Patent Application
20150232969
The invention refers to hardening thermal treatment of the perlite-class steel.
A process of thermal treatment is known where selection or development of the steel chemical composition was frequently based on the necessity to ensure both through- and surface thermal hardening of the part by alloying so as to increase hardenability, the ability of the steel to harden in weak cooling media to prevent crack formation. Alloying was used a lot less frequently to improve steel quality features compared to carbon steels—increase strength, plastic and dynamic properties, heat resistance, lower cold brittleness threshold, etc.
A process of surface hardening with deep induction heating is known, that later on was called through-surface hardening, developed under the supervision of K. Z. Shepeliakovsky, Doctor of Technical Sciences, Professor, Honored inventor of the Russian Federation [1].
It was, therewith, shown that a high complex of mechanical properties can be achieved on 1st generation low (LH) and specified (SH) hardenability carbon- and low-alloy steels [2], [3].
These are steels whose hardenability complies with the effective loaded cross-section of the part; in this case, following TSH, the surface layers of this cross-section amounting to .1-.2 of the diameter (thickness) have a martensite structure with HRC ≈60 hardness, while the core hardness is HRC=30−45.
Steel hardenability is characterized by the ideal diameter (DI) value that actually defines the optimum hardened layer depth with reference to the specific part shaped as a cylinder, sphere or plate.
In principle, LH and SH steels have the same purpose and just conventionally differ only by the ideal diameter (DI) value; for LH steels as the earlier products, this diameter is equal to 8-16 mm, whereas for modern SH steels it is over 16 mm.
The necessary DI range of LH and SH steels for the specific type of parts was achieved by combined restriction in the upper limit of one or a group of admixture elements which led to a lower accuracy and a wider DI range.
The disadvantage of the known 2nd generation LH steels (see patent RU 2158320) is that the specified low hardenability value was achieved only by strict limitation of contents of all permanent admixtures—Mn, Si, Cr, Ni, Cu, which made smelting more difficult and resulted in lower DI range accuracy during the development of the chemical composition and, as a consequence, led to a wider deviation in the hardened layer depth that went outside the allowed tolerance range.
So, for example, LH steel with .8% C and containing Mn, Si, Cr, Ni, Cu (<1% each) and .06-.12% Ti (LH steels)—(see patent RU 2158320) has minimal hardenability—DI<12 mm for austenite size 10 grain and finer grain (#11) DI<11 mm and <10 mm for #12 grain size, whereas similar steel with .8% C; .05% Mn; .12% Si; .11% Cr; .25% Ni; .3%Cu; .05% Al; .22% Ti (with a wider range of permanent admixtures of Ni, Cu) has the same DI value.
The objective of this invention is the development of a process for thermal treatment during induction and furnace heating of the 3rd generation LH and SH steels.
The technical result is to obtain even finer austenite grain ##11-13 GOST5639 (ASTM), more stable preset hardenability level (DI), with a substantially smaller deviation and which strictly corresponds to the depth of the hardened layer obtained on parts subjected to thermal treatment using the proposed procedure, the ability of treating thinner, smaller and other parts with the through-surface and through-thickness hardening.
To achieve the technical result, a method for thermal treatment of parts made from low (LH) and specified (SH) hardenability structural steel is proposed for parts shaped as a sphere, cylinder, plate, including through-surface hardening by through-surface heating of the part or its effective cross- section, up to the temperatures of austenitization and cooling with a liquid refrigerant at the rate of more than 40,000 kcal/m2.h.0C. and tempering, with a distinction that subjected to hardening are parts made from steel containing the following components, weight %:
with the ideal diameter determined by the following mathematical expression (1);
D
kp
.=K·{square root over (C)}·(1+4.1·Mn)·(1+0.65·Si)·(1+2,33·Cr)·(1+0.52·Ni)·(1+0.27·Cu)·(1+3.14·Mo)·(1+1.05·W)·[1+1.5(0.9−C)]·(1−0.45C′)·(1−0.3Ti)·(1−0.35V)·(1−0.25 Al)
where Der is the ideal diameter (DI), mm,
K—is the coefficient whose value depends on the actual austenite grain size ##6-13 according to the ASTM scale, GOST5639, and is, respectively, equal to: 5.4 for #13 grain; 5.8 for #12 grain; 6.25 for #11 grain; 6.75 for #10 grain; 7.3 for #9 grain; 7.9 for #8 grain; 8.5 for #7 grain; 9.2 for #6 grain;
C, Mn, Si, Cr, Ni, Cu, Mo, W are the components, weight % contained in the austenite solid solution at the final heating temperature preceding hardening cooling;
[1+1.5(0.9−C)] is the multiplicand taken into account only if boron is present in steel in the amount of .002-.007%;
C′, Ti, V, Al are components, weight %, not contained in the austenite solid solution, but present in the form of structurally-free secondary carbonitride phases at the final heating temperature preceding hardening cooling, in which case C′ is the weight % of carbon content in excessive hypereutectoid steel cementite;
Structural steel with the specified chemical composition and with the same grain size has ideal diameter (DI) ranges:
To prevent hot-brittleness, the total content of manganese and titanium in the structural steel should be more than six times the maximum sulfur content.
To prevent hardening cracks, the part made from steel containing carbon in the amount of <3 weight % is heated to ensure the actual austenite grain size of not greater than #6.
To prevent hardening cracks, the part made from steel containing carbon in the amount of <.3 weight % is heated to ensure the actual austenite grain size of not greater than #11.
To prevent hardening cracks, the part made from steel containing carbon in the amount of <.3 weight % is heated to ensure the actual austenite grain size of not greater than #8, cooling is done with self-tempering at 150-300° C. for 1.0-30 seconds.
To prevent hardening cracks, the part made from steel containing carbon in the amount of <0.3 weight % is heated to ensure the actual austenite grain of not greater than #8, cooling is done with multiple self-tempering at 150-300° C. for 1.0-30 seconds.
After hardening, the part is subjected to low tempering during heating in the furnace at 150- 300° C.
After hardening, the part is subjected to tempering at over 300° C., i.e. the above the temperatures of total decomposition of martensite within the surface hardened layer into a thin quasi-eutectoid structure—troostite, troostosorbite, sorbite of tempering to HRC 25-50, with perservation of a pearlite-sorbite hardening structure in the core.
The specialty of the proposed steel thermal treatment process is that ultimate concentration, in the steel, of some indicated permanent admixtures that drastically increase hardenability can be limited to 0-.005% or 0-.1%, whereas that of other, weaker ones, can be increased to the range of 0-.3% and, in some cases, to 0-.5% without deterioration in quality. The simplifies selection of the initial charge during smelting and makes steel less expensive, since the final objective is to obtain the pre-specified calculated DI value by combining the composition of the residual admixtures after steel deoxidation with the amount of alloying elements added as per formula (1) that is based on the classical method of calculating hardenability per Grossman [4] and which turned out to be most acceptable for LH and SH steels. Practical experience confirmed its veracity for parts of various shapes and sizes.
However, this calculation has been subjected to refining by the authors in view of its potential further development. Thus, in the formula, the range of the K-coefficient that depends on the austenite grain size was extended up to ##11, 13. Multiplecands for hardenability as a function of tungsten and boron were additionally introduced. Multiplicands were introduced for the ideal diameter as a function of modifying elements of secondary carbonitride phases that do not enter the austenite solid solution prior to hardening cooling, i.e. titanium, vanadium, aluminum, carbon present in the structurally-free cementite of hypereutectoid steels, and suphur and phosphor were excluded from the formula since their content in the above-shown amounts does not have any practical effect on the ideal diameter (DI) value.
Under the circumstances, primarily from the economic point of view, addition of the weighed amount of manganese is expedient, as the most effective and relatively inexpensive component, alone, or along with inexpensive silicon in the amounts of 0-1.8% of each instead of more expensive ones that had been unreasonably added to the steel earlier only with the purpose of increasing hardenability.
Qualitative addition of boron to the steel in abnormally small amounts of .03-.005% also leads to increase in hardenability, which becomes even more effective as the content of carbon in steel decreases (see formula 1). Using the formula during the development of the steel chemical composition provides for optimal, rather than excessive, alloying steel.
Therefore, addition of other alloying elements—Cr, Ni in the amount of 0-0.5% only with the purpose of bringing steel hardenability (DI) to the required level of their presence in the form of permanent admixtures is less reasonable, since it will not practically change the mechanical properties compared to their lower content or absence with the same DI value and austenite grain size.
Addition of modifying elements, weight %, —titanium in the amount of not more than .4, vanadium—not more than .4, aluminum—.03-.1, nitrogen—not more than .1, present in steel in the form of finely dispersed carbides and nitrides insignificantly dissolved in austenite helps to reduce grain size, widen the optimal temperature range when heating prior to hardening, improve strength and plastic properties of the 3rd generation LH and SH steels. In this case, the total content of manganese and titanium should be more than size times the maximum sulfur content, since titanium, like manganese, binds sulphur into high-melting sulfides.
Addition of other alloying elements, weight %, —chrome, nickel in the amounts of not more than .6 (not more than 1.8 of each), molybdenum and tungsten (not more than .5 Mo and not more than 1.5 W), individually or together, selectively (as complex), also in compliance with the above formula in order to achieve a specified calculated DI value and improve quality properties, i.e. improve mechanical properties, heat resistance, lower the cold brittleness threshold, etc.
Given below is the justification of the chemical composition of LH and SH steels used for the proposed hardening process. The ultimate manganese content of 1.8 weight % is determined by steel susceptibility to overheating when its content is higher; excessive content of silicon above 1.8-2.0 weight % is fraught with steel changing from perlite class to ferrite class that is insusceptible to strengthening by hardening; chrome content for perlite class steels is not in excess of 1.8-2.0 weight % either due to the higher brittleness of the martensite-structured hardened layer; the ultimate content of nickel in the amount of 1.8-2.0 weight % is selected based on its high cost compared to that of manganese, silicon, chrome and relative low increase in its hardenability factor (.52). Besides, austenite grain size reduction to 10-13 in the inventor's application presented for the steel results in a substantial increase in plasticity and viscosity which excludes the role of nickel with subsequent increase in its quantity; copper is usually a practically non-removable admixture, its maximum content is usually limited to .25&, which, for the proposed steel, can, in some cases, be increased to .3-.5 weight % without affecting the quality and taken into account during weighed alloying of steel; molybdenum and tungsten are expensive components, as well, and are also added to the steel in weighed amounts, along with chrome and nickel mostly to increase its heat resistance; exceeding the content limit—.5 weight % for molybdenum and 1.5 weight % for tungsten can, even with small amounts of manganese and chrome, result in the steel changing to the martensite class, i.e. in through-hardening irrespective of the part size.
Sulphur and phosphor present in steel in the above-mentioned amounts do not practically affect the ideal diameter (DI) value.
Presence of carbide-forming elements—titanium and vanadium in the above-mentioned amounts, as well as aluminum and nitrogen that form aluminum nitride—contributes to formation of a finer austenite grain, inhibits its growth when subjected to heating prior to hardening and lowers hardenability. Under the circumstances, the lower aluminum content boundary of .03 weight % guarantees rather complete steel deoxidation, exceeding the upper limit of .1 weight % is unreasonable due to the beginning of aluminum dissolution in austenite and uncontrollable growth of steel hardenability making it more expensive.
Exceeding the ultimate content of nitrogen in steel—.1 weight % will cause an irreversible coagulation (coarsening) of aluminum and titanium nitrides which is analogous to titanium and vanadium carbides with the content of these elements above .4 weight %, all this also making steel more expensive.
The minimal ideal diameter value of 6 mm was obtained by the authors experimentally on LH40 steel of the following chemical composition, weight %: 41 C; .03 Mn; .04 Si; .06 Cr; .05 Ni, .3 Cu, .05 Al; .22 Ti (the calculated DI value is 5.4 mm for #13 grain austenite).
This steel ensured formation of a 1.3 mm through-surface hardened layer on a 10 mm dia cylinder-shaped rod, 1.55 mm layer on a 15 mm dia sphere, 1.3 mm layer on an 8 mm thick plate, which practically was in compliance with the calculated data of 1.2 mm, 1.3 mm and 1.1 mm.
Further decrease in the DI to <6 mm leads to a drastic growth in the critical hardenability velocity (Ver) to over 1500° C./sec., resulting in use of extremely pure steels that are free from permanent admixtures which is very difficult.
The most important distinctive features of steels proposed in this method is that during the development of the chemical composition of steel to be subjected to hardening it is possible to theorectically predetermine with sufficient accuracy not only the DI value, but also the optimal hardened layer depth,. as applied to the specific part in the shape of a cylinder, sphere or plate.
For parts exposed to bending or twisting loads during operation, the hardened layer depth equal to .1-.2 of diameter (thickness), i.e. the zone between sections I-I and II-II, see
However, quite often the required hardened depth can be either <.1 D (δ), i.e. 1.0 mm÷.1 D (δ) (the zone located below section I-I (see
Thermo-physical calculation results shown on the graph (see
As far as crack formation is concerned, practical research showed that for all steels with the actual austenite grain not coarser than ♯11 mentioned in this invention, crack formation does not take place. This is true with respect to both surface-hardened sections of the part with compression residual stresses within the hardened layer that lower the sensitivity to crack formation, and separate zones with through-depth hardening (thin splines, thread, keyslots, end faces of thin bushings, etc.) with stretching residual local tension stresses in the surface layers, that greatly contribute to crack formation and propagation. The reason lies in the fact that the tension stresses that occur during formation of fine-structure martensite are reliably offset by its high brittle strength.
For steels with .15-.30% carbon content, due to lower tensile strength of the low-carbon martensite featuring some plasticity even in the untempered state, experimental research made it possible to expand the range of the actual austenite grain growth to ♯6 size, where cracks are not present.
For steels with a higher carbon content (>0.3%), with no local sections in parts with through-depth hardening the maximum actual austenite grain size at which crack formation does not take place corresponds to ♯8, provided that steel was self-tempered at 150-300° C. for 1.0-30 sec.
For relatively massive parts (>100 mm dia) made from steels with increased carbon content, single-, doubhle- or tripple self-tempering is recommended for reliable prevention of crack formation.
Low tempering while heating in a furnace at 150-300° C. is performed to ensure final stabilization of the steel part microstructure and properties.
Medium and high furnace tempering of parts is carried out at temperatures over 300° C. is performed to ensure final stabilization of the steel part microstructure and properties.
Medium and high furnace tempering of parts is carried out at temperatures over 300° C., i.e. above temperatures at which martensite totally decomposes within the surface-hardened layer into a thin quasi-eutectoid structure—troostite, troostosorbite, temper sorbite with HRC 25-50, and preserving the pearlite-sorbite hardening structure in the core. This tempering is carried out for parts requiring a combination of strength, higher plasticity properties and surface impact strength—shafts, axles and similar parts.
Shown below are Tables 1-3 listing chemical compositions of LH and SH steels with their ideal diameters and examples of technological processes based on the proposed process.
Example .1 LH steel with the following chemical compositions, %: .78 C; .04 Mn; .08 Si; .07 Cr; .15 Ni; .08 Cu; .04 Al; .15 Ti; .015 S; .018 P has a calculated DI=.78 mm when treated for ♯12 grain size. According to the graph, (see
An 8-mm wall roller bearing ring made from this steel was through-heated in an induction furnace up to 850° C. for 20 sec and then subjected to cooling with a sharp water shower and tempered in a furnace at 150° C. with soaking for 2 hours.
As a result, the hardened layer on the outer and inner surfaces of the ring was 1.7 mm and 1.5 mm, i.e. .18-.2 of the wall thickness, which corresponds to the through-surface hardening (TSH); the hardened layer microstructure was cryptocrystalline martensite martensite (♯1 grain size), hardness was HRC65-66, in the core—troostite, troostosorbite, sorbite with HRC 38-48 hardness.
Example 2, SH steel with the following chemical composition %: .61 C; .5 Mn; .08 Si; .13 Cr; .25 Ni; .03 Cu; .04 Al; .05 Ti; .015 S; .018 P has a calculated DI=22.5 mm when treated for ♯11 grain size. According to the graph (see
A 45 mm dia cylindrical center pin made from this steel was through-heated in an induction furnace up to 900° C. for 50 seconds and then subjected to cooling with a sharp water flow and tempered in a furnace at 180° C. with soaking for 2 hours. As a result, the hardened layer depth on the part surface was 5 mm, i.e. .11 of the diameter, which corresponds to the graph (see
A 30 mm dia grinding ball made from this steel was through-heated in a furnace up to 850° C. and then subjected to cooling with a sharp water flow with self-tempering at 180° C. for 5 seconds, followed by final cooling with a water flow. As a result, the hardened layer on the part surface was 12 mm, i.e. .4 of the diameter, which corresponds to hardening close to through-depth hardening and to
Example, 3. SH Steel with the following chemical composition, %: .5 C; .1 Mn; .15 Si; 1.0 Cr; .8 Ni; .03 Cu; .05 Al; .35 V; .5 W; .015 S; .018 P has a calculated DI=47 mm when treated for ♯10 grain size. According to the graph (see
A 150×200×200 mm parallelepiped-shaped part made from this steel was through-heated in a furnace up to 850° C. and then subjected to cooling with a sharp water flow with double self-tempered at 180° C. for 5 seconds, followed by final cooling with a water flow and tempering in a furnace at 450° C. with soaking for 3 hours. As a result, the hardened layer along the part surface perimeter was 9 mm, i.e. .06 of the thickness (150 mm), which corresponds to the FIG. 3b graph; the microstructure of the hardened layer was in the form of temper troostite, hardness was HRC48, in the core—hardened troostosorbite, hardness HRC48-50.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2010154543 | Dec 2010 | RU | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/RU11/00279 | 4/28/2011 | WO | 00 | 10/6/2014 |
