CLASSIFICATION OF LIGATURES AND METHODS OF THEIR PRODUCTION

2.1. Requirements for ligatures

In foundry production, alloys occupy a significant share in the volume of charge materials: depending on the chemical composition, up to 50% of alloys. A master alloy is an intermediate alloy containing a sufficiently large amount of alloying metal that is added to the melt to obtain the required chemical composition, structural and technological properties of castings and ingots. As a rule, alloys for aluminum and magnesium alloys contain only one alloying component, but sometimes triple and quadruple alloys are prepared. The composition of complex alloys is selected in such a way as to ensure that the desired chemical composition of the alloy is obtained within specified limits for each alloying component.

The need to use alloys is due to the low rate of dissolution of refractory components in pure form in liquid aluminum and magnesium, as well as by increasing the degree of absorption of easily oxidized alloying elements. In most aluminum and magnesium alloys, the alloying component is in the form of crystals of intermetallic compounds, in some magnesium alloys - in the form of small particles in pure form. Taking into account the nature of the distribution of the component in alloy materials and the rate of its dissolution in melts of aluminum or magnesium, it is possible to obtain a given content of the alloying component in the alloy by adding a certain amount of alloy to the solid charge or directly to the melt. Important property The ligature has a significantly lower melting point than the refractory component. Thanks to this, alloys based on aluminum or magnesium do not need to be overheated to high temperatures, as a result, the loss of the base and alloying metal is reduced. The use of alloys with low-melting elements makes it possible to reduce losses of the latter due to evaporation and oxidation. With the help of alloys, it is much easier to introduce into the melt elements that have a melting point that is sharply different from the main melt, have high vapor elasticity and are easily oxidized at melt preparation temperatures, as well as in cases where the introduction of an alloying element directly into the melt is accompanied by a strong exothermic effect, leading to significant overheating of the melt, or when the evaporation of an alloying element is accompanied by the release of toxic vapors into the workshop atmosphere.

Since the alloy is an intermediate alloy, there are no requirements for mechanical properties. But due to the introduction of it in large quantities into the main melt, the hereditary influence of charge materials on the structure of castings and ingots, as well as increased requirements for the quality of castings and semi-finished products, a number of requirements are imposed on alloy ingots:

1. A sufficiently low melting point of the alloy, which will ensure the minimum temperature of the element additive, which is 100-200 ° C above the liquidus temperature. The low temperature of the liquidus of the alloy contributes to the rapid dissolution of the alloying element and its uniform distribution throughout the volume of the melt, especially under the condition of sufficiently intense and uniform mixing of the latter. Only alloys of the Al-Cu, Al-Si systems have a liquidus temperature close to or lower than the melting temperature of the base, as follows from Table. 20.

The liquidus temperature of the remaining alloys continuously increases with increasing content of the refractory alloying component in them.

From an economic point of view, it is better to have alloys with a high content of alloying component due to the saving of working space for storing the alloy, vehicles, consumption of primary aluminum and its waste. Since at present alloys are prepared mainly in reverberatory furnaces from pure metals, the content of titanium, zirconium and chromium in melts is usually 2-5%. With a higher content of these metals in alloys, a very high (1200-1400 ° C) temperature is required. With an increase in the component content in the master alloy, with the existing organization of casting it in ingots, coarse accumulations of intermetallic compounds are formed, the dissolution of which requires additional holding time of the alloy or an increase in the temperature of the latter.

2. Uniform distribution of alloying elements over the cross section of the pig. To avoid heterogeneous chemical composition of the pigs, it is necessary to thoroughly mix the melt before casting, and the casting itself must be done as quickly as possible. The heterogeneous distribution of the element in pigs can be a consequence of two reasons. Firstly, the low rate of solidification of the pig, and secondly, the non-uniform distribution of the element in the liquid alloy before casting. In turn, the heterogeneous composition of the liquid alloy depends on the difference in the density of the phase components of the alloy. In magnesium alloys, in which the alloying element is usually present in pure form, this factor operates constantly; in aluminum, the segregation of intermetallic compounds by density develops when the temperature of the alloy decreases below its liquidus.

3. Low evaporation and oxidation of the alloying element when introducing it into the melt from the alloy.

4. Easy crushing of master alloy pigs into small pieces for more accurate weighing of the charge; at the same time, the ligature must be sufficiently technologically advanced during casting. For example, an increase in the manganese content in a double master alloy by more than 15% leads to cracking of the pig, which complicates its transportation and storage.

N. E. Kalinina, V. P. Beloyartseva, O. A. Kavac

MODIFICATION OF CASTING ALUMINUM ALLOYS WITH POWDER COMPOSITIONS

The influence of dispersed refractory modifiers on the structure and properties of cast aluminum alloys is presented. A technology has been developed for modifying aluminum alloys of the L!-81-Md system with a powder modifier of silicon carbide.

Introduction

The development of new components of rocket and space technology poses the task of increasing the structural strength and corrosion resistance of cast aluminum alloys. Ukrainian launch vehicles use silumins of the aluminum-silicon system, in particular, AL2, AL4 and AL4S alloys, the chemical compositions of which are given in Table 1. Alloys AL2 and AL4S are used to cast critical parts that make up the turbopump unit of a rocket engine. Foreign analogues of domestic silumins are alloys 354, C355 of the A!-B1-Si-Md system, alloys 359 of the A!-B1-Md system and A357 of the A!-B1-Md-Be system, which are used for casting housings for electronic units and guidance systems rockets.

Research results

Improving the mechanical and casting characteristics of aluminum alloys can be achieved by introducing modifier elements. Modifiers for cast aluminum alloys are divided into two fundamentally different groups. The first group includes substances that create a highly dispersed suspension in the melt in the form of intermetallic compounds, which serve as a substrate for the resulting crystals. The second group of modifiers includes surfactants, the effect of which is reduced to adsorption on the faces of growing crystals and thereby inhibiting their growth.

Modifiers of the first kind for aluminum alloys include elements I, 2g, B, Bb, which are included in the composition of the studied alloys in amounts up to 1% by weight. Research is underway on the use of such refractory metals as BS, H11, Ta, V as modifiers of the first type. Modifiers of the second type are sodium,

potassium and their salts, which are widely used in industry. Promising directions include the use of elements such as Kb, Bg, Te, Fe as modifiers of the second kind.

New directions in the modification of cast aluminum alloys are being pursued in the field of using powder modifiers. The use of such modifiers facilitates the technological process, is environmentally friendly, and leads to a more uniform distribution of the introduced particles over the cross-section of the casting, which increases the strength properties and ductility characteristics of the alloys.

It should be noted the results of research by G.G. Krushenko. The powder modifier boron carbide B4C was introduced into the composition of the AL2 alloy. As a result, an increase in ductility was achieved from 2.9 to 10.5% with an increase in strength from 220.7 to 225.6 MPa. At the same time, the average macrograin size decreased from 4.4 to 0.65 mm2.

The mechanical properties of hypoeutectic silumins mainly depend on the shape of eutectic silicon and multicomponent eutectics, which have the shape of “Chinese characters”. The work presents the results of modifying alloys of the A!-B1-Cu-Md-2n system with particles of TiN titanium nitrides less than 0.5 microns in size. A study of the microstructure showed that titanium nitride is located in the aluminum matrix, along grain boundaries, near silicon wafers and inside iron-containing phases. The mechanism of influence of dispersed TiN particles on the formation of the structure of hypoeutectic silumins during crystallization is that the bulk of them is pushed out by the crystallization front into the liquid phase and takes part in the grinding of the eutectic components of the alloy. Calculations showed that when using

Table 1 - Chemical composition

Alloy grade Mass fraction of elements, %

A1 Si Mg Mn Cu Zn Sb Fe

AL2 Base 10-13 0.1 0.5 0.6 0.3 - 1.0

AL4 8.0-10.5 0.17-0.35 0.2-0.5 0.3 0.3 - 1.0

AL4S 8.0-10.5 0.17-0.35 0.2-0.5 0.3 0.3 0.10-0.25 0.9

© N. E. Kalinina, V. P. Beloyartseva, O. A. Kavac 2006

formation of titanium nitride particles with a size of 0.1-0.3 microns and when their content in the metal is about 0.015 wt.%. the particle distribution was 0.1 µm-3.

The publication discusses the modification of the AK7 alloy with dispersed refractory particles of silicon nitrides 813^, as a result of which the following mechanical properties are achieved: stB = 350-370 MPa; 8 = 3.2-3.4%; HB = 1180-1190 MPa. When introducing titanium nitride particles into the AK7 alloy in an amount of 0.01-0.02% wt. temporary tensile strength increases by 12.5-28%, relative elongation increases by 1.3-2.4 times compared to the unmodified state. After modifying the AL4 alloy with dispersed particles of titanium nitride, the strength of the alloy increased from 171 to 213 MPa, and the relative elongation increased from 3 to 6.1%.

The quality of foundry compositions and the possibility of their production depend on a number of parameters, namely: the wettability of the dispersed phase by the melt, the nature of dispersed particles, the temperature of the dispersed medium, and the mixing modes of the metal melt when introducing particles. Good wettability of the dispersed phase is achieved, in particular, by introducing surface-active metal additives. In this work, we studied the effect of additives of silicon, magnesium, antimony, zinc and copper on the assimilation of silicon carbide particles of the fraction up to 1 micron by liquid aluminum of grade A7. BYU powder was introduced into the melt by mechanical mixing at a melt temperature of 760±10 °C. The amount of introduced aluminum was 0.5% by weight of liquid aluminum.

Antimony somewhat impairs the absorption of administered BYU particles. Elements that produce alloys of eutectic composition (B1, 2p, Cu) with aluminum improve absorption. This effect is apparently associated not so much with the surface tension of the melt, but with the wettability of the SC particles by the melt.

A series of experimental melts of aluminum alloys AL2, AL4 and AL4S, into which powder modifiers were introduced, was carried out at the State Enterprise PA "Yuzhny Mashinostroitelny Zavod". Melting was carried out in an induction furnace SAN-0.5 with casting into chill molds from of stainless steel. The microstructure of the AL4S alloy before modification consists of coarse dendrites of the α-solid solution of aluminum and the α(D!)+B1 eutectic. Modification with silicon carbide BS

made it possible to significantly refine the dendrites of the a-solid solution and increase the dispersion of the eutectic (Fig. 1 and Fig. 2).

The mechanical properties of AL2 and AL4S alloys before and after modification are presented in Table. 2.

Rice. 1. Microstructure of AL4S alloy before modification, x150

Rice. 2. Microstructure of AL4S alloy after modification B1S, x150

Table 2 - Mechanical properties

Alloy grade Casting method Type of heat treatment<зВ, МПа аТ, МПа 8 , % НВ

AL2 Chill T2 147 117 3.0 500

AL2, modified 8Yu Chill 157 123 3.5 520

AL4S Chill T6 235 180 3.0 700

AL4S, modified 8Yu Chill 247 194 3.4 720

In this work, the effect of temperature on the degree of assimilation of refractory particles T1C and B1C was studied. It has been established that the degree of assimilation of powder particles by the AL4S melt changes sharply with temperature. In all cases, maximum absorption was observed at a temperature specific to a given alloy. Thus, the maximum assimilation of Tiu particles was achieved at the melt temperature

700......720 °C, at 680 °C absorption decreases. At

When the temperature rises to 780......790 °C, the absorption of TI drops by 3......5 times and continues to decrease with a further increase in temperature. A similar dependence of assimilation on the melt temperature was obtained for BU, which has a maximum at 770 °C. A characteristic feature of all dependences is a sharp drop in absorption upon entering the two-phase region of the crystallization interval.

Uniform distribution of dispersed silicon carbide particles in the melt is ensured by stirring. With increasing mixing time, the degree of absorption of dispersed particles worsens. This indicates that the particles initially assimilated by the melt are subsequently partially removed from the melt. Presumably, this phenomenon can be explained by the action of centrifugal forces, pushing foreign dispersed particles, in this case BS, towards the walls of the crucible, and then bringing them to the surface of the melt. Therefore, during smelting, stirring was not carried out continuously, but was periodically resumed before selecting portions of metal from the furnace.

The mechanical properties of silumins are significantly affected by the particle size of the introduced modifier. The mechanical strength of casting alloys AL2, AL4 and AL4S increases linearly as the particle size of powder modifiers decreases.

As a result of the theoretical and experimental

Experimental studies have developed technological regimes for producing high-quality cast aluminum alloys modified with refractory powder particles.

Research has shown that when dispersed particles of silicon carbide are introduced into aluminum alloys AL2, AL4, AL4S, the structure of silumins is modified, primary and eutectic silicon is crushed and takes on a more compact form, the grain size of the a-solid solution of aluminum decreases, which leads to an increase in strength characteristics of modified alloys by 5-7%.

Bibliography

1. Fridlyander I.N. Metallurgy of aluminum and its alloys. - M.: Metallurgy, 1983. -522 p.

2. Krushenko G.G. Modification of aluminum-silicon alloys with powdered additives // Materials of the II All-Union Scientific Conference "Patterns of formation of the structure of eutectic type alloys." - Dnepropetrovsk, 1982. - P. 137-138.

3. Mikhalenkov K.V. Formation of the structure of aluminum containing dispersed particles of titanium nitride // Casting processes. - 2001. -№1.- P. 40-47.

4. Chernega D.F. The influence of dispersed refractory particles in the melt on the crystallization of aluminum and silumin // Foundry production, 2002. - No. 12. - P. 6-8.

Received by the editor on May 6, 2006.

The infusion of dispersed refractory modifier1v into the structure of that power-east is given! Livarnyh aluminum1n1evih alloy1v. The technological modification of the aluminum alloy in the Al-Si-Mg system was completed with a powder modifier of silicon carb1d.

The influence of fine refractory modifiers on structure and properties of foundry aluminum alloys is given. The technology of modifying aluminum alloys of the Al-Si-Mg system by the powder modifier carbide of silicon is developed.

At the initial stage of development of aluminum alloys, it was noted that small impurities or special titanium additives (hundredths or tenths of a percent) sharply refine the grain of cast aluminum. In 1914, K. Schirmeister published an article in which he showed the beneficial effect of small titanium additions on the fracture structure of small aluminum ingots. The effect of grain refinement of cast aluminum by introducing special additives was called modification.

In further widespread work on the modification of aluminum alloys, it was found that, in addition to titanium, the aluminum grain is crushed during crystallization by small additions of zinc, tungsten, molybdenum, boron, rhenium, tantalum, hafnium, vanadium, scandium, strontium and, to a much lesser extent, iron , nickel, chromium, manganese.

Due to the great importance of surface phenomena in modification processes, researchers have tried to determine criteria for surface activity that would allow the selection of modifiers necessary for a given change in structure.

Based on experiments by A.M. Korolkov put forward the ratio of atomic volumes of the additive as a criterion U d and solvent V p. If U d > U r, then the additive is surface active. Based on this criterion, he obtained data on assessing the activity of certain additives to aluminum at concentrations ranging from thousandths and hundredths of a percent to 10-20%. It has been shown that lithium, calcium, magnesium, tin, lead, antimony and bismuth are surface active towards aluminum. Alloying aluminum with copper, chromium, germanium and silver did not lead to a noticeable change in surface tension.

V.N. Elagin proved that the refinement of aluminum grains during crystallization is the result of a special interaction of transition metals with aluminum.

In table Table 1.3 shows the results illustrating the influence of the most powerful modifiers (titanium, tantalum, boron, zinc) when casting aluminum A99 into a chill mold.

Table 1.3

Results of the influence of the most powerful modifiers

According to V.I. Napalkova and S.V. Makhov, the structure of pure aluminum and its alloys depends on many parameters, which can be divided into two groups. The first group of parameters is determined by the physicochemical properties of refractory modifier particles. Taken together, these properties are expressed by chemical nature, structural, dimensional and adsorption factors. The second group should include the temperature-time regime of melting and casting alloys, the concentration of the modifier, the cooling rate of the ingot and the particle size of the intermetal and dov.

According to the mechanism of influence on the crystallization of the melt, all modifiers are divided into two classes: nucleation and surface-active action, and the first class modifiers are most important for grain refinement.

The ideal modifier is a particle that satisfies the following requirements: it must effectively grind the grain at a minimum concentration; in the melt be in a thermally stable and dispersed state; have minimal structural difference with the lattice of the modifying alloy; do not lose its modifying properties during remelting. None of the modifiers currently known has the full range of these properties.

The work presents the following mechanism for modifying aluminum and its alloys. When a modifier element is introduced into the aluminum melt, fluctuation phenomena occur, resulting in the formation of a pre-nucleus, the formation of which is due to the presence of suspended particles such as aluminum oxide, titanium carbide and others with a size of less than 1-2 microns. Fluctuation phenomena arise as a result of thermal supercooling of the melt, the magnitude of which is determined by the type of modifier element. The greater the thermal supercooling value, the greater the number of fluctuations, and the greater the number of impurities present in the melt becomes activated. The modifying ability of elements is determined by the interaction of their valence electrons with the valence electrons of aluminum. This interaction is due to the ability of the valence electrons of two atoms to collectivize to form an electron gas determined by the ionization potential.

Most authors note that with the addition of 0.10-0.15% Ti to high-purity aluminum and 0.07% Ti to technical-purity aluminum cast at temperatures of 690-710 ° C, noticeable modification is achieved. Particularly strong grain refinement is observed with the introduction of 0.20% Ti or more.

The work examines the effect of boron on grain refinement, but mainly the addition of boron is used for aluminum used in the electrical industry. R. Kissling and J. Wallas note that at a melt temperature of 690-710 ° C, the most effective addition is 0.04% B immediately before casting.

In wrought alloys of the Al-Mg and Al-Mn systems, the addition of 0.07% Ti ensures the production of a fine-grained structure in ingots cast by the continuous method and a fine-grained recrystallized structure in sheets.

M.V. Maltsev and his colleagues discovered the greatest grain refinement in ingots of aluminum wrought alloys at a titanium concentration of 0.05-0.10%. The obtained dependence of aluminum grain refinement on titanium concentration was explained by them by the nature of the aluminum-titanium phase diagram. Analysis of this dependence showed that a characteristic inflection appears on the “number of grains - additive” curve, the position of which is associated with the formation of TiAl 3 crystals at a titanium concentration of more than 0.15%. The strongest effect on the structure of aluminum is observed at titanium concentrations of 0.15-0.30%. When the titanium content is less than 0.15%, the refinement of the aluminum grain is practically very small. This is due to the uneven distribution of additives in macrovolumes of the liquid alloy. At a titanium concentration of more than 0.30%, slight grinding occurs, and at a concentration of 0.70% and above, aluminum grains become larger. In semi-finished products from modified aluminum alloys, due to the elimination of zoning in the structure, the mechanical properties are smoothed out, and their values ​​increase by 10-20% compared to semi-finished products from modified alloys. As established by M.V. Maltsev and his co-workers, a fine-grained structure of aluminum casting is obtained with the introduction of 0.05-0.10% B. The strongest refinement of the aluminum grain is observed with the addition of 0.20% B, and with a further increase in the boron concentration, the grain coarsens again.

Add boron in the amount of 0.05-0.10 % in alloy B95 significantly reduces the grain size in ingots, while the tensile strength of semi-finished products with the addition of boron is 15-20 MPa higher compared to semi-finished products from modified ingots. The introduction of boron in greater quantities than indicated leads to a sharp decrease in the ductility of semi-finished products from the B95 alloy.

The first experiments on grain refinement of aluminum alloys with combined additions of titanium and boron were carried out by A. Kibula and his colleagues from the British Association for the Research of Non-Ferrous Metals. In this work, to obtain the optimal modification effect, the following concentrations are recommended: 0.01-0.03% Ti and 0.003-0.010% B. Since pure aluminum does not contain impurities, it is the most difficult to modify. The Cavecchi company recommends adding 0.0025-0.0075% Ti and 0.0005-0.0015% B to pure aluminum, and 0.003-0.015% Ti and 0.0006-0.0003% B to aluminum wrought alloys. As the ingot size increases, the alloy addition must be increased. The alloy must be introduced only into primary aluminum and added to the melt 15-20 minutes before the start of casting.

The modification process was based on A. Kibula and later M.V. Maltsev, when studying the grain refinement in aluminum alloy ingots with the addition of titanium and jointly titanium and boron, laid down the theory of nucleation. As has been established, during the crystallization of alloys without titanium additives, supercooling occurs, the value of which reaches 1-2 °C, while with the introduction of 0.002-0.100% Ti, no supercooling is observed. In this case, a fine-grained structure is obtained over the cross-section of the ingot. All this gave reason to believe that the grain is crushed due to the presence of nuclei, on which crystallization of the melt begins. Such particles can be carbides, borides and aluminides of transition metals, having lattice parameters corresponding to the lattice parameter of the aluminum solid solution (4.04 A).

According to A. Kibula, the additive introduced as a modifier must meet the following requirements:

• sufficient stability in molten aluminum at high temperatures without changing the chemical composition;
• the melting point of the additive is higher than the melting point of aluminum;
• structural and dimensional correspondence between the additive and aluminum lattices;
• formation of fairly strong adsorption bonds with atoms of the modifying melt.

The criterion for the strength of these bonds, apparently, can be the surface tension at the melt-solid particle interface. The greater the surface tension, the worse the particle is wetted by the liquid phase and the less likely it is to use the particle as a crystallization center. Work on a large number of systems has shown that the catalytic activity of the substrate with respect to nucleation is determined not by the value of the lattice matching, but by the chemical nature of the substrate.

Studying the industrial alloy A1-5TMV produced by the company "Kavekki", the authors of the work came to the conclusion that the grain refinement of aluminum alloys is associated with the formation of TiAl 3 particles due to the structural and dimensional correspondence of their lattice to the lattice of the aluminum solid solution. Crystals of titanium diboride and boron aluminide do not participate in the modification process, as shown by the results of electron microscopic analysis. The addition of boron to the aluminum-titanium alloy promotes the formation of aluminide at concentrations

Experiments have shown that maximum degree modification is observed at a ratio of titanium to boron concentration of 5:1; with larger or smaller ratios, the modification effect decreases. Obviously, modification occurs when titanium aluminide predominates, although borides can also be nuclei during the solidification of aluminum. The main difference between these two types of nuclei is that the solidification of aluminum on titanium aluminide occurs without supercooling, whereas for borides some supercooling is required.

Most researchers argue that the modification effect is determined by the ratio of titanium and boron. So in the work this is explained by the fact that the introduction of a master alloy containing 2.2% Ti and 1% B into the aluminum melt provides the same modification effect as the addition of a master alloy containing 5% Ti and 1% B. But in the alloy Al-2 ,2Ti-lB titanium aluminide is present in small amounts or absent and the main constituent is titanium diboride, which serves as a nucleus for the solidification of aluminum. In the A1-5Ti-lB alloy, the main modifier is titanium aluminide, the nucleus of which is titanium diboride. It can accumulate along the crystallization front and dissolve a limited amount of aluminum. According to D. Collins, titanium aluminide and other intermetallides formed as a result of the peritectic reaction are very effective modifiers and refine grain even at low cooling rates.

As J. Morisot points out, the modification process is greatly influenced by the rate of crystallization, the presence of alloying components, which expand the crystallization range of the alloy and create concentration supercooling, as well as thermal supercooling in the melt near the interface.

The work outlines the following grain grinding mechanism. Before the crystallization front, the melt contains a sufficient amount of primary particles of TiB 2, ZrB 2, etc. In the Al-Ti-B alloy, the main modifier is the TiB 2 particle, the lattice of which is similar in structure and size to the aluminum lattice. Solidification of aluminum on titanium diboride particles is possible only with supercooling equal to 4.8 °C. A layer with an increased concentration of titanium is formed near the titanium boride due to its diffusion from the boride. The formation of a layer with an increased concentration of titanium makes it possible to explain why the ratio of titanium to boron in the master alloy exceeds the corresponding stoichiometric ratio in the TiB 2 compound. The size factor between the nucleus and the base of the alloy is not decisive, at least for borides.

It should be noted that the experimental data on melt supercooling in the presence of modifying additives are inconsistent. The work shows that supercooling in aluminum alloys with 0.3-0.8% Ti is a fraction of a degree. In this case, alloys with titanium crossing the peritectic horizontal are characterized by greater supercooling than extraperitectic ones.

The work carried out a study of the effect of titanium additives on the supercooling of aluminum in a volume of 10 μm 3 at a heat removal rate of 5-10 °C/min. The addition of 0.025% Ti reduced the supercooling of aluminum from 47 to 16 °C. The degree of supercooling is also significantly affected by the volume of the melt. Directly measure the temperature of the supercooled melt and adjust the heat removal rate to obtain reproducible results V.I. Danilov recommends in volumes of 0.25-0.50 cm 3.

According to the Japanese researcher A. Ono, the reason for the grinding of primary grains is the factor that determines the appearance of equiaxed crystals. Using the example of the Al-Ti alloy, it is shown that rapid cooling itself does not lead to the formation of equiaxed crystals in the rapid cooling zone. To form them, it is necessary to stir the melt. In this case, the growth of crystals deposited on the walls of the crystallizer during the solidification process is stopped. Due to supercooling and changes in solution concentration, crystal growth on the crystallizer wall is limited, and tensile stresses act at their base. As a result, the crystals are separated from the walls of the crystallizer, and an equiaxed structure is formed. A. It believes that in the grinding of grains, the main role is played by the effect of enveloping the bases of the crystals grown on the walls of the crystallizer with modifier elements; this is also observed when modifiers are introduced. Titanium envelops the bases of the crystals, which accelerates their separation from the walls of the crystallizer, and is an impurity for aluminum that is selectively captured by growing crystals. As a result, titanium segregation is observed at the bases of the crystals, which causes the crystals to envelop and inhibit their growth. Thus, in studies, the slowdown in crystal growth is explained by the segregation of dissolved elements during the solidification process and the mixing of the melt during solidification.

There is another original way to control the crystallization process, especially of thick-walled castings, developed in detail in relation to steel casting. In this case, sharp cooling of the melt in its entire volume is achieved by introducing metal powders into the metal stream during casting into a mold or other mold. During suspension solidification, due to the sharp cooling of the melt throughout the entire volume, high rates of crystal growth develop from many simultaneously formed crystallization centers. In this case, volumetric crystallization of the ingot is observed.

Recently, suspension casting has been used to eliminate columnar structure, axial porosity, segregation and hot cracks in steel castings. It will also be tested as a means to improve the structure of aluminum alloy castings. When choosing microrefrigerators, it is recommended to observe the principle of crystallographic correspondence, i.e. the material of microrefrigerators must be identical or close in its crystallographic characteristics to the alloy being processed. For the greatest effect, it is necessary that the melting temperature of microrefrigerators be close to the melting temperature of the alloy being processed.

It is also possible to introduce into the head part of the ingot solid bodies of the same composition as the alloy being poured, which, when melted, take away part of the heat from the liquid well of the ingot. E. Sheil achieved effective grain refinement of aluminum alloys by adding wire or tape of a certain thickness to the stream of the poured alloy. By this time in our country V.I. Danilov studied in detail the mechanism of grain refinement in ingots of various alloys by introducing seed material.

V.E. In 1940, Neumark proposed using a seed made of the same metal as the melt to refine the structure of the ingot. The seed was introduced in the form of pieces or chips in an amount of 1-2% into a slightly overheated melt before it was poured into the mold. The influence of the seed on the structure of the ingot depends on the overheating temperature of the melt, on the thoroughness of mixing the seed into the melt and on the casting method. Pure metals are more difficult to refine grains using seeds than alloys. An important circumstance is the value of surface tension at the crystal-melt interface, therefore, the lower the surface tension, the lower the work of formation of a crystalline nucleus and the greater the likelihood of obtaining a fine-crystalline ingot. The possibility of using a seed for certain metals and alloys is determined by the degree of deactivation of impurities when the melt is overheated. The higher the deactivation temperature, the more effective the effect of the seed on the structure of the ingot. To increase the temperature, a seed was used containing a small amount of an element that modifies the structure of the ingot: the seed was made of aluminum with 0.5% Ti. The use of such a seed led to a more significant refinement of the aluminum structure than when using a titanium seed.

Studies on refinement of the structure of the D16 alloy with a rod of the same composition have shown that with the introduction of a constant amount of filler material, the effect of grain refinement decreases with increasing temperature in the range of 670-720 ° C. At higher casting temperatures there is very little grinding. Increasing the amount of added material enhances grain refinement to the extent that the casting temperature decreases. These results are in full agreement with those developed by G.F. Balandin's ideas about the modifying and seeding effect of fragments of the solid phase in a crystallizing alloy.

The studies presented in the works convincingly show the hereditary influence of the grain structure of aluminum alloy ingots on the structure and properties of semi-finished products made from them. Since the quality requirements for products made from aluminum wrought alloys are strict, it is very important to correctly assess the feasibility of using a particular modification method and find ways to overcome its negative aspects. A wide variety of aluminum wrought alloys and features of the technological process for producing ingots, as well as a wide range of semi-finished products from these alloys require a differentiated approach to the choice of modification method, taking into account restrictions on the content of impurities, the different susceptibility of alloys to the formation of a columnar structure, and the precipitation of primary crystallizing intermetallic compounds. Often in factory practice it is necessary to find ways to eliminate the inhomogeneous or rough equiaxial structure of ingots. The question of optimal concentration and the feasibility of using one or another modifier when casting ingots of different sizes. In addition, scientists are searching for new materials that have a high modifying ability and have chemical composition, close to the modified alloy. Such materials can be obtained by combined methods of casting and metal forming. In particular, a technology has been proposed for producing a ligature tape used in modifying aluminum ingots in order to form a fine-grained structure in them. This technology consists of using a combined process of high-speed crystallization and hot plastic deformation of the resulting workpiece, resulting in additional crushing of intermetallic particles formed during crystallization. In addition, conditions are provided for the formation of finely differentiated sub-grain structures of the base of the ligature strip (rod, tape), which represents an additional modifying effect.

According to known data, the finest grain of aluminum is 0.13-0.20 mm (respectively, the number of grains in an area of ​​1 cm 2 of a section is 6000 and 2300) is achieved when using the best Al-Ti-B rod ligature from the company " Cavecchi." A significant advantage of the microstructure of the experimental master alloy made from alloys of the Al-Ti-B system, compared to the rod master alloy from Cavecchi, was the predominance of the globular morphology of TiAl 3 particles with smaller sizes and a much more uniform distribution of these particles throughout the volume of the aluminum matrix. The individual plate-shaped particles present in the structure are fragmented into blocks, the size of which does not exceed 10 microns. This advantage is confirmed by analysis of the fine structure of the experimental alloy tape (the size of the subgrains in the cross section ranged from 0.17 to 0.33 μm, and the particle size of titanium diborides was 0.036-0.100 μm). Studies of the fine structure of the alloy strip have shown that the combination of high-speed crystallization of the melt and continuous deformation of the solidified part of the metal forms a fine sub-grain structure. The average cross-sectional size of the subgrains is ~0.25 µm.

Thus, aluminum ingots modified with a master alloy obtained by the proposed method are characterized by a sharp refinement of the grain structure. Alloy alloys of the Al-Ti-B system or technical or high-purity aluminum can be used as the material for the alloy tape. In the latter cases, when modifying an aluminum ingot, grain refinement is ensured while eliminating contamination with impurities, including intermetallic compounds that cause ruptures of a thin strip (foil) during rolling.

The use of the developed technology, including melting of the master alloy, overheating, holding at an overheating temperature and accelerated crystallization on the surface of water-cooled crystallizer rolls, which were used as rolling mill rolls, made it possible to combine continuous high-speed crystallization of the strip with its hot plastic deformation in a single process. The results of studies on the modification of aluminum with alloy materials obtained using the proposed technology are given in Table. 1.4. Analyzing them, it can be noted that the use of alloy materials obtained using the technology of combined casting and pressure treatment gives no less of a modifying effect than the use of well-known alloys, for example, rods from the Cavecchi company. However, the use of Al-Ti-B master alloy does not always lead to the solution of production problems, since the presence of intermetallic inclusions in the modifier is often accompanied by their retention in the finished semi-finished product, which reduces its quality.

The use of fine-grained ingots will reduce the amount of losses from defects (breaks, cracks, inhomogeneities on the foil surface) and improve product quality. In this regard, attempts were also made to obtain a ligature tape from technically pure aluminum of the A5 and AVCh grades (Table 1.5).

Table 1.4

Change in grain size and number of grains per 1 cm 2 in Alkane test samples after aluminum modification, depending on the amount of introduced Al-Ti-B alloy ligature

ligature ligature	Original aluminum,	Amount of titanium, % wt.	Average grain size in Alkan-test sample, µm		Number of grains per 1 cm 2, pcs.		The degree of grain refinement after holding the melt for 5 minutes, times
			after holding the melt for
Known method
Rod with a diameter of 8 mm from Cavecchi (Al-3Ti-0.2B)
Suggested method
Ligature

Table 1.5

The influence of aluminum alloy tape on the grain size in an aluminum ingot after modification

	Amount of aluminum tape, % wt. (grade of aluminum)	Original ingot aluminum grade A7, microns			Average grain size of modified aluminum, microns	Number of grains per 1 cm 2 in modified aluminum, pcs.
			1 minute after inserting the tape		7.5 minutes after tape insertion

The research results showed that the number of grains in the modified aluminum is comparable to the same indicators of the master alloy made from the Al-Ti-B alloy. This gives grounds to assert that using high-speed crystallization-deformation methods it is possible to obtain new modifying materials, including from aluminum.

The use of tape as a modifying material is technologically unprofitable, since almost all foundry installations are equipped with devices for feeding master alloy in the form of a rod, so it is urgent to develop methods for producing modifiers that would have a technologically advantageous shape and size, and also would not make changes to the chemical composition of the alloy ingots undergoing modification.

Thus, to introduce into production technologies for producing deformed semi-finished products with a high level of mechanical properties, it is necessary to produce new modifying materials using high-speed crystallization of aluminum alloy in water-cooled rolls, combined with hot deformation of the metal.

Melting of most aluminum alloys is not difficult. Alloying components, with the exception of magnesium, zinc, and sometimes copper, are introduced in the form of alloys. When smelting small portions of casting alloys in crucible furnaces, protective fluxes, as a rule, are not used. A mandatory operation is refining to remove non-metallic inclusions and dissolved hydrogen. The most difficult to melt are aluminum-magnesium and multicomponent heat-resistant alloys.
When melting wrought alloys, special attention is paid to cleaning the furnace from slag and residue from previous melting. When switching to another brand of alloy, in addition to transition melts, the furnace and mixers are washed to remove remnants of the old alloy. The amount of metal for washing should be at least a quarter of the furnace capacity. The metal temperature during washing is maintained 40-50 °C above the alloy casting temperature before washing. To speed up cleaning, the metal is intensively stirred in the furnace for 8-10 minutes. For washing, aluminum or remelting is used. In cases where the metal is completely drained from the furnace, you can limit yourself to washing with fluxes. Alloys are melted under submerged arc
Charge materials are loaded in the following sequence: pig aluminum, bulky waste, remelting, alloys (pure metals). It is allowed to load dry shavings and small-sized scrap into liquid metal at a temperature not exceeding 730 °C. Copper is introduced into the melt at a temperature of 740-750 °C, silicon - at 700-740 °C using a bell. Zinc is loaded before magnesium, which is usually added before the metal is drained. The maximum permissible overheating for cast alloys is 800-830 °C, and for deformable alloys 750-760 °C.
When melted in air, aluminum oxidizes. The main oxidizing agents are oxygen and water vapor. Depending on the temperature and pressure of these gases, as well as the kinetic conditions of interaction, aluminum oxide Al2O3, as well as Al2O and AlO are formed as a result of the oxidation of aluminum. The probability of formation increases with increasing temperature and decreasing partial pressure of oxygen in the system. IN normal conditions melting, the thermodynamically stable phase is solid aluminum oxide γ-Al2O3, which does not dissolve in aluminum and does not form fusible compounds with it. When heated to 1200 °C, γ-Al2O3 recrystallizes into α-Al2O3. As oxidation occurs, a dense, durable oxide film with a thickness of 0.1-10 microns is formed on the surface of solid and liquid aluminum, depending on the temperature and duration of exposure. When this thickness is reached, oxidation practically stops, since the diffusion of oxygen through the film slows down sharply.
The oxidation process of liquid aluminum alloys is very complex and insufficiently studied. Available literature data show that the intensity of oxidation of the alloy components is a function of the oxygen pressure, the dissociation pressure of their oxides, the concentration of the components in the alloy, the rate of diffusion of atoms towards oxygen atoms, the interaction of oxides with each other, etc. The kinetics of oxidation is determined by the continuity, density and strength of the oxide films. At the same concentration, the most active elements are oxidized first, in which the formation of oxide is associated with the greatest decrease in the isobaric-isothermal potential.
Most alloying elements (copper, silicon, manganese) do not have a significant effect on the oxidation of aluminum and protective properties oxide film, since they have a ratio VMem0/mVMe≥1. The oxide film on binary aluminum alloys with these elements at low concentrations consists of pure γ-Al2O3. At significant contents of these elements, solid solutions of oxides of alloying elements in γ-Al2O3 and corresponding spinels are formed.
Alkali and alkaline earth metals (potassium, sodium, barium, lithium, calcium, strontium, magnesium), as well as zinc (0.05-0.1%) greatly increase the oxidation of aluminum. The reason for this is the loose and porous structure of the oxides of these elements. The oxide film on double melts in this case is enriched with oxides of alkali and alkaline earth metals. To neutralize the harmful effects of zinc, 0.1-0.15% Mg is introduced into aluminum melts.
Alloys of aluminum and magnesium form an oxide film of variable composition. At a low magnesium content of 0.005% (by mass), the oxide film has the structure of γ-Al2O3 and is a solid solution of MgO in γ-Al2O3; with a content of 0.01-1.0% Mg, the oxide film consists of spinel (MgO*Al2O3) of variable composition and magnesium oxide crystals; with a content of more than 1.5% Mg, the oxide film consists almost entirely of magnesium oxide.
Beryllium and lanthanum slow down the oxidation of aluminum alloys. The addition of 0.01% beryllium or lanthanum reduces the oxidation rate of Al-Mg alloys to the level of aluminum oxidation. The protective effect of these elements is explained by the compaction of the oxide film by filling the resulting pores with beryllium and lanthanum oxides.
The oxidation of aluminum melts is greatly reduced by fluorine and gaseous fluorides (SiF4, BF3, SF6, etc.), present in the furnace atmosphere in amounts up to 0.1% (by weight). Adsorbed on the surface of the oxide film, they reduce the rate of oxygen penetration to the metal surface.
Mixing the melt during the melting process is accompanied by a violation of the integrity of the oxide film and the mixing of its fragments into the melt. The enrichment of melts with oxide inclusions also occurs as a result of exchange reactions with the lining of melting devices. The most significant influence on the degree of contamination of melts by films is exerted by the surface oxidation of the original primary and secondary charge materials. Negative role This factor increases as the compactness decreases and the specific surface area of ​​the material increases.
The oxide film of the charge is also a source of saturation of the melt with hydrogen, since it consists of 30-60% Al(OH)3. Chemically bound moisture is difficult to remove from the surface of the charge materials even at a temperature of 900 C. The hydroxide, entering the melt, greatly saturates it with hydrogen. For this reason, it is undesirable to introduce shavings, sawdust, trimmings, spills and other non-compact waste into the charge. Of particular importance is the organization of storage and timely processing of waste and return own production, preventing oxidation and corrosion with the formation of hydroxides. The introduction of own returns into the charge is also associated with the inevitable accumulation of harmful iron impurities in the alloys, which form complex solid intermetallic compounds with the alloy components, reducing the plastic properties and impairing the cutting processing of castings.
Along with oxides and intermetallic compounds, the melt may also contain other non-metallic inclusions - carbides, nitrides, sulfides. However, their number is small compared to the content of oxides. The phase composition of nonmetallic inclusions in aluminum alloys is varied. In addition to aluminum oxides, they may contain magnesium oxide (MgO), magnesium spinel (MgAl2O4), aluminum, magnesium, titanium nitrides (AlN, Mg3N2, TiN), aluminum carbide (Al4C3), aluminum and titanium borides (AlB2, TiB2) and etc. The bulk of inclusions are oxides.
Depending on their origin, non-metallic inclusions found in alloys can be divided into two groups: dispersed inclusions and films. The bulk of dispersed inclusions have a size of 0.03-0.5 microns. They are relatively evenly distributed in the volume of the melt. The most probable thickness of the oxide films is 0.1-1.0 microns, and the length is from tenths of a millimeter to several millimeters. The concentration of such inclusions is relatively small (0.1-1.0 mm2/cm2), and the distribution is extremely uneven. When melts stand, large inclusions may float or settle. However, due to the large specific surface area of ​​the films and the small difference between their density and the density of the melts, the floating (deposition) is slow; most of the films remain in the melt and, when filling the mold, are carried into the casting. Finely dispersed suspensions separate even more slowly. Almost all of them go into casting.
During smelting, aluminum is saturated with hydrogen, the content of which can reach 1.0-1.5 cm3 per 100 g of metal. The main source of hydrogen is water vapor, the partial pressure of which in the atmosphere of gas melting furnaces can reach 8-16 kPa.
The influence of alloying elements and impurities on the equilibrium solubility of hydrogen in aluminum has been little studied. It is known that copper and silicon reduce the solubility of hydrogen, and magnesium increases it. The solubility of hydrogen is also increased by all hydroforming elements (titanium, zirconium, lithium, sodium, calcium, barium, strontium, etc.). Thus, an aluminum alloy with 2.64% Ti can release up to 25 cm3 of hydrogen per 100 g, and an aluminum alloy with 5 % Zr - 44.5 cm3 per 100 g. Alkali and alkaline earth metals (sodium, lithium, calcium, barium), which form hydrides, most actively increase the solubility of hydrogen and aluminum.
A significant proportion of hydrogen dissolved in alloys is the gas introduced by alloys and electrolytic copper. For example, an aluminum-titanium alloy, depending on the smelting technology, can contain up to 10 cm3 of hydrogen per 100 g, and electrolytic copper with build-ups - up to 20 cm3 per 100 g. Cast alloys contain more impurities and non-metallic inclusions than wrought alloys. Therefore, they are more prone to absorb gases
The kinetics of the process of hydrogenation of aluminum melts is limited by the mass transfer of hydrogen in the liquid metal, through the surface oxide film and in a gaseous environment. The most significant influence on mass transfer is exerted by the composition of the alloy and the content of non-metallic inclusions, which determine the permeability of the oxide film, the diffusion mobility of hydrogen and the possibility of its release from the melt in the form of bubbles. The permeability of the film is also significantly influenced by the composition of the gaseous medium. The diffusion mobility of hydrogen in aluminum is reduced by copper, silicon and especially magnesium, manganese and titanium. Finely dispersed non-metallic inclusions, having a high adsorption capacity for hydrogen, greatly slow down its diffusion mobility in aluminum melts.
The aluminum oxide film has low permeability to hydrogen atoms; it slows down the reaction between the melt and atmospheric moisture. With a film thickness of 1-10 microns, gas exchange between the metal and the atmosphere practically stops. The permeability of the film is greatly influenced by the composition of the alloy. All elements that increase the oxidation of aluminum (magnesium, lithium, sodium, strontium, calcium) increase the permeability of the oxide film to hydrogen. Alloying elements (copper, zinc, silicon) have little effect on gas exchange. They somewhat loosen the oxide film and therefore contribute to faster saturation of the alloys with hydrogen.
The hydrogen permeability of the oxide film is significantly affected by the composition of the atmosphere above the melt. The permeability of the film increases significantly if Cl2, C2Cl6, BF4, SiF4, freons and other halogens are present in the gas environment. Chlorides, having a high affinity for aluminum, are adsorbed, penetrate under the oxide film and destroy it as a result of the formation of gaseous aluminum chloride. Fluorides interact less actively with aluminum. Interacting with the oxide film, they contribute to the dehydration of its surface and the desorption of molecules and oxygen atoms. Having a high adsorption capacity, fluorides occupy the vacated active centers on the film and create oxyfluoride complexes such as Al2O2F2, which stop the access of oxygen and water vapor to the melt, making the film thin and permeable to hydrogen. Liquid fluxes containing fluorides also destroy the oxide film and facilitate degassing of melts.
Dissolved hydrogen, released during the crystallization of melts, causes the formation of gas and gas-shrinkage porosity in castings. With increasing hydrogen concentration, the gas porosity of the castings increases. The susceptibility of aluminum alloys to gas porosity is determined by the degree of supersaturation of the solid solution with hydrogen, which is expressed by the ratio η - (Cl-Stm)/Stm, where Cl and Stm are the concentrations of hydrogen in the liquid and solid alloy, cm3/100 g. Gas porosity does not form when Stp=Com. The degree of supersaturation of the solid solution increases with increasing cooling rate.
For each alloy, there are limiting hydrogen concentrations below which gas pores do not form in the castings at given cooling rates. For example, in order to prevent the formation of gas pores during the solidification of thick-walled castings from the Al - 7% Si alloy, the hydrogen content in the melt should not exceed 0.15 cm3 per 100 g. The limiting hydrogen content in duralumin is considered to be 0.12-0. 18 cm3 per 100 g, depending on the intensity of cooling during crystallization.
The protection of aluminum melts from oxidation and hydrogen absorption is achieved by submerged arc melting in a weakly oxidizing atmosphere. As a coating flux when melting most alloys containing no more than 2% Mg, a mixture of sodium and potassium chlorides (45% NaCl and 55% KCl) is used in an amount of 1-2% by weight of the charge. The composition of the flux corresponds to a solid solution with minimum temperature melting point 660 °C. For this purpose, a flux with a more complex composition is also recommended (Table 12).

For aluminum-magnesium alloys, carnallite (MgCl2*KCl) and mixtures of carnallite with 40-50% barium chloride or 10-15% calcium fluoride are used as a coating flux. If the use of flux is impossible, protection against oxidation is carried out by introducing beryllium (0.03-0.05%). Protective fluxes are widely used when melting alloys in reverberatory furnaces.
To prevent interaction with moisture, measures are taken to remove it from the lining of melting furnaces and casting devices, from refining and modifying fluxes; melting and casting tools are calcined and painted, and the charge materials are heated, cleaned, and dried.
However, no matter how carefully the melt is protected, when melting in air it always turns out to be contaminated with oxides, nitrides, carbides, slag and flux inclusions, and hydrogen, so it must be cleaned before pouring into molds.

Melt refining

To clean aluminum alloys from suspended non-metallic inclusions and dissolved hydrogen, settling, purging with inert and active gases, treatment with chloride salts and fluxes, vacuuming, filtration through mesh and granular filters, and electroflux refining are used.
How independent process settling may be applicable in cases where the density difference is large enough and the particle size is not too small. But even in these cases the process is slow, increased fuel consumption is required and it turns out to be ineffective.
Purification of melts by blowing with inert or active gases is based on the occurrence of two processes of diffusion of dissolved gas into bubbles, blowing and floating action of bubbles in relation to inclusions and tiny gas bubbles. Refining is carried out the more successfully, the smaller the size of the bubbles of the purged gas and the more uniform their distribution throughout the volume of the melt. In this regard, the method of processing melts with inert gases using porous ceramic inserts deserves special attention. But compared to other methods of introducing inert gases into melts, blowing through porous inserts is the most effective.
Blowing melts with gases is widely used in foundries for the production of ingots. It is carried out in special lined boxes installed along the path of metal transfer from the mixer to the crystallizer. For refining aluminum melts, nitrogen, argon, helium, chlorine and its mixture with nitrogen (90%), purified from moisture and oxygen, are used.
Blowing with nitrogen or argon is carried out at 720-730 °C. The duration of blowing, depending on the volume of the melt, ranges from 5-20 minutes; gas consumption is 0.3-1% of the melt mass. This treatment makes it possible to reduce the content of non-metallic inclusions to 1.0-0.5 mm2/cm2 according to the technological test of V.I. Dobatkina and BK. Zinoviev, and the hydrogen content is up to 0.2-0.15 cm3 per 100 g of metal.
The treatment of melts with chlorine is carried out in sealed chambers or ladles that have a lid with gases vented into the ventilation system. Chlorine is introduced into the melt through tubes with nozzles at 710-720 °C. The duration of refining at a chlorine pressure of 108-118 kPa is 10-12 minutes; chlorine consumption - 0.2-0.8% of the mass of the melt. The use of chlorine provides a higher level of purification compared to technical nitrogen and argon. However, the toxicity of chlorine, the need to process melts in special chambers and the difficulties associated with its drying significantly limit the use of chlorination of melts in industrial conditions. Replacing chlorine with a mixture of it and nitrogen (90%) provides a fairly high level of purification, but does not solve the problems associated with toxicity and drying.
Degassing by blowing is accompanied by losses of magnesium: when treated with nitrogen, 0.01% of magnesium is lost; when treated with chlorine, these losses increase to 0.2%.
Refining with chlorides is widely used in the shaped foundry industry. For this purpose, zinc chloride, manganese chloride, hexachloroethane, titanium tetrachloride and a number of other chlorides are used. Due to the hygroscopicity of chlorides, they are subjected to drying (MnCl2, C3Cl6) or remelting (ZnCl2). The technology of refining with chlorides consists of introducing them into the melt with continuous stirring with a bell until the release of gaseous reaction products stops. Zinc and manganese chlorides are introduced in an amount of 0.05-0.2% at a melt temperature of 700-730 ° C; hexachloroethane - in an amount of 0.3-0.7% at 740-750 °C in several stages. With decreasing temperature, the efficiency of refining decreases due to an increase in the viscosity of the melts; refining at higher temperatures is impractical, since it is associated with intense oxidation of the melt.
Currently, in shaped casting shops for refining, tablets of the drug “Degaser” are widely used, consisting of hexachloroethane and 10% (by weight) barium chloride, which are introduced into the melt without the use of “bells”. Having a greater density than the melt, the tablets sink to the bottom of the container, ensuring that the entire volume of the melt is processed.
Chloride salts interact with aluminum according to the reaction: 3MnCl2 + 2Al → 2AlCl3 + 3Mn.
Bubbles of aluminum chloride, rising to the surface of the melt, entrain suspended non-metallic inclusions; Hydrogen dissolved in the metal diffuses into the bubbles, and the melt is purified. After mixing is completed, the melt is allowed to stand for 10-45 minutes at 720-730 °C to remove small gas bubbles.
Refining with chlorides is carried out in furnaces or ladles with a small specific surface area of ​​the melt. In furnaces with a small-height melt layer, refining with chlorides is ineffective. In terms of the level of purification from non-metallic inclusions and gas, treatment with chlorides is inferior to purging with chlorine.
Cleaning aluminum melts with fluxes is used in the melting of cast and wrought alloys. For refining, fluxes are used based on chloride salts of alkali and alkaline earth metals with the addition of fluoride salts - cryolite, fluorspar, sodium and potassium fluorides (Table 13).

In the practice of melting most aluminum wrought alloys, flux No. 1 is used for refining.
To clean aluminum and magnesium alloys, carnallite-based fluxes are used - 80-90% MgCl2*KCl, 10-20% CaF2, MgF2 or K3AlF6. Pre-melted and dried fluxes in an amount of 0.5-1% by weight of the metal are poured onto the surface of the melt at 700-750 °C. Then the flux is vigorously mixed into the melt for 3-5 minutes, the slag is removed and the melt is allowed to stand for 30-45 minutes. After the slag is removed again, the melt is used to fill casting molds. When processing large volumes of metal, flux is introduced to the bottom of the melt using a “bell”.
For refining cast aluminum alloys (silumins), fluxes No. 2 and 13 are widely used. They are introduced into the melts in liquid form in an amount of 0.5-1.5% (by weight) and vigorously kneaded. They contribute to the destruction of the foam formed when filling the dispensing ladles and enrich the melts with sodium.
A high level of degassing is obtained by vacuuming. This cleaning method is used mainly in shaped foundries. Its essence lies in the fact that the metal smelted using standard technology in conventional furnaces is poured into a ladle, which is then placed in a vacuum chamber. The metal in the chamber is maintained at a residual pressure of 1330 Pa for 10-30 minutes; The melt temperature is maintained within 720-740 °C. In cases where evacuation is carried out without heating, the melt is overheated to 760-780 °C before processing. The installation diagram for vacuum degassing is shown in Fig. 93.

IN last years To clean aluminum melts from non-metallic inclusions, filtration through mesh, granular and porous ceramic filters is increasingly used on a large scale. Mesh filters are widely used to clean melts from large inclusions and films. They separate those inclusions whose size is larger than the mesh cell. For the manufacture of mesh filters, glass fabric of various brands with cell sizes from 0.5x0.5 to 1.5x1.5 mm and metal mesh(made of titanium). Filters made of fiberglass are installed in distribution boxes and crystallizers, in gating channels and dispensing crucibles (Fig. 94), their use makes it possible to reduce the content of large non-metallic inclusions and films by 1.5-2 times; they do not affect the content of dispersed inclusions and hydrogen.

Grain filters provide a significantly greater cleaning effect. Distinctive feature they consist of a large contact surface with the metal and the presence of long thin channels of variable cross-section. The purification of metal melts from suspended inclusions when filtering through granular filters is due to mechanical and adhesion processes. The first of them plays a decisive role in the separation of large inclusions and films, the second - in the separation of fine inclusions. Due to the mesh effect, granular filters retain only those inclusions whose size exceeds the effective diameter of the intergranular channels. The smaller the diameter of the filter grains and the denser their packing, the higher the achieved level of purification of melts from large inclusions and films (Fig. 95).
As the thickness of the filter layer increases, the cleaning efficiency increases. Melt-wettable filters are more efficient than non-wettable filters.
Filters made from an alloy of calcium and magnesium fluorides make it possible to obtain castings from AL4, AK6 and AMg6 alloys that are 1.5-3 times less contaminated with large inclusions than filters made from magnesite.

The speed and mode of melt flow through the intergranular channels of the filter have a significant influence on the completeness of separation of large inclusions and films. With increasing speed, the possibility of sedimentation of inclusions from a moving flow under the influence of gravity decreases and the probability of washing away already settled inclusions as a result of hydrodynamic action, the degree of which is proportional to the square of the filtration speed, increases.
The efficiency of cleaning aluminum melts from finely dispersed inclusions using granular filters increases as the wetting of the filter and inclusions by the melt deteriorates.
For the manufacture of filters, fireclay, magnesite, alundum, silica, alloys of chloride and fluoride salts and other materials are used. The completeness of removal of suspended non-metallic inclusions depends on the nature of the filter material. The most effective filters are those made from fluorides (active materials) (Fig. 95 and 96).
Active materials, along with large inclusions and films, make it possible to separate up to 30-40% of finely dispersed suspensions and reduce the hydrogen content in alloys that have been refined with flux or chlorides by 10-20%. As finely dispersed suspensions are removed, the grain size in the castings increases, the gas content decreases, and the plastic properties of the alloys increase (Fig. 97). A high level of purification of the AK6 and AL4 alloys from inclusions and hydrogen is observed when using filters made of an alloy of calcium and magnesium fluorides with a grain size of 4- 6 mm in diameter and filter layer height 100-120 mm.

Granular filters, like mesh filters, are installed along the path of metal movement from the mixer to the mold. For continuous casting of ingots, the optimal installation location is the mold; in shaped casting, the filter is placed in a riser, dispensing crucible or sprue bowl.
Typical layouts of granular filters when casting shaped castings and ingots are shown in Fig. 98.
Before use, the filter is heated to 700-720 °C to remove adsorbed moisture and prevent freezing of the metal in the channels.

Filling is carried out in such a way that the upper level of the filter is covered with a layer of metal of 10-15 mm, and the outflow of metal after the filter occurs under the flooded level. If these conditions are met, the residual content of non-metallic inclusions and films in the casting can be increased to 0.02-0.08 mm2/cm2 according to V.I. technological test. Dobatkin and V.K. Zinoviev, i.e. 2-4 times reduced compared to filtering through mesh filters.
Most effective method cleaning aluminum melts from films and large non-metallic inclusions - electroflux refining. The essence of this process is to pass thin jets of melt through a layer of liquid flux while simultaneously applying a constant or alternating current, creating more favorable conditions for the adsorption of inclusions by flux as a result of a decrease in interfacial tension at the boundary with the metal. With an increase in the specific surface area and the duration of contact of the metal with the flux, the cleaning efficiency increases. Therefore, the designs of devices for flux and electroflux refining provide for jet fragmentation (Fig. 99).

The optimal mode of electroflux refining involves passing a stream of metal with a diameter of 5-7 mm, heated to 700-720 °C, through a layer of molten flux 20-150 mm thick with the imposition of a direct current field with a force of 600-800 A and a voltage of 6-12 V with the cathode polarization of the metal. With a flux consumption (carnallite with 10-15% CaF2, MgF2 or K3AlF6 for Al - Mg and Al - Mg - Si alloys and cryolite for other aluminum alloys) of 4-8 kg per 1 ton of melt and careful removal of moisture from the flux and casting devices , the content of large non-metallic inclusions in alloys AK6, AMg6, V95 can be reduced to 0.003-0.005 mm2/cm2 according to a technological test.
Unlike granular filters, electroflux refining does not affect the macrostructure of alloys, which indicates its lower efficiency in removing dispersed non-metallic inclusions.
Wrought and cast alloys are also subjected to refining to remove metal impurities: sodium, magnesium, zinc and iron.
Removal of sodium from aluminum and aluminum-magnesium deformable alloys AMg2, AMg6 is carried out by blowing the melts with chlorine or vapors of chlorides (C2Cl6, CCl4, TiCl4), freon (CCl2F2) and filtering through granular filters made of AlF3 with a grain size of 4-6 mm. The use of these methods makes it possible to increase the residual sodium content in the melt to 2/3*10-4%. The harmful effect of sodium on the technological properties of the alloy can be suppressed by introducing into the melt additives bismuth, antimony, tellurium or selenium, which form refractory intermetallic compounds with sodium.
In some cases, secondary aluminum alloys are purified from impurities of magnesium, zinc and iron by fluxing, vacuum distillation and sedimentation, followed by filtration. Removal of magnesium by flux is based on the reaction 2Na3AlF6 + 3Mg → 6NaF + 3MgF2 + 2A1. The surface of the melt is coated with a flux consisting of 50% cryolite and 50% sodium chloride. Then the alloy is heated to 780-800 °C and intensively mixed together with flux for 10-15 minutes. Reaction products that float to the surface of the melt are removed; with a high magnesium content (1-2.5%), the refining process is repeated several times. Using cryolite, the magnesium content in the melt can be reduced to 0.1%. Refining secondary aluminum alloys from magnesium can be successfully carried out with a flux consisting of 50% Na2SiF6, 25% NaCl and 25% KCl. For these purposes, you can use oxygen-containing fluxes, such as potassium chlorate (KClO3).
Melts are purified from magnesium and zinc in vacuum distillation furnaces at 950-1000°C. As a result of this processing, alloys containing 0.1-0.2% Mr and 0.02-0.05% Zn are obtained. Melts are purified from magnesium by distillation in cases where its content in the alloy is high and the use of purification by fluxing becomes unprofitable.
By settling, it is possible to reduce the iron content in an aluminum alloy to 1.7%, i.e., almost to the eutectic content, according to the aluminum-iron equilibrium state diagram. Further reduction is achieved by combining the settling process with the introduction of chromium, manganese or magnesium into the alloy. The addition of these elements shifts the eutectic point towards aluminum and promotes the separation of excess iron. By introducing 1-1.5% Mn into the melt, the iron content in it can be reduced to 0.7%. Adding magnesium in an amount of 25-30% allows you to increase the iron content to 0.1-0.2%. The process of separation of iron intermetallic compounds is accelerated by combining settling with filtration. Filtration is carried out through a basalt filter heated to 700 °C using a vacuum. Refining from iron with the help of magnesium is applicable for alloys containing no more than 1.0% Si. At a higher silicon content, silicides are formed, which greatly complicate filtration and remove a significant amount of magnesium from the cycle. In addition, the alloy is depleted of silicon.

Modification of alloys

Refinement of macrograins in castings is achieved by introducing small quantities (0.05-0.15% of the mass of the melt) of modifying additives (Ti, Zr, B, V, etc.) into the melt. This method is used to modify wrought alloys (V95, D16, AK6, etc.); It has not found wide application in the casting of shaped castings. Modifiers are introduced in the form of alloys with aluminum or copper at 720-750 °C.
With regard to deformable alloys, titanium is most widely used for refinement of the macrostructure. When it is introduced into melts in an amount of 0.05-0.15%, the macrograin of alloys in diameter is crushed to 0.5 mm. In this case, the crystallization centers are particles of the intermetallic compound TiAl3. To introduce titanium, an Al-Ti alloy containing 2-5% Ti is used.
An even greater degree of refinement of the macrograins of deformable alloys can be obtained by jointly introducing titanium and boron in the ratio Ti: B = 5: 1. The crystallization centers in this case are complex intermetallic compounds, including compounds TiAl3, TiB2, AlB2 with grain sizes of 2-6 μm. This modification makes it possible to obtain a homogeneous macrostructure with a grain size of 0.2-0.3 mm in ingots with a diameter of more than 500 mm. To introduce titanium and boron, an aluminum-titanium-boron ligature, a “zernolit” preparation, or a flux containing fluoroborate and potassium fluorotitanate are used. The compositions of these modifiers and modification modes are given in table. 14. The highest degree of assimilation of titanium and boron is observed when using flux, which, along with a modifying effect, also has a refining effect.
Modification of the macrostructure of aluminum wrought alloys increases the technological plasticity of ingots and the uniformity of mechanical properties in forgings and stampings.

Casting hypoeutectic and eutectic alloys (AL2, AL4, AL9, AK7, AK9, AL30, AL34) are modified with sodium or strontium to grind eutectic silicon precipitates (see Table 14). Metallic sodium is introduced at 780-800 °C to the bottom of the melt using a bell. Due to the low boiling point (880 °C) and the high chemical activity of sodium, its introduction is associated with some difficulties - large waste of the modifier and gas saturation of the melt, since sodium is stored in kerosene. Therefore, under production conditions, melts are modified with sodium salts.
Modification with a double modifier (a mixture of 67% NaF and 33% NaCl) is carried out at 780-810 °C. The use of a triple modifier (62.5% NaCl, 25% NaF and 12.5% ​​KCl) allows modification to be carried out at 730-750 °C.
To modify, the alloy is poured from the melting furnace into a ladle, which is placed on a heated stand, the metal is heated to the required temperature, the slag is removed, and ground and dehydrated modifier (1-2% by weight of the metal) is poured onto the surface of the melt in an even layer. The melt with applied salts is kept at the modification temperature for 12-15 minutes when using a double modifier and 6-7 minutes when using a triple one. In this case, interaction occurs according to the reaction 6NaF + Al → Na3AlF6 + 3Na. The released sodium has a modifying effect. To speed up the reaction and ensure the diffusion of sodium into the melt, the crust of salts is chopped and kneaded to a depth of 50-100 mm. The resulting slag is thickened by adding fluoride or sodium chloride and removed from the surface of the melt. The quality of modification is controlled by sample fractures and microstructure (Fig. 100). The modified alloy should be poured into molds within 25-30 minutes, since longer exposure is accompanied by the removal of the modification effect.

It is advisable to modify silumins with universal flux (50% NaCl; 30% NaF; 10% KCl; 10% Na3AlF6). Dry powdered flux in an amount of 0.5-1.0% by weight of the melt is poured under the metal stream during pouring from the melting furnace into the ladle. The jet vigorously mixes the flux with the melt. The process is successful if the melt temperature is not lower than 720 °C. When using a universal flux, high temperatures are not required, the melt processing time is reduced, flux consumption is reduced, and the alloy is modified and cleared of metal inclusions.
Modification with sodium does not provide the required duration of preservation of the modification effect and is accompanied by an increase in the susceptibility of alloys to oxidation, the absorption of hydrogen and the formation of gas porosity.
Strontium has good modifying properties. Unlike sodium, this element burns out of aluminum melts more slowly, which allows the modification effect to be maintained for up to 2-3 hours, and does not increase the oxidation of alloys and their tendency to gas absorption to the same extent as sodium. To introduce strontium, an aluminum-strontium alloy with 10% Sr is used. Yttrium and antimony are also used as long-term modifiers.
Hypereutectic silumins (13% Si) crystallize with the release of large silicon particles, which reduce the mechanical properties of the alloys (especially ductility) and complicate mechanical processing due to increased hardness. The grinding of primary silicon crystals is carried out by introducing phosphorus (0.05-0.1%) into the melt - a material surface-active towards silicon (Fig. 101). For modification, the modifiers given in table are used. 14.

Modification

MODIFYING THE STRUCTURE OF CASTINGS AND INGOTS

Using materials from the book Theoretical Fundamentals of Crystallization of Metals and Alloys. Zadiranov A.N., Kats A.M.

1. General ideas about modification

It has been experimentally established that the more nuclei there are per unit volume of the melt, the more crystals are formed, the smaller they are and the higher the mechanical properties of the metal. For this reason, alloys deliberately try to facilitate the formation of crystallization nuclei. The substance that promotes the formation of embryos is called a modifier, and the operation itself is called modification.

Modifiers according to their action can be classified into three groups:

modifiers that increase the wettability of one component of the alloy by another, i.e. reducing the surface tension at the boundary between them and thereby facilitating the formation of the solid phase in contact with the liquid;

1. modifiers, which are direct nuclei of crystallization;

inoculators are modifiers that change the cast structure by reducing the overheating of the crystallizing metal melt.

Modifiers of the second type can be such in very rare cases - when their size and the temperature of the modified metal melt are so close to the solidification temperature that it will not be enough to melt the modifier introduced into the bath and the metal layer that has already crystallized on it (frozen). Particles of the solid phase already present in the melt (non-metallic inclusions or particles of a more refractory metal introduced quite a long time ago, and therefore having the same temperature as the crystallizing melt) cannot be nuclei of the solid phase, since in accordance with the second law of thermodynamics (heat transfer from cold to hot is impossible) they simply cannot take on (into themselves) the heat of crystallization released during the formation of the solid phase. Therefore, statements often found in the literature that oxides, nitrides and sulfides can be crystallization nuclei are very controversial. In addition, the controversial proposition that sulfides and nitrides in steel can be nuclei of the solid phase is caused by the fact that at the moment of the onset of crystallization (temperature 1400...1500 °C) the formation of such compounds is possible only in exotic cases, in particular at very high concentrations of nitrogen and a strong nitride former (for example, zirconium), and in exceptional cases, the release of solid CaS particles is possible when the metal is treated with an excessive amount of calcium at a high sulfur concentration. But even if these inclusions are present in the metal, they have the same temperature as it and therefore cannot accumulate the additional amount of energy released during crystallization in the form of heat of fusion.

Modification is also a widespread technological technique in the production of materials for industries such as astronautics.

Modifiers of the third type - inoculators - exert their effect through cooling the crystallizing metal melt. A higher cooling rate contributes to an increase in the rate of crystallization and a decrease in the development of segregation processes, which, naturally, has a positive effect on the structure.

2. Theoretical foundations of modification

By modifying the macrostructure we mean the production of castings and ingots with a fine-grained structure. The ultimate goal of modification is to improve mechanical, technological and operational properties castings, ingots, as well as products and semi-finished products obtained from them by grinding the cast structure.

The dispersion of the cast structure is characterized by the distance between the axes of the first order or the size of the so-called cast grain. The latter is an area visually distinguished on the thin section, differing from neighboring areas in color and having pronounced boundaries. Cast grains are formed under different thermophysical conditions, the difference in which determines a different direction and possible value of the temperature gradient and, accordingly, the direction of growth of the solid phase; accumulates at the junction of similar areas increased amount liquids and crystal lattice defects, which causes increased etchability of these places and, accordingly, the possibility of their visual identification.

An example of products that require high dispersion of the cast structure is spacecraft.

The cast grain may contain one or more dendrites, the directed growth of which actually contributed to its formation. The grain boundary cannot be crossed by the dendrite itself that formed it. Inside the grain, the axes of the corresponding orders are parallel.

Since the size of the cast grain depends on the ratio of nucleation rates ( n ) and growth ( v ) crystals, then modification is essentially aimed at changing these parameters in the desired direction. The lower the rate of crystal growth and the higher the rate of nucleation of crystallization centers, the smaller the distance between the first-order axes. According to the theory of crystallization, under conditions of spontaneous nucleation of crystals, the rates of their growth and nucleation depend not only on supercooling, but also on the surface tension at the melt-crystal boundary and the activation energy of atoms in the melt ( U )

n= K 1 ·exp[ - U 1 /(R· T)]· exp[ -IN·σ 3 /(T Δ T 2)] (1)
v = K 2 · exp[- U 1 /(R· T)] exp [-E· σ 2 /(TΔ T)] (2)
Where TO 1 - proportionality factor, approximately equal to the number of atoms in the volume of melt under consideration (for one mole K 1~10 23);
K 2 - proportionality factor, approximately equal to the number of atoms on the surface of the volume under consideration (for one mole K 2~10 16);
U- activation energy of atoms in the melt;
U 1 - activation energy, which determines the rate of exchange of atoms between a two-dimensional nucleus and the melt ( U 1 = 0.25 U);
σ - surface tension at the melt-crystal boundary;
σ 1 - surface tension of the melt at the periphery of a two-dimensional embryo;
IN- substance constant = (2 / k)·2;
M andρ - molecular weight and density of the crystal substance;
q- heat of fusion of one mole of a substance;
k - Boltzmann's constant;
E - constant of matter (E σ 2 ~ 10 -3 · IN·σ 3);
R- gas constant;
T - temperature;
Δ T- hypothermia.

From the above equations it follows that an increase in the rates of nucleation and growth of crystals is possible with a decrease in the activation energy and the value of surface tension.

The role of surface tension at the melt-crystal boundary is more clearly visible from the expressions for the total work of nucleation formation ( A r ) and critical radius of the nucleus ( r cr )

A p = Bσ 3 /(T·Δ T 2) (3)
r cr = 2·σ·T/(L·Δ T·T) (4)

The equation for calculating the critical radius of a solid phase nucleus was obtained based on the following considerations.

Education new phase accompanied by the appearance of a new liquid-solid surface. Therefore, in order for a nucleus to form, it is necessary that the decrease in the energy of the mass of matter from which it was formed exceeds the energy expended on the formation of the interface. Therefore, the formation of a new phase (cluster) is possible only when it reaches a certain critical radius. Until the embryo reaches a critical size, its growth is accompanied by an increase in energy. Such a process is possible only due to fluctuations.

Thus, denoting the molar energy of the liquid and solid phases as G L And G S, and the surface of the formed new phase as S, we write down the conditions for the appearance of a new phase

Δ G = V·ρ/ M r·( G S - G L) + S· σ L-S
Where V- volume of one mole of substance, m 3 /mol;
ρ - density of the substance, kg/m3;
M r- molar mass, g/mol;
σ L -S- surface energy, J/m 2.

If we assume that the embryo has a spherical shape, we obtain

Δ G= 4/3·π· r 3 ρ/ M r·( G S - G L) + 4·π· r 2 · σ L-S (5)

At temperatures above the melting point G S > G L and, accordingly, the existence of the solid phase is energetically unfavorable. Cooling the metal to lower temperatures T pl leads to the fact that the difference ( G S - G L) becomes negative. Due to this, in a liquid supercooled to a certain temperature at a certain critical value r = r toΔ value G reaches its maximum value. Further increase r leads to a decrease in Δ G.

The radius of the critical nucleus can be found from the condition that at the maximum ∂ΔG/∂ r= 0. Thus, from equation (5) it follows that

r to= 2∙σ L-S ∙ M r Fe/[(G S - G L)∙ρ Fe]

Magnitude ( G S - G L) can be expressed in terms of the latent heat of fusion and T pl using the known thermodynamic relation:

Δ G = Δ H - T·Δ S = -L - T·Δ S

At T = T pl difference Δ G is equal to zero, therefore

Δ S = -L/Tpl

Assuming that for relatively small supercoolings it does not depend on temperature, we find

Δ G Tmel- Δ G T = (Δ N Tpl - T pl·Δ S Tpl) - (Δ N T - T·Δ S T) = -Δ T·Δ S = Δ T· L/T pl

As a result we get

r to= 2∙σ L-S ∙ M r Fe∙T pl/(ρ Fe∙L∙Δ T)
Where r to- cluster radius, m;
r Fe- radius of the iron atom, Å;
M r Fe- molecular weight of iron, g-atom/mol;
Ρ Fe- density of iron, g/cm 3 ;
σ L-S - surface tension, J/cm 2;
L- heat of fusion, J/mol;
T pl - melting point, K;
Δ T- hypothermia, K.

From these expressions it is clear that the lower the surface tension, the lower the work of nucleation formation and the lower the critical size of a stable nucleus. Thus, a decrease in surface tension at the melt-crystal boundary facilitates the nucleation of crystallization centers, because increases the rate of center nucleation, proportional to the index

y = exp [ -IN· σ 3 /(T· Δ T 2)] (6)

According to these solutions, an increase in supercooling acts in a similar direction, which also contributes to the birth of new crystallization centers. Based on a comparison of equations (1) and (2), we can conclude that of the two processes (nucleation and growth), the process of nucleation of crystallization centers is the limiting one. This is due to the fact that supercooling enters into the nucleation rate equation (1) with a degree of 2 (in contrast to the expression for the growth rate, where the exponent for supercooling is equal to 1). Therefore, the nucleation of crystallization centers requires much greater supercooling than for their growth. Taking this into account, when considering modification, the greatest attention is usually paid to increasing the rate of nucleation of crystallization centers under the influence of modifier impurities.

3. Goals of modification

The modification is aimed at solving a number of problems:

• macrograin grinding;
• grinding micrograins (dendritic cells);

grinding of phase components of eutectics, peritectics, incl. brittle and low-melting phases (with a change in their composition by introducing additives that form chemical compounds with these phases);

grinding of primary crystals that fall out during crystallization in pre- or hypereutectic alloys;

grinding the shape and changing the size and distribution of non-metallic inclusions (intermetallic compounds, carbides, graphite, oxides, sulfides, oxysulfides, nitrides, phosphides).

The simultaneous solution of all these problems often turns out to be impossible. Thus, refinement of the macrostructure is often accompanied by coarsening of micrograins. At the same time, sometimes it is possible to simultaneously achieve several of the listed goals.

Modification differs from alloying:

modifiers have a shorter duration of action (usually 10...15 minutes), but some modifiers have a long-term effect.

4. Modification methods

The following classification of modification methods is proposed:

introducing modifier additives into the melt;

the use of various physical influences (melt temperature regulation, preliminary cooling of the melt during overflow, suspension casting, casting in the crystallization temperature range, vibration, ultrasound, electromagnetic stirring);

combined methods combining the above (input of modifiers + ultrasound, etc.).

5. Types of modifier additives and their effectiveness

At the very beginning of the article it was already said that, according to the nature of their effect, modifiers can be divided into three types: modifiers of the 1st kind, 2nd and 3rd kind. Modifiers of the 1st kind affect the structure by changing the energy characteristics (activation energy and surface tension) of the nucleation of a new phase; modifiers of the 2nd kind, as is believed in most literature sources, change the structure by influencing it as nuclei of the solid phase (however, such an influence of modifiers, in our opinion, is doubtful and subject to revision); modifiers of the 3rd kind - refrigerators / inoculators - reduce the temperature of the metal and increase the rate of crystallization, thereby inhibiting the development of segregation of elements.

5.1. Modifiers of the 1st kind (soluble)

These modifiers are most widely used. Modifiers of this type include impurities that are unlimitedly soluble in the liquid phase and slightly soluble in the solid phase (0.001...0.1%). These impurities, in turn, can be divided into two types: those that do not change the surface properties of the crystallizing phase (a) and those that change the surface tension at the melt-crystal boundary (b). Soluble impurities like " A"can inhibit the growth of the solid phase only due to the concentration barrier at the crystal-melt interface (at a distribution coefficient k < 1 концентрация второго компонента в приграничном слое жидкой фазы выше, чем в твердой фазе). (HOWEVER, DOESN'T THIS ALWAYS REDUCE MOF?) In this case, there is no change in the energy characteristics of the process. Additives like " b", Reducing the surface tension at the melt-crystal boundary and selectively concentrating for this reason on the surface of the crystals (dendrites) are called surface-active.

Surfactants are capable of creating a continuous adsorption layer. This means that with virtually no solubility of the surface-active modifier in the solid phase, a liquid shell enriched with modifier elements is formed around it. In this case, the viscosity of the shell melt can increase significantly ( BUT IS THIS CLEAR?), which, in turn, will reduce the rate of diffusion of atoms to the nucleus

D=k· T/(4 · n · n · r M) (8)
Where D- diffusion coefficient;
k- Boltzmann's constant;
T - melt temperature;
n - coefficient of dynamic viscosity;
rm - radius of the modifier atom.

As the flow of atoms to the nucleus decreases, crystal growth becomes more difficult.

The formation of such a layer enriched with an impurity/modifier in front of the crystallization front under conditions of ongoing heat removal leads to an increase in supercooling in the liquid layer ahead of the crystallization front.

The effect of additives like " b"based on a decrease in the value of surface tension σ at the melt-crystal boundary. Such additives (impurities) are called surface-active to the crystallizing phase. They reduce the temperature range of metastability (minimum supercooling, exceeding which ensures the emergence of crystallization centers). The tendency to adsorption is determined by the generalized ratio (moment) of the charge of an ion to its crystallographic radius. If the generalized moment of the ion of a surfactant is less than the generalized moment of the metal, then this additive will lower the surface tension.

The complexity of the action of soluble surfactant impurities is due to the fact that, along with a change in surface tension σ, they can change the activation energy U. Impurities that are soluble in the liquid phase and insoluble in the solid phase, during crystal growth, create, as noted above, an increased concentration in the liquid layer adjacent to the growing crystals. Thus, they prevent the growth of crystals and increase the activation energy necessary for the exchange of atoms between the liquid and solid phases. Therefore, usually a surface-active impurity, along with a decrease in surface tension, which accelerates the nucleation of centers, increases the activation energy, is adsorbed on the surface of growing crystals, and complicates the transition of atoms from the liquid phase to the solid phase. In this case, an increase in the activation energy slows down the generation of new centers and reduces the rate of their growth.

Thus, the introduction of modifiers of the 1st kind is accompanied by a change in surface tension and activation energy in opposite directions. This is complicated by their combined influence on crystallization and the size of the cast grain. From expression (1) it is clear that the exponent (3) at σ higher than with U(1), therefore we can expect a stronger influence on the nucleation rate of surface tension. Thus, the effect of macrograin refinement is most characteristic of modifiers of the 1st type. Since an increase in activation energy due to the adsorption of impurities on crystal faces helps to reduce the rate of crystal growth, this causes a coarsening of the dendritic structure of the grain. Thus, under the influence of modifiers of the 1st kind, the macrograin is simultaneously refined and the micrograin is enlarged, i.e. there is a complex effect on the macro- and microstructure.

The above mechanism of action of modifiers of this type was confirmed in experimental studies when studying the modification of high-alloy steels with magnesium, boron, cerium, and barium. At the same time, a decrease in the surface tension of the metal and its tendency to overcooling was revealed when additives were introduced. The minimum value of the surface tension of the modified metal corresponded to the smallest grain size.

Examples of modifiers of the 1st kind are given in table. 1. More detailed data on rational modifiers and their content in relation to various grades of steel are given in table. 2.

Of interest are the data on the simultaneous reduction in the sizes of macro- and micrograins when modifying steel with small additives, as well as the facts of the disappearance of the dendritic structure when 0.3% zirconium is introduced into the steel (only small austenite micrograins are detected). A simultaneous decrease in the sizes of macro- and micrograins in X25N20 steel upon modification with zirconium was established. Blocking of the dendritic form of crystal growth in steel and difficulties in the growth of macrograins were noted at a sufficiently high concentration of the surfactant additive.

Table 1. Modifiers of the 1st kind for various metals and alloys.

Metal (alloy)	Modifier	Note
Steel	Boron, rare earth metals, calcium cerium, magnesium, lanthanum, zirconium, lithium, barium, uranium
Aluminum and aluminum alloys with silicon (silumins) (AL2, AL4, AL9, AK9, etc.)	Sodium (0.006-0.012%), potassium, lithium, bismuth, antimony 0.1-0.3%, strontium 0.01-0.05% (antimony and strontium are long-acting modifiers), mixture of salts (0.1% sodium and 2% mixture of sodium fluoride and sodium chloride)	Hypothermia 6-15°C. Grinding of eutectic in the Al-Si system with sodium, strontium. The lamellar shape of silicon crystals becomes compact with a size of 2-5 microns
Copper Copper alloys without iron Copper alloys with iron	Tin, antimony Vanadium, zirconium, molybdenum Titanium, boron, tungsten
Cast iron	Scandium, lanthanum
High-strength cast iron with nodular graphite	Primary modification with hundredths of magnesium or cerium plus secondary (graphitizing) modification with ferrosilicon FS75 to prevent the appearance of structurally free carbides in cast iron	Conversion of lamellar graphite precipitates of iron-graphite eutectic into spherical particles
Malleable cast iron, heat treatable	Thousandths of a percent of bismuth, antimony or tin
Magnesium alloys containing aluminum	Carbon-containing substances (0.3-0.6%), ferric chloride, chalk, marble, magnesite, hexachloroethane, carbon dioxide, acetylene. Melt overheating-holding-cooling
Magnesium alloys not containing aluminum	Zirconium 0.5-0.7% or calcium 0.1-0.2%

Table 2. Modifiers for various grades of steel

steel grade	Modifiers	Amount of additive in %
20L	Titanium	0,75
U12	Cerium	0,50
U12	Titanium	0,25
40HL	Titanium	0,50
ZOHNZM	Bor	0,50
1X1 8H9	Titanium	0,50
1X1 8H9	Zirconium	0,25

The effect of modification is different for different grades of steel (Table 3).

5.2. Modifiers of the 2nd kind (insoluble)

Also, the crystallization parameters and the macrostructure reflecting it can be influenced by solid particles introduced into the melt. At the same time, a number of researchers associate this influence precisely with the contact effect on the process of nucleation of crystallization centers. This is explained by the fact that when an insoluble impurity with properties close to the properties of the crystallizing substance is introduced into the melt, a significant decrease in the metastability range of the melt occurs. This position is based on the so-called principle of P.D. Dankov, according to which heterogeneous nucleation is caused by insoluble impurities that are structurally similar to the crystallizing substance. Such impurities are called isomorphic with the crystallizing substance and modifiers of the 2nd kind. They have crystal lattice parameters close to the parameters of this substance, and it is believed that, similar to type 1 modifiers, they provide a reduction in the metastability range and refinement of the macrograin. Isomorphic impurities are those whose lattice periods differ from the lattice period of the crystallizing metal by no more than 10...15%. The usual content of modifiers of this type is less than 0.1%. It is considered necessary that the centers of crystallization be released in a very dispersed form (no more than 1 micron), thereby forming a stable suspension that is not prone to coagulation and stratification during prolonged exposure of the melt in the mixer and during the casting process.

Based on a generalization of various works, the following conditions are formulated for the selection of insoluble additives (particles) with the greatest modifying ability:

it is necessary to use refractory insoluble substances that form an independent phase in the melt;

particles of the solid phase must comply to the maximum extent with the principle of structural and dimensional correspondence;

dispersed particles with a large total phase interface and comparable in size to clusters of the order of 1...10 nm are more effective;

it is desirable that the particles have metallic properties (type of chemical bond);

The most effective are particles of stable chemical compounds of endogenous origin, i.e. formed in the melt as a result of the interaction of the additive with one of the components or base of the alloy;

in most cases, effective additives form intermetallic compounds and eutectic (or peritectic) with the alloy base with the eutectic point strongly shifted towards the base component.

Examples of modifiers of the 2nd kind are given in table. 4.

Table 4. Modifiers of the 2nd kind

Metal (alloy)	Modifier	Note
Aluminum alloys	Sodium chloride, titanium - up to 0.1 5%, vanadium - up to 0.15%, scandium, zirconium, boron	Refractory compounds are formed that are isomorphic to aluminum: TiAl 3, ScAl 3, VAl 6, ZrAl 3, TiB 2
Hypereutectic silumins	Phosphorus 0.05-0.1% or sulfur	Introduction of crystallization centers (aluminum phosphide AlP), grinding of primary silicon
Become	Aluminum, titanium	Refractory compounds Al 2 O 3 and TiN are formed
Gray cast iron with flake graphite	Graphitizing modifier - silicon; stabilizing modifiers - manganese, chromium, tin, copper, antimony, etc.	Input of silicocalcium SK30 (0.3-0.6%) or ferrosilicon FS75 (0.5-0.8% by weight of cast iron). Purpose: grinding graphite and reducing the tendency of cast iron to chill

A number of researchers believe that modifiers of the 2nd kind can also be formed from modifiers of the 1st kind. Thus, the nature of the action of type 1 modifiers, for example boron in steel, can change when chemical compounds of the modifier are formed with other elements. In this case, the new chemical compound will ultimately play the role of an independent modifier. Under some conditions, these compounds can be surface-active, and under others, on the contrary, inactive (not reducing, but increasing surface tension). Thus, boron in steel can form a stable chemical compound with iron FeB 2, which will serve as a crystallization center as a type 2 modifier (insoluble impurity). When introducing aluminum into steel, the formation (along with the compounds noted in Table 4) of aluminum nitrides is possible, which will also create crystallization centers.

When modifying gray cast iron with silicon in order to obtain cast iron with flake graphite, a “silicate turbidity” is formed in the melt (silicon, which is a graphitizer, contributes to the appearance of graphite spel - graphitization centers). At the same time, chill is eliminated and the structure is refined (small graphite plates are formed). At the same time, the number of graphite inclusions decreases and the mechanical properties and their uniformity increase, ensuring high wear resistance and machinability of cast products. The best modification results are achieved with a reduced content of silicon and carbon in the original gray cast iron.

Modification with additives that promote the appearance of crystallization centers is accompanied by a decrease in supercooling (in contrast to modification with surface-active additives adsorbed on the surface of growing crystals).

5.3. Activated (active) impurities (insoluble)

This type of impurities differs in that they enter the melt with the charge, which has undergone preliminary crystallization (for example, pigs). These impurities do not have a structural similarity to the crystallizing substance, but during previous crystallization they acquire the property of becoming a solid substrate for nascent crystal nuclei. This is due to the fact that in the previous crystallization between the particles of this impurity and the solid phase of the crystallized substance, a boundary layer appears due to molecular contact, which has a structural similarity to the crystals of the substance. If the melting temperature of this layer is higher than the melting temperature of the crystals, then after their melting, the marked boundary (transition) layer will remain on the surface of the impurity particles. This will ensure the transformation of these impurity particles into active ones. As a result, it is possible that their grinding effect on macrograins is similar to type 2 modifiers. It is believed that technical metals and alloys always contain active impurities that significantly affect the nature of crystallization and the formation of the macrostructure of castings and ingots.

The manifestation of the role of active impurities is usually associated with the effect of melt overheating on the macrostructure. An increase in the superheating temperature of the melt, as a rule, leads to an enlargement of the macrostructure. This is explained by the deactivation of active impurities, which is understood as the disappearance of the activated layer melt on the surface of impurity particles at an elevated temperature. The decontamination temperature depends on the type (nature) of impurities and the composition of the melt. In particular, complete deactivation of impurities occurs upon overheating (by °C): steel X27 - by 100, steel 12Kh18N9T - by 5-10, steel X23N18 - by 2-3, aluminum - by 50-60. At sufficiently low overheating of the melt, the effect of structure inheritance was observed, also associated with the action of active impurities. Under these conditions, charge billets having a fine-grained structure hereditarily transfer the corresponding fine macrostructure to the castings or ingots cast from them. However, this heredity effect disappeared at higher melt superheat (°C), for example, for aluminum - above 8-10, and for steel 15X28 - above 30-40.

In the last 10-15 years, work has been developing in the field of so-called genetic engineering, aimed at controlling the structure and properties of castings and ingots using the phenomenon of heredity.

5.4. Complex modifiers

The use of complex modifiers is dictated by several reasons:

the combined effect of two or more modifiers enhances the effect obtained when using a single modifier. This is due to the above-mentioned nucleation of crystallization centers on insoluble impurities in the layer of the liquid phase with diffusion supercooling caused by the introduction of a soluble impurity (especially a surface-active one);

when using a complex modifier, it is possible to minimize the content of each of its components, which makes it easier to meet the conditions for limiting the alloy composition by impurities;

the combination of modifiers with physical influences enhances the effect of the modifiers and creates the possibility of obtaining particularly small and special structures.

There are three types of complex modifiers:

• refining, containing active elements Mn, Si, Ca, Mg, Al, P3M, etc.;
• strengthening agents containing carbides, borides, nitrides, which are formed in the alloy as a result of the interaction of the corresponding elements and contribute to dispersion strengthening of the base;
• refining and strengthening agents, which contain active elements and compounds.

Modifiers containing active elements such as REM, Ba, Ca are an effective means of changing the nature and shape of non-metallic inclusions, obtaining the most preferred type of oxide inclusions in the sulfide shell.

Modification of steel with nitrides of vanadium, titanium, zirconium, and aluminum by introducing special alloys or nitrided ferroalloys into the steel has become common. As a result, nitride and carbonitride dispersed particles are released in steel during quenching and subsequent tempering. When modifying structural steels with vanadium nitrides, the austenite grain is refined by 3-4 points, increasing ductility, toughness and strength.

Table 5 Complex modifiers.

Metal (alloy)	Modifier	Note
Gray cast iron	Fe-Si-Ca + Al, Ti, Ce, La
Gray cast iron with reduced carbon equivalent C + 0.3 Si (3.5-3.7)	Si-Mn-Zr	Goal: obtaining cast iron with flake graphite
Steel	REM with silicocalcium, ferromanganese, ferrosilicon	Purpose: removal of cerium sulfides
Steel	Ti-B-Ca, Ti-Ce-B, Mg-Zr-Ce, Ti-V-Ca
Aluminum alloys	Ti+B. mixtures of chloride and fluoride salts (sodium chloride, sodium fluoride, potassium chloride and cryolite)	Goal: obtaining more dispersed and stable intermetallic compounds
Malleable iron	Al+Bi+B	Goal: reducing annealing time
Ductile iron with vermicular graphite	Mg + Ti, Y, Ce, Ca	Goal: obtaining isolated, thickened inclusions with rounded ends (more compact than flake graphite)

A feature of modifying steel with complex alloys is that, in parallel with the refinement of the structure, the nature and shape of non-metallic inclusions change, the contamination of austenite grain boundaries with oxide, sulfide and nitride inclusions decreases by 1.5-2.0 times, and the uniformity of distribution increases structural components, provides an increase in the ductility and toughness of steel.

In the production of nodular cast iron, along with separate modifiers (magnesium or cerium), a complex modifier (magnesium + cerium) is used. The addition of cerium to magnesium neutralizes the effect of harmful impurities (titanium, aluminum, lead, antimony, arsenic, bismuth, tin), which have an extremely harmful effect on the quality of cast iron modified with magnesium. Examples of complex modifiers are given in table. 5.

5.5. Modifiers of the 3rd kind - inoculators

The introduction of inoculants into the crystallizing melt ensures an increase in the homogeneity and dispersion of the cast structure, optimization of the shape and distribution of non-metallic inclusions, a reduction in some casting defects (porosity, friability, axial and extra-axial segregation), which significantly increases the level and isotropy of the properties of the cast metal:

with approximately equal strength, the plastic characteristics of the metal and its impact strength increase by 30...50% or more (up to 2.5...3.0 times);

the maximum effect of increasing plastic properties in the middle (at half the radius) and axial zones indicates a significant increase in the physicochemical homogeneity and isotropy of the properties of the metal over the cross section of the ingots;

the decrease in the anisotropy of the properties of the suspension metal in the longitudinal direction in the surface zone is associated with the elimination of the structure of columnar crystallites, which is usually characteristic of this region.

an increase in the level and isotropy of the characteristics of ductility and toughness of steel, due to the introduction of powders, is maintained after forging (up to 5...10-fold forging);

In terms of ductility, ingots cast with the introduction of exogenous inoculants approach this indicator of forged metal or reach maximum values ​​already at small 1.5- and 3-fold forging, impact strength does not decrease after 5...10-fold forging, as it does place in ordinary ingots.

However, despite the improvement in the macrostructure of ingots and castings, the use of metal powder and cast shot as inoculants leads to an increase in steel contamination with non-metallic inclusions, mainly oxides. The limited use of this technology is caused by the complexity of the technological chain for producing dispersed inoculants (powder, shot), which require protection from oxidation during storage, transportation and input into the ingot. In addition, the existing methods and devices for processing liquid steel with dispersed inoculators have not been widely implemented due to insufficiently developed input technology, complexity of operation and a number of design flaws.

A promising direction in the field of improving the technology for introducing inoculators and controlling the metal structure is the method of forming inoculators in a jet when casting large ingots in a vacuum. With this casting method, proposed by S.I. Zhuliev, the introduced particles have the same chemical composition as the melt. The formation of solid particles in this case is ensured by additional separation of the melt stream with the creation of conditions for the crystallization of droplets as they enter the mold.

Getting into the metal, the inoculators lead to local cooling of the metal melt, while first a crust of the solid phase freezes on them, which subsequently melts due to heating from the surrounding melt, and later the inoculator itself melts. Thus, the inoculators in the melt take away heat for their own heating and melting, resulting in a decrease in the temperature of the melt. The cooling effect it introduces ultimately leads to an increase in the rate of crystallization, which in turn is reflected in a decrease in segregation heterogeneity in the workpiece and an increase in the uniformity of mechanical properties in large forged products for critical purposes. As the mass of introduced inoculants increases, the rate of crystallization increases.

6. Generalized systematization of modifiers

Previously, modifiers were systematized on the basis of the periodic system of D.I. Mendeleev. In the upper part of the diagram, the curve of changes in their melting temperatures was used as a characteristic of the periodicity of changes in the properties of simple bodies. At the bottom of the diagram, high bars indicate elements that give a strong modification effect in steel, cast iron and aluminum alloys, and low bars indicate a weak effect. The absence of a column against the element number meant that it was not a modifier. Shaded bars corresponded to reliably established data, unshaded bars corresponded to questionable data or the absence of data on the proposed effect.

The position of modifier elements on the diagram in most cases corresponded to the first elements of each period, marked by a double line on the melting point curve. The results of the above systematization showed the existence of a direct connection between the structure of the outer electron shells of atoms of elements and their modifying effect. This is consistent with the influence on surface tension of the generalized ratio (moment) of the charge of the surface-active additive ion to its crystallographic radius (compared to the corresponding characteristic for the base metal).

7. Processes occurring during modification

Usually, without modification, the value of supercooling of non-ferrous metals and alloys reaches 7-10 °C. As a rule, during modification, a large number of crystallization centers appear in the melt. As a result, the heat of crystallization is released and supercooling almost disappears. Further growth of crystallization centers depends on the nature of the influence of impurities or physical influences on the situation in the crystal-melt boundary zone. In most cases, soluble or insoluble impurities have an inhibitory effect on crystal growth, and the specific mechanism of growth inhibition depends on the nature of the impurity and the mechanism of its modifying action.

When modifying iron by introducing 0.1% cerium and lanthanum, supercooling decreased from 320 to 40-50 °C, and when introducing rare-earth metals into steel - from 260 to 10-30 °C. At the same time, during the refining action of modifiers (removal of non-metallic inclusions), greater supercooling was observed compared to the unmodified melt. The modifying role of rare earth metals manifests itself only with slight overheating of the steel and in a short period of time. A similar picture occurs when modifying non-ferrous metals and alloys. Therefore, they tend to carry out modification immediately before pouring the melt or introduce modifiers directly into the stream of the poured melt.

8. Results of the influence of modifiers on the structure

Two types of effects of modifiers (Fig. 1) on the structure were established:

monotonous grain grinding with increasing modifier content. At very low concentrations, the influence of the modifier is insignificant, and at concentrations of more than 0.2-0.6% it stabilizes, so usually the content of modifiers is 0.1-0.3%;

non-monotonic grain grinding with an optimal concentration region of 0.01-0.1%, exceeding which leads to an increase in grain size.

The option of a monotonic decrease in grain size with an increase in modifier concentration is characteristic of insoluble catalyst impurities (for example, titanium in aluminum), and the option of non-monotonic grain refinement is characteristic of surface-active soluble impurities (for example, magnesium in zinc).

Rice. 1. Scheme of the influence of the modifier content on the size of the alloy macrograin:
1 - monotonous grain grinding; 2 - non-monotonic grain grinding.

Rice. 2. The influence of modifiers on the structural components of alloys.

The effect of modifiers on individual structural components of the alloy is schematically depicted in Table. 6 and in Fig. 2. It was found that the addition of 0.08% boron to Kh15N25L steel reduces the macrograin size from 9 to 2 mm.

Table 6 Results of the influence of modifiers on the structure.

Alloy type	Result	Structures (Fig. 2)
Alloys - solid solutions (carbon steels with ferrite-pearlite structure)	Primary grain grinding	1,2
-"-	Phase recrystallization	1,3
-"-	Grinding of secondary grains after phase recrystallization	1,2,4
Alloys with primary precipitates and eutectic (gray and high-strength cast iron)	Grinding of both structural components	5,6
-"-	Coarse-crystalline eutectic	7
-"-	Thin-plate eutectic with very short plates	8
-"-	Grinding of individual large structural components	9,10
-"-	Coagulation and spheroidization of structural components	11,12

Along with the concept of “modifiers”, there is the opposite concept of “demodifiers” - additives that increase grain size. They increase the work of embryo formation, delay its formation and reduce the probability of the occurrence of a crystallization center. Demodifiers include: bismuth, lead, antimony - for cast iron; sulfur and carbon - for magnetic alloys of the Fe-Ni-Co-Al-Cu-Ti system (increase the size of columnar crystals).

9. Influence of modification on the properties of castings and ingots

Modification increases the mechanical properties of castings and ingots (Table 7). It was found that in cast iron and silumin the positive effect of the modifiers has a particularly strong effect on the plastic characteristics of the cast metal.

Table 7. Results of the influence of modifiers on mechanical properties.

Metal (alloy)	Result of influence on properties
Nodular cast iron compared to gray cast iron with flake graphite	Increase in tensile strength by 2-4 times, and elongation by tens of times
Steel	Increased strength by 25-30%, wear resistance by 15-50%, heat resistance by up to 45%, ductility, impact strength
Silumin	Increase in tensile strength by 1.14-1.55 times and relative elongation by 2.2-6.5 Abramov V.P., Zatulovsky S.S., Mayorov N.P. et al. Uniformity of a continuous carbon steel ingot after suspension casting // Problems of steel ingot: Tr. IV Conference on Ingots. M.: Metallurgy. 1969. P. 497...499. Skvortsov A. A., Sokolov L. A., Ulyanov V. A. On the use of water-cooled vibration coolers for continuous casting of steel // Izv. Academy of Sciences of the USSR. Metals. 1980. No. 1. pp. 61...65. Kutishchev S. M. Features of casting steel ingots with a cooling inoculator // Physico-chemical influence on the crystallization of steel: Sat. scientific tr. Kyiv: IPL AN Ukrainian SSR. 1982. P. 121...126. Zatulovsky S.S. Suspension casting. Kyiv: Naukova Dumka, 1981. 260 p.