News

Austenitic Manganese Steel

• satyendra
• April 3, 2014
• 53 Comments
• abrasion, AMS, austenitic, Casting, corrosion, embrittlement, Heat teatment, manganese steel, Thermal expansion, Wear resistance, Work hardening,

Austenitic Manganese Steel

It was the need for an alloy which combines hardness and toughness which had motivated Sheffield Robert Abbott Hadfield to embark on the study of alloys of iron (Fe) and other elements in 1878. Four years later, he found that with 10 % manganese (Mn), heat treatment, and water quenching, the materials have both hardness and toughness. The original austenitic manganese steel (AMS), containing 1.2 % carbon (C) and 12 % Mn, invented by Hadfield still holds prominence to-date without much change in its invented composition. Hadfield’s steel was unique since it combined high toughness and ductility with high work hardening capacity and with good resistance to combat wear at service temperatures up to 250 deg C.

AMS offers the best combination of toughness and resistance to high stress and gouging abrasion. Hence, it rapidly gained acceptance as a very useful engineering material. AMS is still used extensively, with minor modifications in composition and heat treatment, in several fields which include (i) equipment for handling and processing earthen materials (rock crushers, grinding mills, dredge buckets, power shovels, bucket teeth, tooth adaptors, pans, pumps handling gravel and rocks, clay crusher rolls, grizzly bar, wear pads liners, beaters, ring-granulators, blow bars, and hammers), (ii) automobile industry (fragmentation hammers, shredders, and grates), (iii) military tank track pads, (iv) railway track work (frogs, switches, and crossings where wheel impacts at inter-sections are especially severe), (v) metal-to-metal wear also known as galling (sprockets, pinions, gears, wheels, sheave wheels, conveyor chains, drag-line chain, wear plates, shoes), and in steel plant (weld overlays on couplings, spindles, pinions and other items working under heavy impact loads). Hard to cut AMS also finds its end use in safes, and prison cell bars. AMS also develops a favourable wear pair with alloy steels so that it can be used a bushing material in demanding mining applications subjected to galling type of wear.

Recently, new light weight alloys with very high Mn and aluminium (Al) contents for armour plates applications are being developed. Because of its austenitic nature, AMS has no magnetic response and hence, makes good wear plates for the bottom of electro-magnets. AMS use is understood to be growing in liquefied natural gas (LNG) applications, and cryogenic LNG transportation equipment and vessels.

AMS has certain properties which tend to restrict its use. It is difficult to machine and normally has a yield strength of only 345 MPa to 415 MPa. Hence, it is not well suited for parts which need close-tolerance machining or which is to resist plastic deformation when highly stressed in service. However, hammering, pressing, cold rolling, or explosion shocking of the surface raises the yield strength to provide a hard surface on a tough core structure. Fig 1 shows microstructures of austenitic manganese steel.

Fig 1 Microstructures of austenitic manganese steel

Several variations of the chemical composition of the original AMS have been proposed, frequently in unexploited patents, but only a few have been adopted as significant improvements. These normally involve variations of C and Mn, with or without additional alloys such as chromium (Cr), nickel (Ni), molybdenum (Mo), vanadium (V), titanium (Ti), and bismuth (Bi). The compositions are available in different national and international standards.

The available assortment of wrought grades is smaller. Some wrought grades contain around 0.8 % C and either 3 % Ni or 1 % Mo. Large heat orders are normally needed for the production of wrought grades, while cast grades and their modifications are more easily available in small lots. A Mn steel foundry can have several dozen modified grades on its production list. Modified grades are normally produced to meet the needs of application, section size, casting size, cost, and weldability considerations.

Mn steels are simple and cheap to produce and offer excellent potential to replace expensive chromium iron alloy, which are known to possess high hardness and wear resistance but also to show critical limitations regarding ductility and toughness. The high C content in Mn steels can be completely retained in solution, which provides best resistance to abrasion. Mn steels normally show freezing ranges as wide as 200 deg C (temperature range between liquidus and solidus lines), making them susceptible to micro-porosity and the occurrence of harmful continuous carbide networks, particularly at grain boundaries.

During industrial production of Mn steel like several other alloys steel, there is always slight, unavoidable variation in composition for reason ranging from prolong holding of liquid bath at high temperature, and heavy oxidation characteristic of alloying elements in the shop floor etc. From the Mn steel equilibrium diagram, a C composition higher than 1.2 % encourages the formation of acicular carbide which in turn can lead to intergranular embrittlement in steel. Different elements present in the AMS and how they influence the AMS is described below.

Carbon and manganese – The compositions of AMS do not permit any austenite transformation when the alloys are water quenched from above the Acm (the temperature which corresponds to the boundary between the cementite-austenite and the austenite regions). However, this does not preclude lower ductility in heavy sections because of slower quenching rates. The effect is because of the formation of carbides along grain boundaries and other inter-dendritic areas and to some extent affects nearly all commercial castings except the very small ones. Fig 2a shows Acm temperatures for 13 % Mn steels containing between 0.6 % C and 1.4 % C. Fig 2b shows the effects of C and Mn content on the Ms temperature, i.e., the temperature at which martensite starts to form from austenite upon cooling, of a homogeneous austenite with all C and Mn in solid solution.

Fig 2 Acm temperatures and variation in Ms temperatures

The mechanical properties of AMS vary with both C and Mn content. It is seen that C increases strength up to the range of 0.4 % C to 1.6 %. A plateau is indicated in the curve at 1.05 % C to 1.35 % C content. Any departure from the curve can be attributed to grain size unless good statistical evidence is found. The plateau is at 827 MPa. As C is increased, it becomes increasingly difficult to retain all of the C in solid solution, which can account for reductions in tensile strength and ductility. However, since abrasion resistance tends to increase with C, C content higher than 1.20 % can be preferred even when ductility is lowered. C content above 1.4 % is rarely used because of the difficulty of achieving an austenitic structure sufficiently free of grain boundary carbides, which are detrimental to strength and ductility. The effect can also be observed in 13 % Mn steels containing less than 1.4 % C since segregation can result in local variations of +/- 17 % (+/- 0.2 % C) from the average C level determined by chemical analysis.

The 0.7% C (minimum) of certain grades can be used for minimizing carbide precipitation in heavy castings or in weldments, and similar low C contents are specified for welding filler metal. Carbides are formed in castings which are cooled slowly in the moulds. In fact, carbides is formed in practically all as cast grades containing more than 1.0 % C, regardless of mould cooling rates. They are formed in heavy-section castings during heat treatment if quenching is ineffective in producing rapid cooling throughout the entire section thickness. Carbides can form during welding or during service at temperatures above around 275 deg C. If C and Mn are lowered together, e.g., to 0.53 % C with 8.3 % Mn or 0.62 % C with 8.1 % Mn, the work-hardening rate is increased because of the formation of strain-induced alpha (body-centered cubic, or bcc) martensite. However, this does not provide improved abrasion resistance (at least to high-stress grinding abrasion) as is frequently desired.

Manganese contributes the important austenite-stabilizing effect of delaying transformation (but not eliminating it.). Hence, in a simple steel which contains 1.1 % Mn, isothermal transformation at 370 deg C begins around 15 seconds after the steel is quenched to that temperature, whereas in a 13 % Mn steel, transformation at the same temperature does not begin until after 48 hours. Below 260 deg C, phase changes and carbide precipitation are so sluggish that for all practical purposes they can be neglected, in the absence of deformation, if Mn content exceeds 10 %. Fig 3 shows the influence of Mn content on the strength and ductility of cast austenitic steel which has been solution treated and water quenched. It confirms the observations of several studies done to find out the influence of Mn content up to around 22 %. Mn content has little effect on yield strength. In tensile testing, ultimate strength and ductility increase fairly rapidly with increasing Mn content up to around 12 % and then tend to level off, although small improvements normally continue up to around 13 % Mn.

Fig 3 Variation of properties with manganese content

Silicon and phosphorus – Silicon (Si) and phosphorus (P) are present in all the grades of AMS. Si is seldom added except for steelmaking purposes. Si content exceeding 1 % is not common, since foundries do not like to have the Si pyramid in melts containing returned scrap. A Si content of 1 % to 2 % can be used to increase yield strength to a moderate degree, but other elements are preferred for this effect. Loss of strength is abrupt above 2.2 % Si, and Mn steel containing more than 2.3 % Si can be of no use. On the other hand, Si levels below 0.1 % show decreased fluidity during casting.

The availability of low-P ferro-manganese since around 1960 has enabled steelmakers to reduce P levels in Mn steel to a large extent. The preferred practice is to hold the P content below 0.04 % even though 0.07 % is permitted in standards. Levels above 0.06 % contribute to hot shortness and low elongation at very high temperatures and frequently are the cause of hot tears in castings and under-bead cracking in weldments. It is particularly advantageous to keep P at the lowest possible level in the AMS grades which are welded, and in Mn steel welding electrodes, and in heavy section castings.

Common alloy modifications – The most common alloying elements are Cr, Mo, and Ni. When added to the normal C level of around 1.15 %, both Cr and Mo increase yield strength and flow resistance under impact.

Chromium additions are less expensive for a given increase, and Cr grades are probably the most common modifications. Normal grade of AMS frequently contains some Cr also. The 2 % Cr does not significantly lower toughness in light sections. However, in heavier sections, its effect is similar to that of raising the C level, the result is a decrease in ductility because of an increase in the volume fraction of carbides in the micro-structure. Cr additions have been used up to 6 % for some applications, sometimes in combination with copper (Cu), but these grades no longer receive much attention. Cr improves resistance to both atmospheric corrosion and abrasion, although the latter effect is not always consistent and depends on the individual application. It is also used up to 18 % in low-carbon electrodes for welding Mn steel. Because of the stabilizing effect of Cr on iron carbide, higher heat-treatment (solutionizing) temperatures are frequently necessary prior to water quenching.

Molybdenum additions, normally in the range of 0.5 % to 2 %, are made to improve the toughness and resistance to cracking of castings in the as-cast condition and to raise the yield strength (and possibly toughness) of heavy-section castings in the solution treated and quenched condition. These effects occur since Mo in Mn-steel is distributed partly in solution in the austenite and partly in primary carbides formed during solidification of the steel. The Mo in solution effectively suppresses the formation of both embrittling carbide precipitates and pearlite, even when the austenite is exposed to temperatures above 275 deg C during welding or in service.

The Mo in primary carbides tends to change the morphology from continuous envelopes around austenite dendrites to a less harmful nodular form, especially when the Mo content exceeds 1.5 %. The 1 % Mo grades are resistant to the reheating effect which limits the usefulness of this grade. Grade with Mo content of 0.9 % to 1.2 % is adapted to heavy-section castings used in roll and impact crushers which are frequently reheated during weld build-up and overlays. Grade containing around 2 % Mo can be given a special heat treatment for developing a structure of finely dispersed carbides in austenite. This heat treatment involves a partial grain refinement by pearlitizing near 595 deg C for 12 hours and water quenching from 980 deg C. This type of micro-structure has been found to improve abrasion resistance in crusher applications.

The tensile properties of samples taken from worn cone crusher parts show yield strength in the range of 440 MPa to 485 MPa, tensile strength in the range of 695 MPa to 850 MPa, and elongation in the range of 15 % to 25 %. The addition of Mo in quantities higher than 1 % can increase the susceptibility of the Mn steel to incipient fusion during heat treatment. Incipient melting is a liquation phenomenon which occurs because of the presence of low-melting constituents in inter-dendritic areas, both within individual grains and along grain boundaries. This tendency is aggravated by higher P levels (higher than 0.05 %), higher pouring temperatures (which promote segregation in the casting), and higher C levels (higher than 1.3 %) in the steel. As a further use, Mo is added to the lean Mn steel grade (6 % to 8 % Mn) partly to suppress embrittlement in both as-cast and heat-treated conditions.

Nickel, in quantities up to 4 %, stabilizes the austenite since it remains in solid solution. It is particularly effective for suppressing precipitates of carbide platelets, which can form between around 300 deg C to 550 deg C. Hence, the presence of Ni helps retain non-magnetic qualities in the steel, especially in the decarburized surface layers. Ni additions increase ductility, decrease yield strength slightly, and lower the abrasion resistance of Mn steel. Ni is used mainly in the lower-C or weldable grades of cast Mn steel and in wrought Mn steel products (including welding electrodes). In wrought steel products, Ni is sometimes used in conjunction with Mo.

Manganese in AMS is normally in the range of 11.5 % to 14 %. One grade of AMS has reduced Mn (6 % to 8 %) to make the austenite less stable, but this needs compensation with 1 % Mo to gain acceptable properties. Work-hardening rates are reported to be higher than that of the standard 13 % Mn grades, with some loss in toughness. This grade has been used in scoop lips, ball mill end liners, discharge grates, and grizzly screens for siliceous ore grinding. In one case, 45 % longer life in ball mill discharge grates is reported compared to pearlitic Cr-Mo steel used earlier. Average properties reported are 415 MPa yield strength, 585 MPa tensile strength, and 12 % elongation for C levels of 1.2% to 1.4 %. This grade is not adapted to heavy sections or to service involving temperatures above 315 deg C. It has poor weldability and is to be avoided if a casting is to be hard-faced or rebuilt.

Several other elements, e.g., V, Ti, Bi, and Cu are added to Mn steel for unique applications. V is a strong carbide former, and its addition to Mn steels considerably increases yield strength, but with a corresponding decrease in ductility. V is used in precipitation-hardening Mn steels in quantities ranging from 0.5 % to 2 %. Because of the stability of vanadium carbonitrides (VCN), a higher solutionizing temperature 1,120 deg C to 1,175 deg C is desired prior to aging (normally between 500 deg C to 650 deg C. Yield strengths of over 700 MPa are achievable depending on the level of ductility which can be tolerated for a given application. Tests of an age-hardened Mn-Ni-Mo-V austenitic alloy demonstrated that the abrasion resistance of this steel is not as good as that of the standard grades.

Ni, Cu in quantities of 1 % to 5 % has been used in AMS for stabilizing the austenite. The effects of Cu on mechanical properties have not been clearly established. Scattered reports indicate that it can have an embrittling effect, which can be because of the limited solubility of Cu in austenite.

Other elements such as Bi and Ti are also added to standard Mn steels. Bi has been found to improve machinability, especially when coupled with higher Mn levels (higher than 13 %). Ti can reduce C in austenite by forming very stable carbides. The resulting properties can simulate those of a lower-C grade. Ti can also somewhat neutralize the effect of excessive P. Some European practice is apparently based on this idea. Micro-alloying additions (less than 1 %) of Ti, V, B(boron), Zr (zirconium), and N2 (nitrogen) have been reported to promote grain refinement in Mn steels. The effect, however, is inconsistent. Higher levels of these elements can result in serious losses in ductility. N2 in quantities higher than 0.2 % can cause gas porosity in castings. An overall reduction in grain size lowers the susceptibility of the steel to hot tearing.

The sulphur (S) content in Mn steels rarely influences its properties, since the scavenging effect of Mn operates to eliminate S by fixing it in the form of innocuous, rounded, sulphide inclusions. The elongation of these inclusions in wrought steels can contribute to directional properties. In cast steels, such inclusions are harmless. However, it is better to keep S as low as is practically possible to minimize the number of inclusions in the micro-structure which are the potential sites for the nucleation of fatigue cracks in service.

Higher Mn content steels – Austenitic steels with a higher Mn content (higher than 15 %) have recently been developed for applications needing low magnetic permeability, low temperature (cryogenic) strength, and low-temperature toughness. These applications stem from the development of super-conducting technologies used in transportation systems and nuclear fusion field and to meet the need for structural materials to store and transport liquefied gases. For low magnetic permeability, these alloy steels have a lower C content than the regular Hadfield steels. The corresponding loss in yield strength is compensated for by alloying with V, N2, Cr, Mo, and Ti. Cr also imparts corrosion resistance, as needed in some cryogenic applications.

Austenitic steels with a higher Mn content are used in the heat-treated (solution-annealed and quenched) condition except for those which are age-hardenable. Wrought alloy steels are available in the hot-rolled condition. The micro-structure is normally a mixture of gamma (face centered cubic, or fcc) austenite and epsilon (hexagonal close-packed, or hcp) martensite. These alloy steels are characterized by good ductility and toughness, both especially desirable attributes in cryogenic applications. Further, the ductile-brittle transition is gradual, not abrupt.

Since the stability of the austenite is composition dependent, a deformation-induced transformation can occur in service under certain conditions. This is normally undesirable since it is accompanied by a corresponding increase in magnetic permeability. Additions of S, Al, and Ca (calcium) are made to improve the machinability of these alloy steels where needed. Because of their lower C content, majority of these alloy steels are readily weldable by the shielded metal arc welding (SMAW), gas metal arc welding (GMAW), and electron beam welding (EBW) processes. The composition of the weld metal is similar to that of the base metal and tailored for low magnetic permeability. The P content is normally maintained below 0.02 % to minimize the tendency for hot cracking.

Physical properties – The coefficient of thermal expansion decreases with increasing Mn content and normally ranges from 0.000008/deg C to 0.000017/ deg C. The electrical resistivity of a 0.7C-15Mn-1Ni alloy steel is reported to be around 710 micro-ohm. mm. The thermal conductivity of the 32 % Mn-7 % Cr alloy has been determined around 3 watts per meter-kelvin (W/m · K) at room temperature.

Another class of austenitic steels with high Mn additions has been developed for cryogenic and for marine applications with good resistance to cavitation corrosion. Potential applications also include those which need high temperature corrosion resistance. These alloy steels have been viewed as economical substitutes for conventional austenitic stainless steels since they contain aluminum (Al) and Mn instead of Cr, and Ni.

These alloy steels are normally of higher strength but lower ductility than conventional stainless steels e.g., grade 304 steel. The microstructure in these alloy steels is a mixture of gamma austenite and epsilon martensite, and in some cases (especially when the Al content exceeds around 5 %) alpha (bcc) ferrite is also obtained by a solution-annealing heat treatment. There is a tendency for an embrittling beta-Mn phase to form in the high-Mn compositions during aging at higher temperatures. The result is a substantial decrease in ductility. The addition of Al to some extent suppresses the precipitation of this compound.

While these alloy steels need special precautions during melting because of their high Al contents, they are fairly easily hot worked and are usable in wrought form. Data on weldability is not readily available. However, it is supposed that the overall weldability is poor.

The oxidation resistance of these alloy steels is inferior to that of the Cr-containing grades of stainless steels because of the poor adhesion of aluminum oxide (Al2O3) scale. This is especially visible under thermal cycling conditions.

Melting practice – AMS can be produced in any of the conventional steel making process. However, the present practice includes electric arc furnace (EAF) and electric induction furnace (IF). Some manufacturers choose acidic furnace lining as a lower cost option, but acidic lining is attacked by Mn oxides, resulting in Si pick up in melts during successive heats. Majority of the manufacturers use basic lined furnaces and prefer neutral lining for modified high- end AMS products range. Advantages of EAF melt route are (i) use of Si burn to stir the metal, (ii) achieve close control of elements in alloy chemistry, and (iii) removal of oxides and lowering of gas levels (de-oxidized melts, improves fluidity, bi-film formation).

Typical charge materials include C and Mn steel scrap, high-C ferro-manganese, ferro-silicon, and silico-manganese. Alloying elements such as Cr, Mo, and V are normally added as ferroalloys, while elements such as Ni are used in a nearly pure metallic state. Deoxidation of the steel is done with Al prior to pouring.

It is highly beneficial to keep levels of non-metallic inclusions and that of H2 (hydrogen), O2 (oxygen) and N2 (nitrogen) below 150 ppm (parts per million), 300 ppm, and 300 ppm respectively. Lower is the S, and P levels, better is the product. In particular, N2 levels above 300 ppm is known to result in severe gas porosity and pin hole gas defects in AMS castings. Customers and end users are sacrosanct about chemical test certificate compliance to alloy chemistry or specification given in the standard, but are frequently unaware to H2, N2, and O2 ppm levels in melt and actual solidified AMS casting which affects considerably the properties and its useful life.

In the case of majority of the castings, the pouring temperature is regulated to less than 1,470 deg C to prevent an excessively coarse grain size and minimize chemical segregation and other related casting defects. Sand castings are produced with olivine rather than silica sand to prevent mould-metal reaction. The grain size of wrought Mn steels is normally much lower than that of castings because of the recovery and recrystallization of the austenite grains during the hot rolling process.

The fluidity of AMS is quite good to enable pouring it into complex shapes and geometries at low degrees of super-heats with ability to place complex cores to achieve desired hollow shapes. Studies have shown that mechanical properties of AMS are highly improved by finer grain size. Strength and ductility can be as much as 30 % higher for fine-grained AMS. Fig 4 represents solidification cross-section of a 50 mm x 50 mm broken bars to show as-cast grain size.

Fig 4 Solidification cross-section of a 50 mm x 50 mm broken bars to show as-cast grain size

One consideration when selecting AMS grade is the section thickness of the desired part. As section thickness increases, it becomes harder to get good casting properties in AMS or to curb casting defects and metallurgical defects. AMS possesses low thermal conductivity when compared to other steels and to get good toughness, it needs to rapid quench post solution annealing. Also, higher the cross-section of casting, more is the probability of precipitation of carbides which cannot successfully phase transform in any heat treatment cycle. Hence, it is desirable to keep the cross-section less than 150 mm or alternatively adopt rapid solidification (chills) and directionally solidification techniques to get good casting properties.

AMS grades are normally produced in EAFs using basic melting practice. Although AMS grades in the as-cast condition are normally considered too brittle for normal use, there are exceptions to this rule. Mechanical properties of as-cast AMS grades indicate that lowering C content to less than 1.1 % and / or adding around 1 % Mo or around 3.5 % Ni results in commercially acceptable as-cast ductilities in light and moderate section thicknesses. This also applies to weld deposits, which are normally left in the as-deposited condition and hence are basically equivalent to material in the as-cast condition.

Adjustments in composition which limit carbide embrittlement and austenite transformation reduce or remove cracking of Mn steel castings during cooling in the moulds or reheating for solution treatment. The steel is normally poured at temperatures just high enough to avoid misruns in the castings and excessive skulling of the metal in the ladle. This practice helps in ensuring rapid solidification of the metal in the moulds, which in turn helps in preventing an excessively coarse grain size. The final micro-structure in majority of the castings is not fully austenitic, but contains carbide precipitates and pearlite in an austenitic matrix.

Commercial use of castings in the as-cast condition results in cost and energy savings and eliminates the problems of decarburization of thin castings during solution treatment and warpage during water quenching. High ductility or toughness are not needed in certain applications of Mn steels, e.g., as-cast Mn steels have been used successfully for pans on pan conveyors and for other light-section applications where solution-treated and water-quenched castings are prone to severe warpage.

The 6Mn-1Mo grade is the most susceptible to strain-induced martensite embrittlement because of its low Mn levels. However, even this grade can be used in applications for which 1 % elongation (either determined in a tensile test or estimated from a transverse bend test) is considered adequate ductility, e.g., the as-cast 6Mn-1Mo grade containing 0.8 % C to 1 % C has been used successfully in grinding mill liners.

As section thickness increases, the rate at which castings cool in sand moulds decreases. This increases the opportunity for embrittlement by carbide precipitation. Shapes which tend to develop high residual stresses, such as cylinders and cones, can be particularly affected. These stresses most probably result from volume changes accompanying the carbide precipitation and austenite transformation which occur during the normal cooling of castings. Fig 5a shows the volume changes which occur during the isothermal decomposition of a 1.25C-12.8Mn steel at temperatures between 850 deg C and 500 deg C, the principal range within which embrittlement occurs when a casting is cooled in its mould or reheated for re-austenization.

Fig 5 Change in length during isothermal transformation and structures of austenitic manganese steel

Between 850 deg C and around 705 deg C, only carbides are precipitated, mainly as envelopes around austenite grains and as lamellar-type patches within grains. The lamellar carbide patches have the appearance of coarse pearlite, but actually they are carbide plates in austenite. Below 705 deg C, and particularly between 650 deg C and 550 deg C, pearlite nodules, nucleated by previously precipitated carbides, grow at a relatively rapid rate.

Transgranular acicular carbides also tend to precipitate below around 600 deg C, especially in austenite containing more than around 1.1 % C. This precipitation can continue down to around 300 deg C in a 1.2C-12Mn steel. It can be followed by the transformation of some of the C-depleted austenite to martensite as the ambient temperature is approached.

Heat treatment – Heat treatment strengthens AMS so that it can be used safely and reliably in a wide variety of engineering applications. Solution annealing and quenching, the standard treatment which produces normal tensile properties and the desired toughness, involves austenitizing followed quickly by water quenching. Fig 5b shows the microstructures of a 76 mm section of AMS in the as-cast condition and after solution annealing and quenching. Variations of this treatment can be used to improve specific desired properties such as yield strength and abrasion resistance.

Normally, a fully austenitic structure, basically free of carbides and reasonably homogeneous with respect to C and Mn, is desired in the as-quenched condition, although this is not always attainable in heavy sections or in steels containing carbide-forming elements such as Cr, Mo, V, and Ti. If carbides exist in the as-quenched structure, it is desirable for them to be present as relatively innocuous particles or nodules within the austenite grains rather than as continuous envelopes at the grain boundaries.

Full solution of carbides needs a solution-treating temperature which exceeds Acm temperature by around 30 deg C to 50 deg C. Soaking for 1 hour to 2 hours at temperature is normally sufficient. Although it can appear that temperatures above 1,100 deg C permit use of 1.4 % C to 1.5 % C, three factors discourage the use of very high temperatures. These are (i) incipient melting occurs in areas of C and P segregation, (ii) scaling and decarburization become excessive, and (iii) commercial quenching rates are limited in their ability to retain high C concentrations in solution.

The commercial heat treatment of Mn steel castings normally involves heating slowly to 1,010 deg C to 1,090 deg C, soaking for 1 hour to 2 hours per 25 millimetres (mm) of thickness at temperature, and then quenching in agitated water. There is some tendency for the austenite grains to grow during soaking, especially in wrought Mn steels, although final austenite grain size in castings is largely determined by pouring temperature and solidification rate.

For AMS grade having C -1.05 % to 1.45 % and Mo – 1.8 % to 2.1 %, a modified heat treatment is frequently specified or desired. This treatment consists of heating castings to around 595 deg C and soaking them for 8 hours to 12 hours at temperature, which causes substantial quantities of pearlite to form in the structure. The castings are then further heated to around 980 deg C to re-austenitize the structure. This step converts the pearlitic areas to fine-grain austenite containing a dispersion of small carbide particles, which remain undissolved as long as the austenitizing temperature does not exceed around 1,010 deg C. Quenching then results in a dispersion-hardened austenite, which is characterized by higher yield strength, higher hardness, and lower ductility than is achieved if the same steel has been given a full solution treatment at a higher austenitizing temperature. This dispersion-hardening heat treatment permits a relatively high C content which in turn can improve abrasion resistance.

Speed of quenching is important, but it is difficult to increase beyond the rate of heat transfer from a hot surface to agitated water or beyond the rate fixed by the thermal conductivity of the metal. As a result, heavy-section castings have lower mechanical properties at the centre than the thinner castings. Fig 6a shows the cooling rates which can be expected when AMS plates of four different thicknesses are quenched in water.

Fig 6 Cooling curves for austenitic manganese steel and true stress versus engineering strain

Residual stresses from quenching, coupled with the lower properties of heavy sections, establish the normal maximum thickness of commercial castings at around 125 mm to 150 mm, although castings with sections up to 400 mm thick have been produced. There have been reports of two-step quenching cycles being used to decrease the overall thermal gradient and hence to reduce the residual stresses associated with heat-treatment operations. However, this is not a normal practice in the industry and is probably restricted to specific products. Residual stresses in Mn steels are not a critical issue because of their inherent toughness.

The relatively high austenitizing temperature leads to marked surface decarburization by furnace gases and to some loss of Mn. Surface decarburization can extend as much as 3.2 mm below the casting surface. Hence, the skin can be partly martensitic at times and normally shows properties less desirable than those of the underlying metal. This characteristic is not important in components subjected to abrasion, such as those used in crushing or grinding, since in these applications the skin is removed by normal wear.

Tensile deformation in service sometimes produces several cracks in this inferior skin, which terminate where they reach the tough austenite of normal composition except along grain boundaries, which contain mostly continuous carbides (e.g., because of slack quenching during heat treatment). Service performance is not seriously affected except under critical fatigue conditions or in very light sections. In such cases, premature failure can result. If considered necessary, a proprietary grade (equivalent to normal ASM grade to which 6 % Cr has been added) which is less prone to surface decarburization can be specified. Under certain conditions, sections such as wrought sheets can be protected with inert or reducing gas atmospheres, or with covers, metal envelopes, or organic or inorganic coatings, to minimize decarburization.

For applications which need non-magnetic properties, alteration of the skin by furnace gases needs attention. If the affected skin, which is normally magnetic, is quite shallow, it can be possible to remove it by pickling. In heavy sections, where skin thickness can approach 3.2 mm, the altered layer is to be removed by grinding if a surface permeability of less than 1.3 is needed.

Mechanical properties after heat treatment – As the section size of Mn steel increases, tensile strength and ductility decrease substantially in samples cut from heat-treated castings. This occurs since, except under specially controlled conditions, heavy sections do not solidify in the mould fast enough to prevent coarse grain size, a condition which is not altered by heat treatment. Fine-grain samples can show tensile strength and elongation as much as 30 % higher than those of course-grain samples. It is important to note that care is to be taken during tensile tests for ensuring that the gauge section in cast samples contains an adequate number of grains as needed by the standards.

Grain size is also the chief reason for the difference between cast and wrought Mn steels (the latter are normally of fine grain size). For cast grade, the standard deviations for tensile strength and elongation are around 69 MPa and 9 % respectively. The midrange values of 825 MPa and 40 % apply to sound, medium-grain cast samples which have been properly heat treated. The scatter bands for this grade extend from 620 MPa to 1,035 MPa for tensile strength and from 13 % to 67 % for elongation.

Mechanical properties vary with section size. Tensile strength, tensile elongation, reduction in area, and impact strength are substantially lower in 100 mm thick sections than in 25 mm thick sections. Since section thicknesses of production castings are frequently from 100 mm to 150 mm, this factor is an important consideration for proper grade specification. Notched-bar impact test values can be exceptionally high. Charpy test samples are sometimes bent and dragged through the machine rather than being fractured. Occasionally, observed values are biased because of the incorrect preparation of samples. Notches are to be cut by precision grinding to minimize work hardening at the apex of the notch.

Austenitic manganese steel remains tough at sub-zero temperatures above the Ms temperature. The steel is apparently immune to H2 embrittlement, although embrittlement has been produced in steels with low C content (less than 0.02 %) and high Mn content. There is a gradual decrease in impact strength with decreasing temperature. The transition temperature is not well defined since there is no sharp inflection in the impact strength-temperature curve down to temperatures as low as -85 deg C.

At a given temperature and section size, Ni and Mn addition are normally beneficial for improving the impact strength, while higher C and Cr levels are not. Resistance to crack propagation is high and is associated with very sluggish progressive failures. Because of this, any fatigue cracks which develop can be detected, and the affected part or parts removed from service before complete failure occurs, a capability which is a distinct advantage in railway track work.

The fatigue limit of AMS has been reported as 270 MPa. Yield strength and hardness vary only slightly with section size. The hardness of the majority of the grades is around 200 HB (Brinell hardness) after solution annealing and quenching, but this value has little significance for estimating machinability or wear resistance. Hardness increases so rapidly because of work hardening during machining or while in service that AMSs are to be evaluated on some basis other than hardness.

The true tensile characteristics of Mn steel are better revealed by the stress-strain curves in Fig 6b which compare Mn steel with gray iron and with a cast alloy steel quenched and tempered of around the same nominal tensile strength. The low yield strength is important and can prevent the selection of this alloy where slight or moderate deformation is undesirable, unless the usefulness of the components in question can be restored by grinding. However, if the deformation is immaterial, the low yield values can be considered temporary, i.e., deformation produces a new, higher yield strength corresponding to the quantity of strain which is absorbed locally.

Bend tests are frequently used as a qualitative indication of ductility in castings. Description of the bend test sampling and testing procedure is given in the standards. Typically, a separate test sample is poured from the same heat and is heat treated with the batch of castings. This sample is then tested by cold bending around a 25 mm diameter mandrel without any further machining or grinding, except which is needed for removing surface decarburization. The ductility is judged to be acceptable in the majority of the cases if the sample can be bent through an angle of 150-degree without breaking into two pieces.

Work hardening – The approximate ranges of tensile properties produced in constructional alloy steels by heat treatment are developed in AMSs by deformation-induced work hardening. In a tension test, yielding indicates the beginning of work hardening, and elongation is associated with its progress. Little or no necking occurs since work hardening is highest at the point of highest deformation. The increase in strength because of cold work stops further elongation, and deformation then occurs elsewhere in the gauge section. This type of behaviour is frequently compared to that shown by transformation-induced plasticity (TRIP) steels. Elongation hence occurs uniformly and without flow stress saturation, until failure.

A 1988 study has shown that fracture occurs because of a combination of micro-void coalescence (without shear localization) and surface cracking within regions of localized plastic flow. The nucleation of both voids and surface cracks has been observed to be a function of carbides and non-metallic inclusions in the steel, to some extent. Plastic deformation is accompanied by the development of texture. X-ray diffraction studies have shown this texture to be in tension and in compression, similar to the behaviour of 70-30 brass. Also, work-hardening rates tend to be higher in compression than in tension.

The mechanism of work hardening has been the subject of several studies. Different mechanisms have been found to contribute to work hardening, depending on factors such as alloy composition (stacking fault energy, strain rate sensitivity), temperature, and strain rate. These mechanisms include twinning or pseudo-twinning, stacking fault formation, and dynamic strain aging. However, it is well established that a deformation-induced transformation from austenite to alpha martensite (bcc) does not occur in ordinary AMSs. The role of such a transformation in work hardening is more important only at lower C and Mn levels. On a macroscopic scale, the work-hardening rate has been observed to increase with increasing C content and decreasing grain size. Work-hardening rates of 1,500 MPa to 2,500 MPa have been measured in AMSs at nominal strain rates of 0.0001/S at room temperature.

Determination of work-hardening rate – The work-hardening rate is normally determined from ordinary tensile or compression tests as the slope of the true stress-true strain curve, which is mostly linear in the plastic region.

One other measure of the tendency of AMSs to work harden is based on a determination of the so-called Meyer index or exponent. The technique uses a 10 mm diameter Brinell ball indentor and a series of loads. The test loads are plotted against the diameters of the corresponding indentations on logarithmic scales. The result is expected to be a straight line which fit the equation ‘P = A x d to the power n’, where ‘P’ is the applied load, ‘d’ is the diameter of the indentation, ‘A’ is a constant, and ‘n’ is a measure of the tendency of the metal to strain harden (also called the Meyer index or exponent). The Meyer index of a variety of AMSs and stainless steels has been determined to be in the range of 2.17 to 2.6.

Manganese steels are unequaled in their ability to work harden which exceeds even the metastable austenitic stainless steels in this feature, e.g., a standard grade of Mn steel containing 1 % C to 1.4 % C and 10 % Mn to 14 % Mn can work harden from an initial level of 220 HV (Vickers hardness) to a maximum of more than 900 HV. After extended service, the hardness at the wearing surfaces of railway frogs typically ranges from 495 HB to 535 HB. Maximum attainable hardness depends on several factors, including specified composition, service limitations, method of work hardening, and pre-service hardening procedures. It appears that rubbing under heavy pressure can produce higher values of maximum attainable hardness than can be produced by simple impact.

Service limitations – In some cases, abrasion can remove surface metal before it can attain maximum hardness. In other cases, work hardening raises the elastic limit to the point which succeeding impact blows cannot use more plastic flow and hence merely bounce off. However, if a rotating tool is applied with enough force and does not cut, accentuated work hardening can be expected. This is a common occurrence during drilling. A sharp twist drill of superior steel can drill 13 % Mn steel provided a deep enough cut is taken, but cutting ceases if the drill becomes dull. When this happens, it is frequently futile to continue drilling, even with a sharp drill, since the bottom of the hole has become so hard that the cutting edge cannot penetrate it.

The low yield strength of Mn steel is sometimes a disadvantage in service, e.g., plastic deformation because of wheel impact, as in railway frogs and crossings, increases yield strength to levels more resistant to flow, but the associated changes in dimensions are undesirable. With time, low spots develop at critically pounded locations, eventually needing rebuilding with weld deposits. This issue is lessened to some degree by hardening the running surface prior to in-service installation. Low yield strength can also be a disadvantage when Mn steel is used in light armour and similar applications.

Since much energy is absorbed during work hardening, Mn steel sheet is an effective light armour against slow-moving projectiles. However, Mn steel is relatively ineffective against high-velocity projectiles which shear through the armour with little accompanying deformation. One reason for this can be the negative strain rate sensitivity behaviour shown by ordinary AMSs. Heat-treated (quenched and tempered) alloy steels of higher yield strength are preferred for armour against high-velocity projectiles. However, majority of the impact blows encountered in industrial service are of low velocity, and for this service, Mn steel is acceptable. It is the preferred choice for applications needing high impact resistance, toughness, and absorption of energy.

Work-hardening methods – Work hardening is normally induced by impact, as from hammer blows. Light blows, even if they are of high velocity, cause shallow deformation with only superficial hardening even though the resulting surface hardness is ordinarily high. Heavy impact produces deeper hardening, normally with lower values of surface hardness. The course of flow under impact and the associated increase in hardness are shown in Fig 7, which compares a standard 12 % Mn steel with an air-hardening Cr-Ni- Mo alloy steel. Less well known is the fact that abrasion itself can produce work hardening.

Fig 7 Plastic flow and work hardening of a manganese steel and an air-hardening steel

Explosion hardening has been developed as a substitute for hammer or press hardening to achieve hardening with less deformation. Pentaerythritol tetra-nitrate (3.1 mg/square mm) in the form of plastic explosive sheet 2.11 mm thick or mixtures of ammonium nitrate and trinitrotoluene (TNT) are cemented to the surface of the steel and detonated. Normally, three explosions are needed to attain the desired hardness in railway track work. More explosions do not significantly change the pattern of hardening, but increase the possibility of cracking instead.

The use of plastic explosive permits the hardening of areas such as track work flange-ways and unsupported sections which cannot be hammered satisfactorily. Fig 8a compares the depth and intensity of hardening from three different explosion treatments and from one hammer peening operation applied to railway track work. Explosion hardening is now considered a satisfactory but expensive substitute for the hammer or press hardening of Mn steel trackwork castings.

Fig 8 Hardening patterns and flow under repeated impact

Explosion treatment has also been applied to components which are to be subjected to abrasive wear. Initial reports have been favourable, but have been reversed by subsequent experience. There is no solid evidence that explosion hardening is advantageous for service involving grinding or gouging abrasion. Explosion hardening is accompanied by insignificant deformation in spite of the fact that work hardening is normally associated with plastic flow and deformation. Studies have centred on the principle that a different hardening mechanism, other than twinning, is involved in explosion hardening.

The addition of alloying elements such as V, Cr, Si, and Mo is also an effective means of raising yield strength, but V, Si, and Cr reduce ductility. The relative effects of alloying and of pre-hardening by deformation are compared in Fig 8b.

Reheating – Before Mn steel components are reheated in the field, the effects of such reheating are to be seriously considered. Unlike ordinary structural steels, which become softer and more ductile when reheated, Mn steels suffer reduced ductility when reheated enough to induce carbide precipitation or some transformation of the austenite. Fig 9a shows the micro-structural effects of such heating. As a general rule, Mn steels are never to be heated above 260 deg C, either intentionally or accidentally, unless such heating can be followed by standard solution annealing and quenching.

Fig 9 Structure of reheated manganese steel and time-temperature relationship

Time, temperature, and composition are variables in the embrittlement process. At lower temperatures, embrittlement takes longer to develop. The time-temperature relationship in 13Mn-1.2C-0.5Si steel is shown in Fig 9b, which presents data based on metallographic examination for structural changes indicate the beginning of embrittlement. At 260 deg C, transformation needs more than 1,000 hours, reheating to as high a 425 deg C, even with close control of temperature, can be done for no longer than 1 hour if transformation is to be avoided. Fig 10 shows the effect of composition on the magnitude of embrittlement.

Fig 10 Embrittlement from reheating manganese steel

Since large castings are sometimes mounted with backings of molten lead or zinc, the temperature of such castings is to be carefully controlled during reheating. The time-temperature relationship is also to be given due consideration for components which are to be welded. When 12 % to 14 % Mn steels are to be heated above around 290 deg C during service or welding, it is desired that the C content is held below 1 %, which suppresses embrittlement for at least 48 hours at temperatures up to 370 deg C. The addition of 1 % Mo suppresses embrittlement completely at temperatures up to 480 deg C and partly suppresses it at temperatures of 480 deg C to 595 C. If the C content is held below around 0.9 %, the addition of 3.5 % Ni completely suppresses embrittlement up to 480 deg C and partly suppresses it above this temperature. These rules can be expected to apply during heating periods of up to 100 hours. For periods of 1,000 hours or more, limiting temperatures are substantially lower. It is to be noted that localized C contents can still exceed 1 % because of chemical segregation. Hence, these guidelines are to be used with caution.

Three 1 % C Mn-steels (12 % Mn steel, 12Mn-Mo steel, and 12Mn-3.5 Ni steel) were reheated for periods of up to 10 hours at temperatures up to 595 deg C. At 370 deg C, none of the three steels was embrittled. In fact, there were indications that the treatment had caused slight improvements in ductility. At 480 deg C, the plain Mn steel and the Mn-Ni steel were embrittled, but there was little change in the properties of the Mn-Mo steel. At 595 deg C, all three steels were embrittled, but the plain Mn steel was more severely affected than the others.

Embrittlement, as revealed by micro-structural investigation, is caused by the formation of acicular carbide and pearlite in the austenite grains of the plain Mn, by carbide nodules surrounded by pearlite within the austenite grains of the Mn-Mo steel, and by an envelope of pro-eutectoid cementite around each grain of the Mn-Ni steel. If the C content in the Mn-Ni steel is lower, the steel is likely to be substantially less susceptible to embrittlement. In another study, a 0.9C-14.3Mn-1.75Si-3.4Ni steel did not become significantly embrittled when it was heated for 1.5 hours at 480 deg C, and its ductility was reduced by only 17 % to 20 % when it was heated for 1.5 hours at 595 deg C to 760 deg C.

Wear resistance – Compared to the majority of other abrasion-resistant ferrous alloys, Mn steels are superior in toughness and moderate in cost, and it is mainly for these reasons that they are selected for a wide variety of abrasive applications. They are normally less resistant to abrasion than are martensitic white irons or martensitic high-C steels, but are frequently more resistant than pearlitic white irons or pearlitic steels.

The type of wear which is sustained has a major influence on the performance of Mn steels. They have (i) very good resistance to metal-to-metal wear, as in sheave wheels, crane wheels, and car wheels, (ii) good resistance to gouging abrasion, as in equipment for handling or crushing rock, (iii) intermediate resistance to high-stress (grinding) abrasion, as in ball mill and rod mill liners, and (iv) relatively low resistance to low-stress abrasion, as in equipment for handling loose sand or sand slurries.

Metal-to-metal contact – In applications involving metal-to-metal contact, the work hardening of Mn steel is a distinct advantage sine it decreases the coefficient of friction and confers resistance to galling if temperatures are not excessive. Compressive loads, rather than impact loads, provide the deformation needed, producing a smooth, hard surface which has good resistance to wear but which does not abrade the contacting component. Sheaves, wear plates, and castings for railway track work are common applications of this type. Mn steel also has been used in some water-lubricated bearings.

Railway centre plates, which are the bearing surfaces where trucks swivel under freight cars, provide a good example of the merit of Mn steel for metal-to-metal frictional wear. Initially lubricated, these plates soon are operating dry and are accessible to air-borne grit. When plates are made of C steel, the mating surfaces become rough from galling (adhesive wear), the increased friction prevents easy motion, the trucks do not swivel properly on curves, and the lateral pressure is so accentuated that early wheel flange wear is induced. Poor swiveling can accentuate thrust loads at the ends of wheel bearings, generating excessive frictional heat, and centre plates and mating bowls on trucks can wear severely.

Several years ago, the substitution of 13 % Mn steel centre plates on heavily loaded ore cars demonstrated their superiority. More recently, several service tests have established that Mn steel not only wears less than C steel, but also develops a low-friction polished surface. The advantages of Mn steel centre plates are most evident for cars which are very heavily loaded. Mn steel wear plates in a blooming mill housing were in good condition after 16 years, whereas C steel plates needed replacement every 2 years. Mn steel wear plates on ore bridge haulage drums were reported to be as good as new after 5.5 years of service, whereas ordinary cast steel plates wore out in around 3 years. Mn steel sheave wheels are expected to last around four times longer than cast C steel wheels and they do not groove and cause excessive rope wear, as do wheels made of alloys which do not work harden. In steel plant applications, welded Mn steel overlays on large mill-coupling boxes, pinions, spindles, and other items working under heavy impact loads perform satisfactorily.

Abrasion – The concept that Mn steel has poor wear resistance unless it has been work-hardened is not a valid generalization. The mis-understanding has probably developed since, where significant impact and attendant work hardening are present, 12 % Mn steel is so clearly superior to other metals that its performance is attributed to surface hardening. However, controlled abrasion tests have indicated that there are circumstances under which the abrasion resistance of AMS is modified little by preservice work hardening, and other circumstances under which this steel outwears harder pearlitic white cast irons without work hardening.

In applications which involve heavy blows or high compressive and structural stresses, the very hard and abrasion-resistant martensitic cast irons can wear more slowly than Mn steel. However, these irons normally fail by early fracture with a considerable portion of the original cross section unworn, whereas Mn steel can become almost paper thin before fracturing. Pearlitic white cast iron, which has a hardness of around 400 HB to 450 HB, is equally brittle but less resistant to wear. Comparative tests on log washer lugs indicated that Mn steel was around 25 % worn out with no breakage, whereas in the same period white iron lugs wore to the point of uselessness with 14 % breakage.

In clay crusher rolls, Mn steel lasted two to three times as long as white or chilled iron. In grinding barrel liners, cast irons lasted 2 to 3 years compared to 10 years for Mn steel. Part of the superiority of Mn steel over white cast iron is attributed to higher freedom from breakage and spalling, but some is probably because of better intrinsic wear resistance. Mn steel chain, with endless links cast in inter-locking moulds, also provides resistance to wear, lasting three to nine times as long as heat-treated steel chain in certain applications. Mn steel is valuable in conveyors as well as in dragline chain subjected to abrasion and used for carrying heavy loads. Mn steel is not satisfactorily resistant to wear by a stream of air-borne abrasive particles (impingement erosion), such as in sand-blasting or grit-blasting equipment, and hence is not to be selected for such service.

Abrasion testing – A number of laboratory abrasion and wear testing procedures have been developed to simulate a variety of applications with differing degrees of severity. These include a jaw crusher test for gouging abrasion resistance, a dry-sand / rubber wheel low-stress abrasion test, testing for slurry abrasiveness, and tests for galling resistance. Other procedures such as a wet-sand / rubber wheel low-stress abrasion test and a pin-on-disk high-stress abrasion test are also there.

Several variations of the pin test are used to measure wear resistance. One test method utilizes abrasive paper or cloth mounted on a revolving disk, the abraded flat tip of a small pin moving in a spiral path. One laboratory used a milling machine platen with a revolving pin taking a zig-zag path. The Bureau of Mines and some workers in Australia have the pin making a spiral path against garnet paper or cloth fastened to a large revolving drum. The common feature of these test methods is a small pin sample, but the pins are subjected to a variety of conditions, e.g., normal load, tangential force, abrasive type, abrasive grain size, abrasive angularity or roundness, temperature, abrasive concentration, and so on. A considerable quantity of test data on abrasion resistance of a wide variety of AMSs has been accumulated.

There is little need for laboratory tests in the case of well authenticated applications, but such tests can be useful to judging suitability for new applications. The usefulness of these tests is dependent on their precision and reliability (versus statistical standards) and their validity as determined from comparisons with service tests. Crusher tests are slow and expensive, motivating attempts to correlate more convenient procedures, such as a pin on abrasive paper or cloth.

Another jaw crusher test on 28 alloys shows a more even hardness spread. The correlation coefficients highlight the fallacy of judging Mn steel by its initial (as-heat-treated) hardness, even though there is a direct relationship between this property and abrasion resistance for several C and low-alloy steels. One reason for the degradation in correlation coefficient because of the inclusion of Mn steel is that the actual hardness developed at the wear surface because of work hardening is considerably higher. Though 500 HB to 600 HB is normally considered the limit in service, laboratory studies have shown that the hardness can be as high as 900 HV.

The mediocre correlation coefficients and large standard errors emphasize that laboratory tests are frequently handicapped by large scatter. Frequently a certain quantity of judgment is necessary to be able to apply the results of laboratory tests to new applications. The abrasion resistance of austenitic 12 % to 14 % Mn steels with different C contents has been compared with the resistance of other steels and white irons in a jaw crusher abrasion test (Fig 11a).

Fig 11 Relative wear ratio and temperature dependence of flow stresses

In Fig 11a, the wear rate for a quenched and tempered low-C low-alloy steel (at 269 HB), which was used as a comparative standard in each test, is also shown. When the relative wear rate (wear ratio) of each test material is plotted against increasing C content on a log-log scale, results for austenitic steels and irons tend to fall on a descending straight line, and results for martensitic steels and irons fall on a parallel line below the line for the austenitic alloys. A decrease in wear ratio represents a proportionate increase in abrasion resistance.

Hence, Fig 11a strongly supports the conclusions that the abrasion resistance of both austenitic and martensitic steels improves with increasing C content and that, for a given C content, martensitic steels have better abrasion resistance than do austenitic steels. However, martensitic steels and white irons have limited resistance to gouging abrasion because of their lack of toughness. The wear ratios of pearlitic steels, if plotted on Fig 11a would lie above those of the austenitic steels. There is considerable scatter in the rates for pearlitic steels because of their wide variation in hardness for any given C content.

The same trends have been observed in ore milling tests and in sand slurry abrasion tests. Normally, austenitic 12 % to 14 % Mn steels have higher wear rates, and hence poorer abrasion resistance than pearlitic or martensitic steels of equivalent C content in a low-stress slurry abrasion. A 6Mn-1Mo composition is normally intermediate between pearlitic and martensitic steels in wear resistance.

It is to be noted that the high stress term has been carelessly used here and does not apply clearly to a definite condition. Originally used to characterize the wet-sand grinding-abrasion test, it is not appropriate when applied to a pin being worn by an abrasive anchored on paper or cloth which can have a cushioning effect.

Corrosion – Mn steel is not corrosion resistant. It rusts readily. Also, where corrosion and abrasion are combined, as they frequently are in mining and manufacturing environments, the steel can deteriorate or be dissolved at a rate only slightly lower than that of C steel. If the toughness or non-magnetic nature of Mn steel is necessary for a marine application, protection by galvanizing is normally satisfactory.

Effects of temperature – The outstanding properties of 13 % Mn steel between -45 deg C and 205 deg C make it useful for all ambient-temperature applications, even in arctic climates. It is not recommended for hot-wear applications because of structural instability between 260 deg C and 870 deg C. At higher temperatures, it can lack the strength and ductility necessary to withstand severe welding stresses, and hence welding is to be done under closely controlled conditions.

The 13 % Mn steel is not resistant to oxidation, and its creep rupture properties are inferior to those of Fe-Cr-Ni austenites. Since the work-hardening rate decreases at higher temperatures, this steel is not suitable for structural application in the red-heat range. The effects of temperature on mechanical properties (in both tension and compression) have been well documented). The normal trend for changes in flow stress versus temperature is shown in Fig 11b for both tension and compression. There is an increase in yield strength with decreasing temperature with a corresponding drop in ductility and ultimate tensile strength. The changes in strength and ductility, however, are not uniform, and the steel retains a major portion of its room temperature ductility down to around -100 deg C.

There has been some controversy with respect to the operative mechanism for work hardening in Mn steels at different temperatures. It appears presently that deformation twinning is predominant at the lower temperatures (below around 0 deg C). However, at temperatures above 0 deg C, strain hardening has been variously attributed to twinning, dynamic strain aging, and stacking fault formation. Some hardening because of the formation of Cottrell clusters and carbide precipitation has been reported at temperatures above 300 deg C. At -75 deg C, cast Mn steels retain from 50 % to 85 % of their room-temperature impact resistance. These steels are considerably more brittle at liquid-air temperature (-185 deg C), but at all atmospheric temperatures encountered by railway track work and mining and construction equipment, these steels have outstanding toughness which provides a valuable safety factor, compared to ferritic steels at sub-zero temperatures.

Associated with the embrittlement produced by reheating above 260 deg C, are changes in physical properties stemming from the same transformations which cause the loss in toughness. Since both composition and time at temperature influence these changes, erratic behaviour and a considerable range in such properties as thermal and electrical conductivity can be expected above 315 deg C.

Thermal-expansion characteristics of AMSs are similar to those of other austenitic materials. The expected change in length upon heating is around 1.5 times that of ferritic steels. A coefficient of linear thermal expansion of 0.000018/ deg C is normally precise enough near room temperature (Fig 12). Transformation to pearlite and the precipitation of carbide influence the values of the expansion of coefficient in the range from 370 deg C to 760 deg C.

Fig 12 Thermal expansion of a 13 % Mn steel

Magnetic properties – The untransformed austenite of 13 % Mn steel is almost non-magnetic, with a permeability of around 1.03 or less. This permits the use of the material where a strong, tough, non-magnetic steel is needed, as in magnet cover plates, collector shoes for traveling cranes, stator core parts for generators and motors, liner plates for storage bins holding materials which are handled by lifting magnets, magnetic-separator parts, instrument-testing devices, and furnace parts located in the magnetic fields of induction furnaces.

The changes which occur in the composition of the surface during heat treatment can produce a magnetic skin, one that is either a martensitic surface layer or a magnetic oxide. Permeability values of around 1.3 have been achieved on samples which have this magnetic surface layer. Frequently, this layer does no harm, but, if necessary, it can be removed by grinding or pickling or it can be prevented altogether by suitable (although frequently expensive) corrective measures during heat treatment. Cast or wrought Mn steel is probably the most economical material for strong non-magnetic parts if machining is not needed. Fabrication costs and anticipated operating temperature are decisive factors in selection with operating temperature does not exceed 260 deg C.

A 20 % Mn steel containing Bi has been developed for non-magnetic parts which need machining. Laboratory tests indicate a machinability comparable to that of the more expensive type 304 stainless steel that this material is designed to replace. It can be lathe turned (horizontal force – 135 N to 380 N at 1.35 metre/second), drilled, and tapped. This steel has not yet been exploited commercially.

Lack of magnetism is a disadvantage when AMS is used in components of material-handling systems which depend on magnetic separators to remove tramp iron from the process stream before it enters crushers, grinders, and other machinery. When there is the possibility that Mn steel components can become detached from working equipment and fall into the process stream, it is advisable to cast mild steel inserts in the Mn steel components. The inserts are to be large enough to provide the level of ferro-magnetism necessary for magnetic separators to detect and remove the lost components so they cannot enter working machinery and cause damage.

Welding – Several of the common applications of AMS steel involve welding, either for fabrication or for repair. Hence, it is important to understand that this material is remarkably sensitive to the effects of reheating, frequently becoming embrittled to the point of losing its characteristic toughness. Oxy-fuel gas welding is so likely to produce embrittlement that it is not accepted as a practical method of welding this alloy. When properly done, electric arc welding is the preferred method of joining or surfacing Mn steels. The P content is kept below 0.03 % to minimize hot cracking. Welding is normally carried out after heat treatment to minimize cracking after deposition of the weld metal.

Electrodes for arc welding AMS steel are commercially available in several compositions. They can be used for surfacing, for repair welds, and for joining Mn steel to itself or to C steels. They have a lowered C content to minimize C precipitation as they cool from the welding temperatures. Though formulated to avoid embrittlement of the deposited filler metal, proper welding procedures still are to be used to avoid damage in the heat-affected zone. It is normally desired that the metal temperature adjacent to the weld does not exceed 315 deg C after an elapsed time of 1 minute.

Electrodes of high Mn content, containing insignificant quantities of other alloying elements, are also available. Normally, these electrodes are desired only for the build-up of worn areas, since they are inherently lower in toughness than more highly alloyed grades. High-alloy low-Mn electrodes have not been normally accepted as equivalent to high-Mn types.

Austenitic stainless-steel electrodes such as type 308 or type 310 can be used for the repair of cracks or for the build-up of worn areas. Factors which are frequently overlooked are the losses in C, Mn, and Si which occur during welding. Although several electrode manufacturers compensate for these losses, improper welding techniques, such as use of excessive arc length and excessive puddling, can cause additional losses. The result is inferior properties in the weld deposit.

Frequently C steels are welded to high-Mn steel using austenitic stainless-steel electrodes. Since the deposit tends to be a mixture or hybrid of the base and the filler metal, it can have quite different properties. Frequently it is air hardening, producing a martensitic zone as the weld cools. The ductility of the martensite is low, but the strength is high, and weldments are frequently satisfactory. The chief adverse factor can be cracks in the martensite.

Cross-weld tensile properties of low-C 14Mn-1Mo steel plate welded to 1045 steel with EFeMn-A electrodes have been 435 MPa yield strength, 650 MPa ultimate tensile strength, and 11 % elongation with fracture in the 1045 steel. These properties are superior to those of several weldments of C steel. Filler metal from 0.75 % C, 15 % Mn, 3.5 % Ni, and 4.0 % Cr seems to be superior to that of EFeMn-A and EFeMn-B.

Grades having Cr at levels above 14 % are also useful for the joining of Mn steel and the build-up of worn parts such as trackwork, but because of their low C content, have sacrificed abrasion resistance. However, they are more machinable than the higher C grades. If used to restore the dimensions of crusher parts, they are to be overlaid with an effective hard-facing alloy.

The main consideration in welding AMS is minimum heating of the parent metal to avoid embrittling transformations or carbide precipitation. This precludes preheating. Under the most favourable circumstances, some precipitation is expected, and the resulting heat-affected zones seldom attain the toughness of normal parent metal. Since Mn steel work hardens in service, it can be assumed that any worn area needing repair or rebuilding has a work-hardened surface. This surface is to be removed before welding in order to prevent cracking in the heat-affected zone. Also, any fatigue cracks which can have nucleated in the work-hardened region can cause delamination of the weld if not removed.

The low heat conductivity and high thermal expansion of Mn steel also cause difficulties, combining to produce steep thermal gradients and high residual stresses. Weld beads are subjected to tension as they cool. For minimizing cracking, it is desirable to peen them while they are hot, producing plastic flow and changing the stress to compression. This hammering is to be done promptly after 150 mm to 230 mm of weld bead has been deposited.

Machining – Mn steels are so tough and work harden at the point of a cutting tool to such an extent that frequently they are considered commercially not machinable. However, these steels are regularly cut by adhering to normally accepted procedures. In addition, a new, more highly machinable grade of Mn steel has been developed which can be helpful in appropriate applications. Although details of practice and tool design differ, there is general agreement on the procedures for machining Mn steels.

Machine tools are to be rigid and in good condition. Any factors which encourage chatter are not desirable. Tools are to be sharp. Dull tools cause excessive work hardening of the cut surface and accentuate the difficulty in machining. Low speeds of around 9 metre per minute to 12 metre per minute are to be used. High speeds are likely to create red-hot chips and to cause rapid tool break-down. Cobalt (Co) high-speed steel tools or tools with cemented carbide and ceramic inserts can be used. The latter are preferred.

The liberal use of a good grade of S-bearing cutting oil is beneficial but not necessary. In castings, holes are to be formed by cores in the foundry, rather than by machining, whenever possible. Coolants are desired for surface grinding operations. Different sources provide statements in favour of both positive-rake and negative-rake tools and both dry cutting and liquid coolants. Since high temperatures at the cutting edge are a large part of the issue, effective cooling seems desirable. Negative-rake tools are likely to need more force and hence to produce more heat. However, the thinner edge of a positive-rake tool is more vulnerable to heat. Machinability is increased by the embrittlement which develops with reheating between around 540 deg C and 650 deg C. Although not normally practicable, such a treatment can be useful if the part can subsequently be properly toughened. Milling is normally not considered practicable.

A machinable grade 20Mn-0.6C steel has been developed specifically for improved machinability.  Even though the yield strength has been deliberately reduced from 360 MPa to a value between 240 MPa and 310 MPa to achieve improved machinability, the ultimate tensile strength exceeds 620 MPa, and elongation in small castings can reach 40 %. The heat treatment of this steel involves water quenching from 1,040 deg C. As-cast properties are lower but are probably adequate for several applications. This non-magnetic modified grade can be lathe turned, drilled, tapped, and threaded. Even holes 6 mm in diameter can be drilled and tapped in this metal. In some machine shops, it is rated only slightly more difficult to drill than plain mild steel, and the quality of the tapped threads is considered very good. Wear resistance has been sacrificed for machinability, and this grade has significantly less abrasion resistance than do the different types AMSs.