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THE STUDY OF MOLTEN LIQUID - REFRACTORY INTERACTIONS – IT IS ALL ABOUT THE PHASE(S)

Andrie Garbers-Craig1

1Centre for Pyrometallurgy, Department of Materials Science and Metallurgical Engineering, University of Pretoria, Lynnwood Road, Hatfield, Pretoria, 0002, South Africa

Keywords: Chemical interaction; Wear mechanisms; Slag; Metal; Matte; Refractory material

Abstract

The study of chemical interactions between slag / metal / matte and refractory materials has been an ongoing theme from the time when the first metals and alloys were produced. Over the years the refractory industry has seen extraordinary technological progress such as the development of the magnesia-carbon bricks and self-flowing castables. The advances that take place in high temperature processes, specifically the iron and steel industry, drive innovations in the refractories industry. Since these pyrometallurgical processes are dynamic, continuously improving and changing to process inter alia lower grade ores in energy more efficient and environmentally more friendly ways, the refractories industry has to keep current. This paper discusses the important role that the study of molten liquid – refractory interactions plays in understanding wear mechanisms and directing the development of refractory materials. It also gives relevant examples from refractories used in the production of different commodities.

Introduction

The refractory materials industry has been called the ‘hidden industry’ [1] and the ‘enabler of civilization’ [2]. The origin of this industry can be traced back to the ability of humans to produce and contain fire [1]. During preindustrial times, furnaces were small, reaction temperatures low (1100 to 1250ºC) and times at maximum temperature short [3]. Metallurgical processes were such that the development of exotic refractory materials was not necessary. Innovation in refractory materials therefore only started with changes and developments in pyrometallurgical practices, which started to occur in the 15th century [3]. Today the refractories industry is highly sophisticated, with 38.6 Mt of refractory materials produced in 2013 [4]. In 2013, 73% of the produced refractory materials were used by the iron and steel industry, followed by the cement and lime industry (13%) and the non-ferrous metals industry (4.5%) (Figure 1) [4]. Developments in iron and steel production technology remain one of the main drivers of innovation in the refractories industry, and are likely to continue to do so for the foreseeable future. Nearly 95 years ago Francis Pyne listed the requirements for refractories used in the copper industry as having “(1) resistance to temperatures, (2) resistance to chemical action, (3) the ability to withstand sudden temperature changes, (4) minimum absorption of slag and metal, and (5) the absence of manufacturing defects” [5]. These requirements still hold today, but with the added demands of higher productivity via longer campaign lives, emphasis on vessel

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Advances in Molten Slags, Fluxes, and Salts: Proceedings of The 10th International Conference on Molten Slags, Fluxes and Salts (MOLTEN16)

Edited by: Ramana G. Reddy, Pinakin Chaubal, P. Chris Pistorius, and Uday PalTMS (The Minerals, Metals & Materials Society), 2016

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Figure 1. Global refractory market by end-user industries, 2013 [4]

availability, more efficient use of energy, reduction in greenhouse emissions and a ‘greener’ approach through recycling of refractories and sustainable use of materials [2, 6]. In addition, raw materials suppliers and refractory companies are required to adapt to a long-term decline in commodity prices as illustrated in Figure 2 [7].

Figure 2. Trend in ‘raw industrials’ prices, 1800 - 2000 [7] The refractories industry has developed from initially following a trial-and-error approach to an industry that innovates through a fundamental understanding of how the phase composition and microstructures of these materials are interrelated with their properties, performances and processing methods. Technological developments have been driven by studies on how chemical and physical interactions between slag / metal / matte and refractory materials impact on wear mechanisms, and also on how the refractory materials impact on product quality in the industries where they are used. This paper describes the fundamental theory on and interpretation of refractory wear through penetration and corrosion by molten liquids. It also includes a

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discussion on how investigations on charge – refractory interactions have guided refractory choices and developments in the steel and platinum group metal (PGM) industries.

The study of molten liquid - refractory interactions

Fundamentals Excellent explanations of how thermodynamic and kinetic factors impact on wear through liquid penetration, corrosion, erosion and thermal shock are available in the literature [8-12]. Equations described below are used in the design of refractory materials and slag compositions to limit refractory wear. Penetration can be purely physical, which is then a capillary force driven flow of liquid into the refractory material, or it can be a chemical penetration, which includes chemical interaction that can lead to densification and chemical or structural spalling of the material [8]. The rate of physical penetration of a liquid into a refractory material is conventionally described by Poiseulle’s law, which is dependent on the radius r of the capillary (open pore or microcrack), the capillary suction pressure ΔP, the dynamic viscosity of the liquid η, the penetration depth l of the liquid into the refractory and the time t [10]: dl/dt = r2ΔP/(8ηl) (1) The ΔP term in equation (1) can be expressed in terms of the properties of the penetrating liquid, where γ is the surface tension of the liquid and θ is the wetting or contact angle: ΔP = 2γcosθ/r (2) Substitution of Equation (2) into Equation (1), followed by integration, then gives: l2 = [r γcosθ/(2η)] t (3) Decreasing the capillary radius and surface tension, and increasing the contact angle and viscosity of the liquid can therefore reduce penetration of a liquid into a refractory material. Corrosion is the chemical attack of refractory, mostly by molten liquid (slag, metal and matte), but also by gas. When liquid chemically reacts with the refractory material the reaction product(s) can directly dissolve into the liquid (direct dissolution) or a solid reaction product can form at the refractory – liquid interface, which is then dissolved into the liquid (indirect dissolution) [10]. In ‘direct dissolution’, when the diffusivity of the reaction product(s) is faster than the rate of the chemical reaction at the interface, the initial dissolution rate can be described by: J = K(Ac/Ao)Cm (4) where J is the dissolution rate, K the first order rate constant, Ac the actual area of refractory, Ao the apparent area of the refractory and Cm the concentration of the reactant in the melt. In the case of ‘indirect dissolution’, the corrosion rate can be expressed by the Nernst equation [10]:

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J = D(Cs - Cm)/δ (5) where D is the diffusion coefficient, Cs is the concentration of the component at which it is saturated in the liquid, Cm is the concentration of the component in the liquid and δ is the effective boundary layer thickness. Corrosion is also predicted and explained through thermodynamic considerations. These include assessing phase stabilities before melting (e.g. quartz, zircon, andalusite), predicting possible chemical reactions that can take place between the charge and refractory material (whereby a reaction product interlayer could form at the hot face), the onset of melt formation in the refractory, compatibility between the slag / metal / matte and the main refractory component (oxide / carbon based), reactions between different types of refractory materials, the concept of ‘local thermodynamic equibrium’ (where the penetrating and reacting liquid changes its composition as it reacts deeper into the refractory) and oxidation-reduction reactions [8, 12]. Post mortem and laboratory scale analyses on refractory materials Analysis of wear mechanisms associated with refractory materials usually include post mortem analyses of refractory samples removed from production as well as samples from laboratory scale test work. Laboratory scale tests typically include static tests (the dipping / finger test; slag pot test; induction furnace slag test) and dynamic tests (slag drip test; rotating finger test; rotary kiln slag test). Refractory samples are then analysed using mostly chemical analysis (XRF or ICP), XRD, reflected light microscopy and SEM-EDS analyses. Wear mechanisms are then examined and explained through the use of phase diagrams and thermodynamic calculations, using software such as FactSage and Thermocalc. It should however always be kept in mind that the observed phase relations at room temperature reflect the thermal history of the sample, and do not reveal phase relations at operating temperatures. Phase relations at room temperature should therefore always be extrapolated back to operating temperatures before corrosion mechanisms are interpreted [13]. More recent developments include in-situ analysis methods in which phase relations and corrosion mechanisms are related to dimensional changes of the refractory materials as well as their thermomechanical properties.

Refractory Development in the Steel and PGM Industries Refractory evolution in the steel and PGM industries, which was guided through studies on charge – refractory interactions, is discussed in this section. Steelmaking: In 1957 silica brick was still the number one steel plant refractory, even though it was meeting severe competition from basic refractories, which have far higher melting points [14]. The main reason for the successful use of silica brick was the high temperatures at which it could be used in contact with iron oxide (close to the liquidus temperature in the two-liquid region, which is in excess of 1600ºC, Figure 3) and its high load-bearing capacity at these high temperatures [15, 16]. However, silica brick had the drawback of cracking and spalling when

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they were rapidly heated or cooled over low temperature ranges [1], due to dimensional changes associated with the conversion of one polymorphic form of SiO2 to the other.

Figure 3. Pseudo binary iron oxide – SiO2 system, indicating the variation in size and liquids temperature of the miscibility gap with changing Po2 [15]

The rise of the LD (Linz-Donawitz) or BOF (Basic Oxygen Furnace) process resulted in the replacement of silica bricks by basic refractories [14]. The success of one such basic refractory, magnesia bricks, lies in its high melting point (2800°C) as well as its excellent resistance to attack by iron oxide, whereby the (Mg,Fe)O, (Mg,Fe)2SiO4 solid solution phases form under reducing conditions (Figure 4a) [15]. Its principal limitation is however, its high thermal expansion (Figure 4b), which makes the production of MgO bricks with high thermal shock resistance difficult.

(a) 1650°C isothermal section of the MgO-FeO-SiO2 system in contact with metallic iron [15]

(b) Percentage linear expansion of different types of refractory materials as a function of temperature [17]

Figure 4. Benefits and limitations of magnesia bricks

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The earliest reference to chemically bonded basic refractory was made in 1905, referring to sodium silicate and calcium chloride as binders, while direct-bonded basic brick with low impurity content has been manufactured since the late 1950s [18]. During the early 1960s it became clear that magnesia bricks of higher purity and density were required, as impurities (which have lower melting points than MgO) cause a loss in strength of the material and provide pathways along which liquid penetration and slag attack can take place [19]. One such impurity is B2O3, as it combines with MgO (and other impurities such as CaO) to form a low melting liquid, which easily wets the MgO grains and prevents direct bonding (Figure 5a [20]) [19]. Since the mid-1980s the trend has been to obtain the highest MgO content possible when manufacturing MgO brick. The importance of the CaO/SiO2 impurity ratio in MgO bricks was also understood, as this ratio determines the temperature of initial melt formation in the brick (Figure 5b) [19].

(a) The MgO – B2O3 phase diagram [20] (b) The MgO-SiO2-CaO phase diagram [15].

Figure 5. Influence of impurities on the refractoriness of MgO-based materials

The next step was to further densify magnesia bricks in order to improve their corrosion resistance. Penetration and corrosion resistance of a MgO refractory material can be improved by reducing its porosity, i.e. increasing its density [21]. However, when the porosity becomes too low, the thermal shock resistance of the refractory deteriorates. The design of a MgO refractory material that has a high thermal shock resistance as well as excellent corrosion resistance was therefore a matter of a delicate balance. Adding carbon in the form of pitch addressed this problem to some extent when high-strength pitch-impregnated burned magnesia bricks were developed for the high wear areas of BOF linings in the late 1960s, with MgO contents in the order of 95% [19]. Since carbon has low thermal expansion and high thermal conductivity, a material was produced with improved thermal shock resistance as well as corrosion resistance. However, since the steel industry moved from ingot to continuous casting by the late 1970s, the maximum temperature in the converter increased from 1650 to 1700°C, which necessitated the production of higher quality refractories [22]. Between 1975 and 1980 resin-bonded magnesia-carbon refractories were introduced in Japan, first for electric arc furnace hot spots and shortly

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thereafter for BOFs [19, 22]. Shortly thereafter graphite became the principal carbon source. Since high carbon contents in the MgO brick introduced the severe disadvantage of oxidation and subsequent increase in porosity of the brick, the addition of antioxidants to oxide-carbon bricks followed. These bricks were the first truly composite refractory materials with a ‘self-repairing’ function when coming in contact with oxygen. The ‘self-repairing’ function includes a decrease in carbon loss by reduction of CO(g) to C(s) with an accompanying decrease in porosity, crystallization of amorphous carbon from the binder and the formation of a protective surface layer (Figure 6) [21]. Magnesia-carbon bricks are nowadays successfully used due to their excellent slag corrosion resistance and thermal shock resistance. The benefits of MgO-C bricks include low thermal expansion and high thermal conductivity (i.e. good thermal shock resistance), improved penetration resistance (due to the non-wettability of graphite by slag; the generation of CO and/or Mg vapour inside the brick which resists liquid infiltration; and the formation of a dense MgO layer at the slag/brick interface, physically preventing slag infiltration), and the reduction of iron oxide in the slag to Fe by carbon, thereby avoiding attack of the MgO by iron oxide [21, 23, 24]. However, the high carbon content in the oxide-carbon brick implied higher release of carbon dioxide during tempering and preheating, which increased pollution and carbon footprint of the brick, together with high energy losses [25]. The challenge that then followed was the reduction of the fixed carbon content of the MgO-C brick without decreasing its thermal conductivity or downgrading its corrosion resistance. Research has shown that the addition of high surface area carbon sources (such as nano-carbon) to the matrix of the brick, can reduce the fixed carbon content of the brick without decreasing its corrosion resistance and decreasing its thermal conductivity [25]. The success of oxide-carbon bricks has initiated research into the development of castables that contain graphite. However, incorporating graphite into castables is technically very difficult as graphite has a low aqueous wettability and is therefore difficult to disperse, while aluminium-based antioxidants has a tendency to hydrate [26]. Current research includes investigations into coating of the graphite and antioxidants.

Figure 6. Illustration of the ‘self-repairing’ function of oxide-carbon materials [21]

It is also important to touch on the developments that took place in the manufacture of castable

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materials, from when they where first introduced into the refractories market in the early 1900s with the industrial production of calcium alumina cement (CAC) [27]. Although the Lafarge Company in France started the commercial production of calcium alumina cement in 1918, it was only in the 1950s that high purity CACs, developed specifically for the refractories industry, came on the market [28]. By 1960 conventional castables based on high purity CAC (up to 30 wt%) and tabular alumina were in common use. Disadvantages associated with the conventional castables proved to be low strength (due to their high water content required for placement, which increased porosity), loss in strength during dehydration (in the 538 - 982°C temperature range, as the hydraulic bond was broken down and sintering was still too sluggish to allow the development of a ceramic bond) and deteriorated high temperature properties (as high concentrations of CaO favours the formation of CaO – Al2O3 – SiO2 - based liquid phases on firing, Figure 7) [28]. By the end of 1960 it was clear that the cement content in castables should be reduced, i.e. to reduce the water content for placement but maintain the strength [28]. This lead to the development of low (LCC) and ultra-low (ULCC) cement castables whereby flow was achieved through more efficient particle packing, together with submicron size matrix additions and the use of dispersants. These castables consist of the aggregate and a binding system, where the aggregate phases have pronounced effects on the microstructural evolution on heating, and therefore also the corrosion mechanisms in these materials. In the development of the LCCs and ULCCs the composition of the material was moved from the SiO2-CaO-Al2O3 – based system of the CC to either the SiO2-Al2O3 or the Al2O3-CaO systems [29]. In the SiO2-Al2O3 system, when CaO concentrations are very low, the high melting point mullite phase is readily formed, while the Al2O3-CaO system aims to form the high temperature bonding phase CA6 (Figure 7).

Figure 7. SiO2-CaO-Al2O3 phase diagram, which describes the compositions of CCs (which contains the low melting phases CAS2 and C2AS) and the LCC and ULCCs, which contain either A3S2 or CA6 [15]

The concern with steel cleanliness and efficiency of steel desulphurization in secondary steelmaking has sparked the development of castables based on the alumina-spinel and alumina-

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magnesia systems [28, 30]. Refractories with preformed MA spinel were developed in Japan in the late 1980s, while in the 1990s in-situ spinel-forming compositions were developed [31]. The advantage of using spinel-containing alumina-based castables in contact with steelmaking slag, is that the MA spinel traps FeO and MnO from the steelmaking slag in the spinel lattice. The further infiltration of the slag is then inhibited as its liquidus temperature and viscosity are increased [32]. The pre-formed spinel alumina-based castables have in general high volumetric stability and improved chemical resistance to steelmaking slag, compared to alumina castables, while in-situ spinel forming castables have a significant lower volumetric stability, but higher corrosion resistance when compare to added MA-containing castables. The in-situ spinel castables have however the added challenge of magnesia hydration [32]. Resin-bonded AMC (alumina-magnesia-carbon) bricks for steel ladles combine the advantages of the alumina and spinel phases with those of carbon-containing refractories [33]. Through time the refractory castables have undergone major developments as a result of available raw materials of improved quality, optimisation of particle size distributions and packing design (to achieve higher density materials) and control of particle dispersion and therefore flowability [31]. Current trends in steelmaking include a focus on high quality steel grades with very low limits of residual elements and steel cleanliness, as required by end users [34, 35]. This implies that extended processing of liquid steel in secondary metallurgy can take place and increasingly complex production routes can be followed to achieve metallurgical targets. More emphasis is also placed on reduced energy consumption and increasing resource efficiency [35]. In order to achieve these aims the refractory materials that are used in specifically secondary steelmaking must continuously be re-evaluated and improved as they directly impact on steel quality and cleanliness. The steel ladle has become a “metallurgical reactor” where refractory requirements include corrosion resistance against metallurgical reactive slag, thermodynamic stability of refractory oxides to avoid re-oxidation of steel, and lower carbon content to avoid or limit C-pick up [35]. Higher purity refractories are therefore required for clean steel making. Synthetic materials such as tabular alumina or corundum have replaced natural ones such as andalusite (which has a high SiO2 content and therefore has a low lining life against basic slags) and bauxite (which also contains significant amounts of silica and impurities such as TiO2 and Fe2O3) in high alumina refractories. Significant developments have also taken place regarding energy efficiency, with for example the development of new insulating alternatives. PGM (platinum group metal) smelting: The PGM smelting process involves the smelting of spray dried flotation concentrate to produce a Fe-Ni-Cu-based furnace matte. The matte is then converted to a Ni-Cu-based matte, which contains the PGMs. This converter matte is then treated in the base metals refinery to extract the Cu and Ni, upon which the leach residue is sent to the precious metals refinery for final separation of the PGMs. PGM-bearing nickel-copper concentrate smelted in South Africa are mostly derived from the Merensky Reef (which has high concentrations of base metals and sulphur) and the UG2 Reef (which is rich in chromite and contains low quantities of base metal sulphides). PGM smelting has some distinct challenges as compared to copper and nickel sulphide smelting (Figure 8) [36]: Typical operating temperatures of primary copper smelting furnaces are in the 1180 – 1380°C temperature range, with matte liquidus temperatures in the range of 940 to 1125°C, while in the nickel sulphide industry slag operating temperatures vary between 1240 and 1400°C, with matte

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temperatures ranging from 1150 to 1360°C [37]. Average PGM slag liquidus temperatures in primary furnaces range from 1460 to 1700°C [37], with exceptionally high matte superheats of up to 690°C [36]. The two main components in the primary PGM slags that are responsible for these high liquidus temperatures are MgO and Cr2O3. High concentrations of Cr2O3 in the slag also increase the potential for precipitation of spinel at the matte-slag interface or as build-up in the hearth [38]. Dwindling supplies of high base metals and sulphur containing Merensky concentrate have forced PGM producers to mainly use UG2 concentrate [39, 40]. The main challenge in smelting PGM concentrates from the Bushveld Complex, South Africa, is therefore the increasing chromium content of the charge [39]. Solutions to this problem include the dilution of the high chromium-containing concentrate with Mogalakwena concentrate, which originate from the Platreef ore (which is metallurgical similar to the Merensky ore [40]) and operation at deep electrode immersion and high hearth power densities [38, 41, 42].

Figure 8. Comparison of operating temperatures and matte superheats across the copper, nickel, and PGM smelting industries [36]

Electric smelting of platinum concentrate started in 1969 at the Rustenburg Platinum Mines Ltd, in Rustenburg, South Africa with the commissioning of a 19.5 MVA Elkem rectangular electric furnace [43]. The sidewalls were externally water cooled, while pelletised concentrate was fed to the furnace. The furnace (27.2 m long, 8.0 m wide, 6.0 m high) was lined with magnesia bricks, while superduty fireclay bricks were used in the wall above the slag level and in the roof. Lonmin commissioned its first furnace in 1971, which was a 12.5 MW Merensky six-in-line furnace, followed by their first two 2.3 MW circular furnaces in 1982 to smelt UG2 concentrate [42]. High purity, direct bonded magnesia and magnesia-chrome bricks are typically used in the primary PGM smelter and the converter. The primary smelter slag is SiO2 – MgO – FeO – based, but also contains some CaO, Al2O3 and Cr2O3, while a fayalitic-type slag, which contains some nickel and copper oxides, forms in the converter. The presence of high levels of FeO, and nickel and copper oxides in the converter, necessitate the presence of MgO, since MgO-based refractory can absorb substantial amounts of these oxides, without losing refractoriness. Mag-chrome and chrome-mag bricks are in general more resistant to fayalitic slags than magnesia bricks, while the chromite grains are in turn more resistant than the periclase.

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Currently the largest capacity six-in-line submerged-arc furnace is the Polokwane smelter of Anglo American, which is rated at 68MW (168 MVA) and designed to treat 650 000t/a of concentrate [44]. Furnaces with such high hearth power densities, with a concentrate layer that is heat-insulating, have high sidewall heat fluxes. This necessitated the practice of using furnace sidewall cooling whereby a slag freeze lining is formed [38]. However, the use of copper coolers brings its own unique problems: Sulphur vapour, sulphur dioxide, water vapour and chlorine gas are given off during smelting [45, 46]. Sulphur vapour originates from sulphides and labile sulphur contained in the feed [45], chlorine gas from the dissociation of halides such as KClMg(SO4) and KMgCl3 which can potentially be present in the feed [46], and water vapour from the moisture and chemically bound water in the feed material [47, 48]. Water vapour dissociates at high temperatures, which can then also lead to the formation hydrogen sulphide and hydrogen chloride gas [49]. High temperatures associated with the smelting process drive these gases towards a cooler place, which is in front of the copper coolers [46]. This then leads to corrosion and possible catastrophic failure of the copper coolers [44, 50]. This mechanism was confirmed by a post mortem study on the slag freeze lining and magnesia-chrome bricks removed from the concentrate-slag interface and bottom section of the slag layer. The slag freeze lining was highly porous, thereby allowing base metal sulphides and sulphur and chlorine containing gases to move through the freeze lining to the copper cooler where it sulphidised the copper, forming a non-adherent copper sulphide layer (Figure 9) [51, 52]. This corrosion mechanism has been named ‘chloride accelerated sulphidation’. Great benefits have subsequently been obtained from replacing the magnesia-chrome bricks with graphite blocks in current smelter designs. A post-mortem analysis of a graphite block, which was removed from the same PGM smelter, confirmed that graphite plays a significant role in reducing the extent of penetration and reaction of corrosive gases and base metal sulphides with the copper cooler [51, 53].

Figure 9. Freeze lining from a primary PGM smelter: (a, b) Side views; (c) Front view [52]

Magnesia-chrome and alumina-chrome bricks have been used in the matte tap hole – another area of high wear. Thermodynamic modeling [54] predicted that as the matte temperature increases the matte – tap-hole brick wear mechanism changes from solely matte penetration to penetration with accompanying chemical reaction with the brick. Laboratory-scale experiments at 1500 - 1700°C confirmed that at matte temperatures above 1500°C wear of the magnesia-chrome refractory bricks are due to extensive matte penetration, with subsequent chemical interaction between the matte and brick that caused chromium pick-up in the matte [55]. Wear of alumina-chrome brick was also characterized by extensive matte penetration. Chemical

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interaction between the matte and brick occurred at temperatures above 1550°C, forming the FeCr2O4 spinel phase at the matte – brick interface [56].

Conclusions Very little was known about raw materials and how their phase compositions impacted on brick-making and brick quality in the early years of refractory material production. Harbison wrote in 1866: ‘…the manufacture of brick up to this time had been purely chance work…. frequently whole kilns of brick would be…found so defective that it was unsafe to put them on the market’ [1]. Today, 150 years after Harbison’s ‘confession’, the refractories industry can pride itself on being a mature and innovative industry, with the knowledge and expertise to engineer and manufacture microstructures for specific applications with a specific set of target properties. Challenges in R&D include the development of standard in-situ refractory test methods, whereby refractory performance in service can more accurately be predicted. Exciting current research topics also include the development of bendable, flexible and self-healing refractories. These developments are all dependent on comprehensive studies on liquid – refractory interactions. - It is indeed all about the phase!

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