinger and Funk model [15] for a grain distribution coefficient n equal to 0.37. To prepare the green compact, the weighed raw materials were mixed well with 3.5 wt% addition of sulphite waste liquor. Subsequently, the samples were shaped into cylindrical samples 50 mm in diameter and 50 mm in height under pressing at 120 MPa with a pressure vent at 60 MPa. Then the green samples were dried in natural conditions at room temperature and in a laboratory dryer at 120 °C for 4 h. The dried samples were sintered to establish their final properties. This process was conducted in a high-temperature laboratory gas furnace at 1660 °C with 4 h of soaking time at maximum temperature. Finally, the ceramic samples were naturally cooled until they reached room temperature.

2.1. Raw materials 2.3. Properties of obtained refractories and methods The study presented here was focused on the influence of both grain size and amount of zirconia on the final properties of magnesia-zirconia refractories. The raw materials used in this study were sintered magnesia (Magnesita; Brazil) as well as commercially available zirconia. According to the technical data sheet, magnesia clinker was obtained based on double firing of natural magnesite with a high content (Z98 wt%) of MgO. The particle size of the applied magnesia clinker was o5 mm. Zirconia was used as a non-stabilised form, i.e. the monoclinic form of zirconium oxide. It was introduced into the ceramic masses in two types which differed in particle size distribution, i.e. medium– large size (median grain size: d50 ¼100 mm) or fine-sized grains of ZrO2 (median grain size: d50 ¼10 mm). The chemical compositions and applied mark for the raw materials used here are presented in Table 1.

After the sintering process, different ceramic properties were measured which included:

 firing shrinkage according to diameter and height measurement  bulk density and open porosity according to standard PN-EN   

993-1:1998 (via the hydrostatic water immersion method using paraffin oil to avoid hydration) cold crushing strength according to standard PN-EN 993-5:2001 gas permeability according to standard PN-EN 993-4:1999 thermal shock resistance according to the procedure described below.

Each parameter was determined at least in three parallel samples.

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Table 3 Compositions of corrosive agents used in the corrosion test. Composition; wt% Cementitious corrosive agent Ordinary Portland cement 45 Chemical composition; wt% SiO2 Al2O3 9.12 2.27

Potassium sulphate 50 Fe2O3 1.11

CaO 29.05

MgO 0.43

Glassy corrosive agent Quartz sand Sodium carbonate 30 36

Sodium sulphate 20

Chemical composition; wt% SiO2 Al2O3 Fe2O3 33.04 2.66 0.05

MgO 2.39

CaO 3.38

Carbon 5

K2O 24.01

Na2O 0.11

SO3 26.70

Dolomite 10 K2O 0.09

Thermal shock resistance of the obtained ceramic samples was determined by the following method. The samples were heated at 950 °C for 25 min in an electric furnace and then rapidly immersed for 3 min in a container filled with room temperature water. After this cycle, the samples were put into the furnace again and the procedure was repeated. The test result is determined by the number of cycles the sample can withstand until a loss of 20 wt% of the original mass has taken place. 2.4. Corrosion of samples Corrosion studies were carried out on selected samples by the crucible test, for which two corrosive agents were prepared. Their composition was designed to simulate conditions in a cement kiln or glass-furnace regenerator; a detailed composition of these materials is given in Table 3. Dried components were weighed and then homogenised in a laboratory ball mill. The refractory ceramic samples for the corrosion test had a cylindrical form (ؼ h¼50 mm) with a cavity (ؼ20 mm; depth of about 20 mm) where the corrosive agent (
7 g) was placed and hand-pressed. The ceramic samples with the corrosive agent were heated to 1500 °C in a gas kiln with 10 h soaking time at the temperature above (for the cementitious agent), or heated to 1400 °C with 6 h

Na2O 25.24

L.O.I. 6.96

Total 99.76

Aluminium oxide 2 MnO 1.64

SO3 13.38

Manganum oxide 2 L.O.I. 18.10

Total 99.97

soaking time (for the glassy agent) and then cooled down together with the furnace. After cooling, the samples were cut along the height of the cylinder and the corrosion areas were determined by planimetric measurements. The percentage of the non-corroded surface of the samples was taken as the corrosion-resistance result. For detailed corrosion studies, the polished samples were observed by SEM equipped with an EDS chemical microanalyser. The chemical composition of the raw materials was determined by X-ray fluorescence spectroscopy (Panalitycal Magix PW2424 spectrometer). The microstructure of the obtained ceramic samples was observed by a scanning electron microscope (Mira III, Tescan) in combination with the EDS system (Oxford Instruments).

3. Results and discussion 3.1. Raw materials characteristics The starting magnesia clinker was characterised by high purity. The content of the magnesium oxide reached
98 wt% and was close to that declared by the manufacturer. The remaining part was principally represented by Ca, Si and Fe compounds. The magnesia clinker used here was characterised by a coarsely

Fig. 1. SEM image of a magnesia clinker fracture: general view (a); local area of isolated grains of periclase (b).

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Fig. 2. Electron mapping images of magnesia clinker: general image (a); magnesium (b), calcium (c), silicon (d), iron (d).

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crystalline structure, with the average size of the crystals reaching 100 mm (Fig. 1a). Individual grains of periclase were mostly directly connected to one another. Occasionally, they were connected through the intermediate phase (Fig. 1b), thereby forming local areas of isolated grains of periclase. The intermediate regions were dominated by Ca and Si compounds, which resulted from the chemical analysis and was confirmed by elemental mapping using SEM-EDS (Fig. 2). The iron compounds were homogeneously distributed in the clinker sample, probably as magnesioferrite inclusions in the periclase crystals. The magnesium oxide content in the raw material is important, but current technology requires the smallest amount of impurities which determine the properties of the finished products. One such index is the molar ratio of CaO to SiO2, which determines the presence of other phases in the system. Dicalcium silicate (C2S) has high refractoriness among the silicate phases that may appear. Other silicates affect the appearance of the liquid phase at a much lower temperature [2]. Based on a chemical analysis of the magnesia clinker used here, the calculated molar ratio of CaO/SiO2 was 1.86. This result indicates, in accordance with the Nadachowski classification [16], that the received magnesia clinker belonged to the third phase group of basic refractories and, in consequence, the presence of dicalcium silicate (C2S) as well as merwinite (C3MS2) should be expected in the firing material. The chemical analysis results for the zirconia materials used here were practically the same (Table 1). Material Zr1 was characterised by the presence of various-sized grains of asymmetric and irregular shape. These grains had a lamellar structure and locally formed a twinning. Inclusions of other phases were observed on the grain surfaces (Fig. 3a). Zirconia Zr2 was a fine powder (Fig. 3b). The SEM-EDS analysis of the inclusions contained in zirconia raw materials showed the presence of calcium zirconate (CaZrO3) as well as magnesium silicates, most likely forsterite (Mg2SiO4). 3.2. Properties of the obtained materials The results of all the physical and mechanical properties of the obtained composite refractories produced by incorporating

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various ratios and forms of zirconia (ZrO2) into the magnesia materials are presented in Table 4. These data are supplemented with graphs (Figs. 4 and 5) illustrating the influence of the amount and form of the zirconia materials on the properties of the ceramics. All of the obtained refractory materials from the MgO–ZrO2 system were characterised by negligible firing shrinkage. For the 8Zr2 sample, a maximum linear change occurred which reached 0.8% of the original size. For samples with a higher amount of fine zirconia (Zr2 series), the results showed slightly higher values of firing shrinkage in comparison to the addition of medium–large zirconia (Zr1 series). The external manifestation of this fact was the appearance of visible, slight cracks on the fired ceramics. There was no clearly visible relationship between the amount of introduced zirconia and the values of firing shrinkage. An interesting result was found for the 8Zr1 sample, i.e. there were no changes in the linear dimensions for this composite – shrinkage in height and diameter shrinkage were equal to zero. Presumably, this effect could be associated with the incomplete stabilisation of zirconium oxide by magnesia clinker. Under ambient pressure, zirconium oxide occurs in three major polymorphic forms. The monoclinic low temperature phase is stable up to
1200 °C. Above this temperature it converts to the tetragonal form and next to the cubic high temperature form (above 2300 °C) [17]. The conversion is reversible and is accompanied by an abrupt volume change due to differences in the densities of the respective forms of zirconium oxide. The monoclinic-to-tetragonal phase transformation is accompanied by a reduction in volume (an increase in density from 5.56 to 6.10 g cm  3 [2]) and, in consequence, intensive crack formation takes place during temperature changes. Additives of magnesium oxide, calcium oxide or yttrium oxide keep the high temperature phases metastable until room temperature is reached; they also reduce thermal expansion. Depending on the amount of additives (stabilisers), partially or completely stabilised zirconia can be achieved [1]. According to the MgO–ZrO2 phase diagram [18], zirconium oxide does not form any compounds with magnesium oxide but can reach complete stabilisation at a relatively low temperature

Fig. 3. SEM images of zirconia materials: Zr1 (a); Zr2 (b).

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Table 4 Properties of the obtained samples.

Bulk density / g·cm-3

0 2Zr1 4Zr1 6Zr1 8Zr1 2Zr2 4Zr2 6Zr2 8Zr2

Ø

h

0.3 0.3 0.3 0.2 0.0 0.5 0.6 0.6 0.8

0.1 0.2 0.2 0.0 0.0 0.1 0.2 0.2 0.4

Bulk density; g cm  3

Open porosity; %

Cold crushing strength; Gas permeability; nPm Thermal shock resistance; numMPa ber of cycles

2.80 2.84 2.86 2.87 2.89 2.81 2.87 2.90 2.92

16.8 16.3 16.2 16.2 16.4 16.8 16.4 15.8 15.7

48 72 70 66 64 59 54 53 49

3.3

18

3.2

17

3.1

16

3.0

15

2.9

14

2.8

13 Zr1

Open porosity / %

Sample designation Firing shrinkage; %

Zr2

2.7

12 0

2

4

6

8

Zirconia content / wt%

80

15

70

13

60

11

50 9 40 7 30 5

20

Gas permeability / npm

Cold crushing strength / MPa

Fig. 4. Bulk density and open porosity values as a function of zirconia additives.

3

10 Zr1

Zr2

0

1 0

2 4 6 Zirconia content / wt%

8

Fig. 5. Cold crushing strength and gas permeability values as a function of zirconia additives.

(
1300 °C). Fully stabilised zirconia is generally achieved with a dopant level of 15–16 mol% magnesia. Completely stabilised materials have a linear thermal expansion and the monoclinic form is

2.14 3.31 3.73 3.70 4.06 4.10 4.84 5.01 6.13

8 7 7 7 8 8 8 8 7

missing or only very small residual amounts are present. In the presented studies, the assumption was complete stabilisation of zirconia due to the significant excess of magnesia in the system during the firing cycle. On the other hand, due to the presence of Ca-compounds in magnesia clinker raw material (Table 1), the stabilisation of zirconia by calcium could also occur during the firing cycle. According to the CaO–ZrO2 phase diagram, complete stabilisation of zirconium oxide is achieved with a dopant level of 8 wt% calcium oxide [2,18]. Calcium zirconate (CaZrO3) is presented in the system. This is a phase with a very high melting point (
2340 °C) without polymorphic transformations [19]. With respect to the results obtained for the 8Zr1 sample, no shrinkage was observed due to the high probability of incomplete stabilisation of large grains of zirconium oxide. As the material cooled down, phase transformation of the unstabilised part of the zirconia grain's core occurred which locally generated an increase in volume and, in consequence, compensation of sintering shrinkage was present. For series containing Zr2 zirconia, a systematic increase of firing shrinkage was observed due to the better sintering of fine powder and practically complete ZrO2 stabilisation during the firing cycle. The SEM observation confirmed the above. In comparison to the Zr2 series (Fig. 7), a clearly visible difference in the zirconia grains was observed for material containing medium–large grains of zirconia. For the Zr1 series (Fig. 6), unstabilised core grains were observed around which there were radial cracks as a result of stresses caused by phase transformation of zirconium oxide and their various thermal expansion. Chemical analysis in the microareas by SEM-EDS analysis from the core to the edge of the grains showed the presence of pure zirconia and two different calcium zirconates, respectively. The latter were created during the reaction of zircon oxide with Ca-compounds, as presented in the magnesia clinker used as a secondary phase. Microanalysis by the EDS system in the calcium zirconate micro-region showed that the outer layer was rich in calcium as compared to the rest of the grain, while the inner layer contained a larger amount of zirconium with respect to the outer layer. EDS studies for the inner layer showed the following composition (in atomic %): Ca ¼
6%, Zr ¼
30% and O¼
60%. This composition indicates the presence of ZrO2 and the compound with formula CaZr4O9. The last compound is an intermediate phase named type φ1 and is described by the general formula: (Me,Zr)5O9. This phase is specific for materials stabilised by calcium oxide which is formed during thermal treatment and is derived from regular solid solution [20]. In turn, the outer layer of grain was characterised by a composition with the Ca:Zr:O atomic ratio equal to about 1:1:3, thus indicating the presence of calcium zirconate CaZrO3. Due to the marked differences in density between magnesium and zirconium oxides (3.58 vs. 5.56 g cm  3, respectively [1]), increasing the amount of zirconia caused a systematic rise in the

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Fig. 6. (continued)

Fig. 6. SEM images of Zr grains in obtained 2Zr1 refractory materials and EDS analysis suggesting the presence of: ZrO2 (point 1); CaZr4O9 (point 2); CaZrO3 (point 3); MgO (point 4).

bulk density of the obtained materials (Fig. 4). The highest values of bulk density were obtained for ceramics with 8 wt% of ZrO2, i.e. samples 8Zr1 and 8Zr2. The addition of the zirconium compound also caused a decrease in open porosity, which was more clearly visible for samples with fine-grained zirconia (Zr2 series). This

provided better packing of the grains in the product matrix. The influence of zirconia on magnesia refractories becomes noticeable when cold crushing strength and gas permeability are considered (Fig. 5). The addition of 2 wt% of zirconium oxide in each series caused a significant increase in the crushing strength values as compared to the reference material without any addition. For the Zr1 series it increased from 48 to 72 MPa, while for the Zr2 series it increased to 59 MPa. These values were the highest. A further increase of ZrO2 in the masses reduced the value of cold crushing strength, but it is worth noting that for each of the obtained samples this parameter was still favourable in relation to material without any addition. When comparing the influence of the introduced zirconia form, more favourable properties were obtained for medium–large zirconia grains (Zr1 series). This phenomenon may be related to the strengthening mechanism of the material via the martensitic transformation in zirconia grains, which is initiated by a propagating crack. This mechanism leads to the formation of a compressive stress field around the crack which inhibits its propagation, thus consequently the cold crushing strength of the ceramics increases [21,22]. Along with an increase in the amount of introduced zirconium oxide, the gas permeability values of the obtained refractory samples systematically increased for both series. These were in a range of from 2.14 nPm for the reference sample to 4.06 nPm (8Zr1 sample) or 6.13 nPm for the 8Zr2 sample (Table 4; Fig. 5). Gas permeability reflects the material's characteristics which directly affect the ease of gas flow through the material, especially pore distribution and, in consequence, direct influence on corrosion resistance. The increase in gas permeability values in combination with more addition of zirconia could be attributed to the formation of a heightened amount of micro-cracks in the ceramics. These result from differences in the thermal expansion coefficient

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Fig. 8. Macroscopic image of samples after the corrosion test.

20 12 entitious corrosive e agent ceme sive agent glasssy corros

80 8

60 6

8Zr1Ng

4Zr1Ng

g 2Zr1Ng

0N 0Ng

8Zr1Ng

g 0Ng

2 20

4Zr1Ng

4 40

2Zr1Ng

Corrosion resistance / %

00 10

0 Sample Fig. 9. Relationship between corrosion resistance and type of obtained refractory sample.

Fig. 7. SEM images of Zr grains in obtained 2Zr2 refractory materials and EDS analysis suggesting the presence of: CaZr4O9 (point 1); CaZrO3 (point 2); MgO (point 3).

of the main components of ceramics (i.e. periclase and zirconia). Consequently, the higher the content of zirconium oxide, the higher the structure relaxation of the ceramics may occur. This is in accordance with Ceylantekin and Aksel's studies [14,23], where the authors showed that the introduction of a higher amount of zirconium oxide caused the generation of too many micro-cracks, which could connect to one another. The addition of zirconium oxide did not significantly affect the materials' resistance to sudden temperature changes. Generally, all of the obtained materials were characterised by the same level of thermal shock resistance (Table 4). Level 7–8 of the water cycles was reached for each obtained ceramic sample. It should be emphasised that the obtained results were very good. As a rule,

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Fig. 10. SEM images of reference sample after corrosion test by glassy agent – general view (a); selected microareas at the higher magnification (b and c) with accompanying EDS point analysis in at% (d).

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Fig. 11. SEM images of zirconium-containing sample after corrosion test by glassy agent – general view (a); selected microareas at the higher magnification (b and c) with accompanying EDS point analysis in at% (d).

magnesia refractory products are not resistant to sudden thermal shock (both air and water cycles). It is widely accepted that these products are resistant to only a few air cycles.

3.3. Corrosion study Due to the more favourable results for Zr1 addition (i.e. lower shrinkage and gas permeability as well as higher cold crushing

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Fig. 12. SEM images of reference sample after corrosion test by cementitious agent – general view (a); selected microareas at the higher magnification (b and c) with accompanying EDS point analysis in at% (d).

strength of the samples), as was previously discussed, corrosion studies were conducted on selected obtained refractory ceramics. Materials designated as 0, 2Zr1, 4Zr1 and 8Zr1 were chosen for the test samples. The macroscopic image of selected samples obtained after cutting marked corrosion areas is presented in Fig. 8. It is clearly evident that increasing the amount of zirconia in the

ceramic masses caused an increase in the material's corrosion resistance against a chemical attack by the corrosive agents used in this study. A comparison of the obtained corrosion resistance results as a percentage of non-corroded surfaces of the samples is shown in Fig. 9. Due to the basic nature of the cementitious corrosive agent,

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Fig. 13. SEM images of zirconium-containing sample after corrosion test by cementitious agent – general view (a); selected microareas at the higher magnification (b and c) with accompanying EDS point analysis in at% (d).

corrosion resistance of the reference sample (0) was already high. For the second corrosive agent (glassy), this agent's penetration of the reference sample was significantly higher and was due to the conditions of the experiment procedure. During soaking at the

assumed temperature (1400 °C for 6 h), a high amount of liquid phase with probably low viscosity was created from the glassy corrosive agent and, in consequence, it could easily infiltrate the porous ceramic material.

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For both of the corrosive agents used here, the zirconia addition improved corrosion resistance. The greatest amelioration was observed with 2 wt% zirconia addition. For the corrosion test with a cement-like agent, resistance increased at a level of ten percentage points (from
85% to 96%), while for the glassy agent it ranged as high as ca. twenty percentage points (from
40% to
60%). A further increase in the zirconia content in the ceramic samples did not significantly affect the obtained results. For the cementitious corrosive agent, corrosion resistance slightly increased (reaching
98% for the 8Zr1 sample), while for the glassy agent a systematic decrease was found (
55% for the 8Zr1 sample). The reason for this fact may have been too much structure relaxation of the material with the higher amount of zirconia content, which was discussed previously with respect to mechanical strength. The formation of significant amounts of micro-cracks in the ceramics contributed to easier penetration by the melted glassy corrosive agent and, in consequence, corrosion resistance decreased. The SEM observations were complemented by the abovementioned results. The areas of intensive penetration by the glassy corrosive medium were large. For the sample without zirconia addition after the corrosion test (Fig. 10), an area with high porosity (in a range of approximately 4 mm from the surface) was observed. This could indicate a chemical reaction leading to leaching or evaporation of part of the material. There were also numerous cracks passing by both the matrix and the grain, which could indicate the mechanical stress that occurred in the material. Despite the lack of sharp boundary zones of the analysed material, the chemical composition near the surface and the depth of the material subjected to the glassy corrosive showed differences in the chemical composition of the phases formed in the examined areas. From the corrosion side, besides the periclase grains, phases consisting of Na, Mg, Si, Ca and O with varying proportions of components were identified. These were probably in the form of various silicates. In turn, in the depth of the material only traces of sodium were observed. The microstructure of the zirconium-containing samples after glassy corrosion (Fig. 11) was also characterised by cracks extending through the material, the net of the cracks was slightly thicker in the area of the corrosive agent's activity. In the area close to the surface, besides MgO and ZrO2, the formation of Mgsilicates as well as Ca- and Mg-silicates or Na- and Mg-silicates was identified. In the depths of the material, additionally calcium zirconate appeared. It was created primarily due to the presence of the Ca compound in the magnesia raw material used here. In the case of the corrosion test by a cementitious agent, the microstructure of the obtained refractories was different. For the reference sample after corrosion (Fig. 12), numerous cracks extending mainly by the matrix were also observed. The region of intensive corrosion reached approximately 11 mm into the material. In the area close to the sample surface, besides the grains of MgO, dicalcium silicate and calcium aluminate in varying ratios of components (partially doped with Fe) were identified. In the deeper parts of the material, only the formation of dicalcium silicate was found and traces of Al and Fe were identified. For the zirconium-containing samples after cementitious corrosion (Fig. 13), the region of intense corrosion was markedly lower and reached approximately 7 mm. Cracks appeared mainly by the matrix of the material. Grains of magnesia were present near the surface of the corrosion material but there were no zirconia grains. Instead, isometric grains of calcium zirconate were present that had appeared in the presence of significant amounts of calcium compounds originating from Portland cement, which is a major part of the cementitious corrosion medium. Furthermore, calcium and iron aluminate (with a possible dopant by Mg) were identified as well as calcium silicates with different molar ratios of

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ingredients. In the depths of the material, zirconium oxide doped by Ca appeared (possible solid solutions of CaO–ZrO2 or metastable phase), as type φ1 (CaZr4O9).

4. Conclusions The results obtained in this study showed the influence of zirconia addition with different grain size on the final properties of modified magnesia refractory products. The overriding goal of this study was to investigate the possibility of using zirconium oxide as an additive in ceramic masses to obtain modified magnesia refractory products that could be characterised by high resistance to thermal shock and showing high corrosion resistance. Consequently, these materials could replace the magnesia chromite refractory ceramics which are still being used today. Based on the obtained results, it could be concluded that:

 zirconium oxide addition in the amount of 2 wt% positively af-

  

fected the primary refractory ceramics’ properties – in comparison to the reference material, an increase in cold crushing strength and decrease in open porosity were observed; 2 wt% addition of ZrO2 significantly improved the corrosion resistance of ceramics against a chemical attack by a cementlike and glassy-like corrosion medium; a further increase in the amount of zirconium oxide in the ceramics contributed to a gradual decrease in cold crushing strength and increase in gas permeability; particle size of the introduced zirconia affected the obtained results – generally, better properties of MgO–ZrO2 refractories were achieved for medium–large zirconia.

Considering the high cost of zirconium-containing raw materials and the effect of zirconia addition on the properties of MgO– ZrO2 refractory materials, it can be concluded that 2 wt% addition of zirconia to magnesia refractory ceramics appears to be an optimal solution. The good corrosion behaviour of the obtained magnesia-zirconia materials supports their potential use as an alternative to magnesia-chrome bricks for different applications, such as for the highest part of a glass-furnace regenerator or for the burning zone of cement kilns.

Acknowledgements This work is the result of statutory activity at the Institute of Ceramics and Building Materials, Refractory Materials Division in Gliwice (no. 2N019S15) in 2015, supported by the Polish Ministry of Science and Higher Education.

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