STUDY OF AUTOMATICALLY DEPOSITED HIGH-SPEED STEEL

I. G. Lyapichev

(Presented by Prof. A. N. Dobrovidov, Doctor of Technical Sciences)

The advantage of using certain types of deposited tools in comparison with tools made of forged or rolled high-speed steel is well known. To further increase the effectiveness of introducing this new method of tool manufacture into industry, further investigation is needed of questions connected with automatic deposition of cutting tools. This was noted in the resolution of the conference on cast and deposited tools held in 1950.

One such question is the study of new grades of deposited high-speed steels whose cutting properties are higher than the cutting properties of the standard grades R9 and R18.

The aim of the work is to study the technology of automatic deposition and deposited steels of the R9 type, additionally alloyed with carbon, vanadium, molybdenum, titanium, and tungsten.

Standard high-speed steel is a proven basis for obtaining new steel grades. Additional alloying of R9 steel should lead to an increase in cutting properties, since the completeness of utilization of the alloying elements—the saturation of austenite under the conditions of obtaining deposited steel—increases. Moreover, along with quenching from the highest temperatures—from the solidus-line temperature and above—grain growth does not occur. The crystallization conditions of deposited high-speed steel prevent the formation of large accumulations of carbides, which ensures the necessary toughness, sufficient for its service.

Chemical Composition and Technology of Experimental Deposits of High-Speed Steel

The high cutting properties of a deposited tool are ensured by high hardness after tempering, red hardness, wear resistance, and satisfactory impact toughness. These requirements are met by the cast structure of high-speed steel with a favorable arrangement of primary carbides and a highly alloyed solid solution.

The service conditions of a cutting tool, and consequently the requirements imposed on it, are varied. Experience shows that the composition of the standard steels R9 and R18 is not the most rational for cast and, consequently, deposited tools \[1, 2, 3, 4, 5, 6\].

Studies of the influence of carbon, vanadium, molybdenum, and titanium on the properties of deposited high-speed steel have not been carried out in the works known to us. In a number of works \[13, 12\], only the technology of obtaining the chemical composition of deposited metal close to that of steels R9 or R18 was studied.

A preliminary study of various methods of alloying the deposited metal with a ceramic flux and with bars of ferroalloys (electrode wire of grade Sv1) showed that the most convenient method is the latter—alloying the deposited metal with bars of ferroalloys [7].

The chemical composition of the deposits is shown in Table 1. Basically, this is R9 steel, additionally alloyed with carbon (Group I), titanium (II), vanadium (III), molybdenum (IV), molybdenum and vanadium (V), vanadium, molybdenum and titanium (VI), and tungsten (Group VII).

The metallurgical features of surfacing (considerable overheating, “incomplete”—without Al—deoxidation) and of the crystallization process make it necessary to raise the question of the composition of the deposited high-speed steel.

A study of the microstructure of automatically deposited high-speed steel showed the following.

An increase in the carbon content in the deposits of Group I leads to an increase in the amount of eutectic and of the white constituent—retained austenite. The hardness after surfacing is thereby reduced. The alloying of austenite with carbon apparently increases, which promotes an increase in the solubility of other elements.

At a carbon content of 1.14%, the austenite decomposition temperature rises to 600°. When the content is reduced to 0.72% carbon (Fig. 1), areas of the light constituent are observed inside the grain. Etching for carbides does not color these areas. At a content of more than 1.14% C, the main field is occupied by retained austenite—the light constituent. The carbides in the eutectic occur in the form of a chain, and not as a continuous interlayer.

Fig. 1. Structure of high-speed steel after surfacing and air cooling.
(Deposit No. 1. ×500).

Titanium, with a simultaneous increase in the carbon content, was introduced in order to change the structure and alloy the solid solution. At 0.15% Ti and 1.04% C, coagulation of the eutectic carbides is observed; special inclusions appear within and along the grain boundaries. With slight etching of the polished section, these inclusions have a light-blue tint. The inclusions are uniformly dispersed. With an increase in the titanium content to 0.5% and carbon to 1.4% (Fig. 2), the number of inclusions increases. In addition, martensitic crystals are visible against the background of the light constituent. The hardness of the deposits of Group II decreases in comparison with the deposits of Group I, despite, for example, the same carbon content. Determination of red hardness confirms the assumption that the appearance of martensite crystals is a consequence of

Table 1

(Continuation of Table 1)

decomposition of the depleted austenite. Depletion of the austenite in carbon occurs owing to the introduction of titanium.

However, the alloying of the austenite is sufficient to ensure an increase in hardness on tempering from 58 to 63 \(H_{Rc}\).

Fig. 2. Structure of high-speed steel after surfacing and cooling in air. (Surfacing No. 13. \(\times 500\)).

The introduction of titanium worsens the formation of the bead and leads to the appearance of undercuts. Apparently, this is explained by an increase in the viscosity of the steel. This is confirmed by the sharply defined outline of the “last” pool at the end of the bead.

No influence of titanium on the size of the austenite grain was observed.

An increase in the vanadium content leads to coagulation of the carbides of the eutectic, which in places is broken up. Within the grain of the light constituent, secondary carbides are visible under intensified etching.

At a content of 1.39% C and 5.98% V, readily etched central regions of grains, penetrated by carbides, can be seen against the background of the light constituent. The effect of vanadium on the form of carbide occurrence and on the increase in alloying of the solid solution is analogous to the effect of titanium, but somewhat weaker.

An increase in the molybdenum and carbon content in deposits of group IV promotes an increase in the amount of eutectic and of the light constituent. This leads to an increase in hardness after surfacing. The hardness after tempering increases only slightly.

At 1.42% C and 3.02% Mo (Fig. 3) the metallic matrix is penetrated by a carbide skeleton. Accumulations of eutectic are observed at the grain junctions. In testing the cutting properties, this deposit showed increased durability, which is ensured by sufficient red hardness and high wear resistance. When the feed was increased above 0.15 mm/rev on cutoff tools, brittleness was observed.

The simultaneous introduction of molybdenum (to increase red hardness) and vanadium (for coagulation of carbides) in deposits of group V did not give the expected results. Apparently, the alloying of the solid solution in this case rules out the possibility of its decomposition. The hardness and microstructure

after tempering do not change: against the main background of the light constituent, carbides are distributed.

Fig. 3. Structure of high-speed steel after surfacing and cooling in air. (Surfacing No. 15. ×500).

The introduction of titanium into molybdenum steels instead of vanadium proved more favorable for increasing hardness during tempering and for carbide coagulation. With 0.94% C and 0.2% Ti (Fig. 4), and 1.59% C and 0.5% Ti (Fig. 5), the most favorable structure is observed. At 0.3–0.4% Ti, the primary carbides occur in the form of isolated plates, while the secondary carbides are distributed uniformly throughout the light constituent. Titanium compounds are observed as isolated inclusions in the eutectic. An increase in the carbon and titanium content leads to a decrease in the easily etched regions. The main field is occupied by the light constituent with uniformly dispersed carbides.

Fig. 4. Structure of high-speed steel after surfacing and cooling in air. (Surfacing No. 10. ×500).

High-tungsten deposits are characterized by an increase in the amount of eutectic. Deposits 20 and 21 are close in composition to R18 steel. Increasing the carbon content in them from 0.78 to 1.15% increases the amount of the light constituent; the hardness after deposition decreases. No loss of hardness is observed at tempering temperatures of 300–500°.

Fig. 5. Structure of high-speed steel after deposition and cooling in air. (Deposit No. 18. ×500).

With an increase in vanadium content from 3.39 to 3.51%, coagulation of carbides is not observed. Located along the dendrite boundaries, the carbides assume a distinctive shape (Fig. 6). The increased alloying of the solid solution leads to a further decrease in hardness after deposition and to an increase in the tempering temperature. In testing the cutting properties, brittleness was not observed.

Fig. 6. Structure of high-speed steel after deposition and cooling in air. (Deposit No. 23. ×500).

Structural Zones in the Deposited Metal

Studies of the temperature field in the liquid metal of the pool are not known to us.

The result of nonuniform cooling of the deposited layer is the presence of three structural zones along the height of the bead section.

1. The zone adjacent to the workpiece consists of fine crystals that grew under conditions of strong heat removal.
2. The zone of columnar crystals (Fig. 7). Dendrites grow in the direction of greatest heat removal and slow their growth as it decreases.
3. The zone of randomly oriented crystals (Fig. 8).

Fig. 7. Structure of high-speed steel after deposition and cooling in air. (Deposit No. 23. ×200). On the left, the zone of fusion with the base metal is visible.

Fig. 8. Structure of high-speed steel after deposition and cooling in air. (Deposit No. 23. ×200). In the left part the end of the zone of columnar crystals is observed. On the right is the zone of unoriented crystals.

The degree of development of each zone depends on the composition of the deposited layer, since the surfacing and cooling conditions are identical. The extent of the first zone does not exceed 0.10–0.15 mm, that of the second 4–4.5 mm, and that of the third 2.25–2.50 mm. The grain size and the form of occurrence of the eutectic are not constant over the cross section of the deposit. Fluctuations in grain size reach 30–40%. However, the hardness of the light constituent in the different zones of the bead cross section remains constant. Microhardness was measured under a load of 20–50 g on a PMT-3, at intervals of every 0.25 mm along the height of the section. The hardness in the different zones of the deposit, measured under a load of 200 g, also remains unchanged.

Microhardness measurements (under a load of 200 g) were carried out on specimens with different heat treatments: immediately after surfacing and cooling in air, and also after high-temperature tempering.

These data, together with layer-by-layer (by height) chemical analysis, show that the deposited layer is homogeneous. Such homogeneity is apparently explained as follows: under steady-state process conditions, each layer of the bead solidifies under identical conditions. The composition of the liquid phase thereby remains unchanged, in contrast to the crystallization of a casting in an ordinary mold. Continuous oscillation of the liquid metal and the presence of convection currents lead to equalization of the bath composition. The high cooling rate and the small volume of solidifying metal eliminate zonal segregation.

Tempering of Deposited High-Speed Steel

The study of high-speed steel during tempering begins with an investigation of the change in hardness after one-hour tempering at increasing temperatures. In this way the optimum tempering temperature is determined, one that causes vigorous decomposition of retained austenite and an increase in hardness. In such investigations, one specimen is usually used over the entire range of tempering temperatures, from 100 to 600–650°.

A. P. Gulyaev [8, 9, 10] investigated structural transformations during the tempering of forged high-speed steel as a function of the quenching and tempering temperatures, the holding time at these temperatures, and the number of tempering cycles.

Repeated tempering ensures a more complete decomposition of austenite than single tempering. In this case a higher alloying of the secondary martensite is observed [10].

Changes in hardness during tempering of experimental deposits by the indicated method [10] show that deposited and quenched forged high-speed steels behave similarly during tempering. It is necessary to note the absence, or reduction, of the drop in hardness in the range 300–500°. This is apparently explained by reduced work-hardening of the austenite because of the weak development of the martensitic reaction and by the influence of tensile stresses in the deposited layer.

The weak development of the martensitic reaction is the result of the high alloying of austenite when it is quenched from solidus-line temperatures. The austenite thereby acquires only slight work-hardening from the course of the martensitic reaction. The compressive stresses in the austenite, arising because of the decomposition of neighboring grains, are reduced by tensile stresses in the deposited metal.

The influence of tensile stresses is confirmed during tempering of a deposit of the same composition, which was made on a workpiece previously cut into elements of 5–6 mm each and then reassembled. Such a workpiece provides “free” shrinkage of the deposited layer and prevents the appearance of tensile stresses. During tempering of this deposit in the range 300–500°, the usual drop in hardness is observed.

The experiments carried out confirm the hypothesis of A. P. Gulyaev [10] on the origin of the hardness drop in the range of 300–500°.

The foregoing makes it possible to regard the cooling of deposited steel in air as quenching from the highest possible temperature. In this case there is no growth of the austenite grain.

It is known that the saturation of austenite with alloying elements during quenching from the solid state is determined by the rate of dissolution of carbides in austenite during heating and holding.

An increase in the quenching temperature or in the holding time at that temperature leads to an increase in the alloying of austenite. However, the growth of the austenite grain or melting limits the holding time at that temperature.

The conditions for saturation of the solid solution when it precipitates from a highly alloyed liquid steel are more favorable than its saturation at the usually adopted quenching temperatures.

The composition of the carbide precipitate [9] shows that the content of elements in the liquid phase is 2–3 times higher than their average content in the steel. Intensive convection currents continuously equalize the composition of the liquid phase—the weld pool. Rapid cooling after crystallization ensures the highest alloying of the solid solution. At the same time, it provides better saturation with vanadium, titanium, and other elements whose carbides dissolve insufficiently at a quenching temperature of 1240–1290°. Growth of the austenite grain of the deposited steel does not occur, since cooling reduces the reserve of free energy. Growth of the austenite grain during prolonged holding and at quenching temperatures of 1240–1290° occurs with the expenditure of energy, as an endothermic reaction. During rapid cooling after surfacing, growth of the austenite grain is retarded, as an endothermic reaction during cooling.

We consider it necessary to make some comments on the research procedure. Under normal quenching, high-speed steel of the P18 type contains 18% carbides, 29–37% retained austenite, and 54–45% primary martensite [10]. The greatest decomposition of austenite occurs during the first tempering; little of it remains for the subsequent ones [10].

Apparently, double tempering, with the correct choice of temperature, can ensure a sufficiently complete decomposition of austenite and relief of stresses.

Holding at the tempering temperature leads to precipitation of carbides from austenite. During cooling, the following occur: decomposition of impoverished austenite into secondary martensite (increase in hardness) and decomposition of both primary and secondary martensite from the preceding tempering (decrease in hardness). Tempering regimes must promote the development of the first transformation and retard the second.

The studies of P. M. Yushkevich [11] show that the composition of the austenite of cast high-speed steel of the P18 type does not change at tempering temperatures up to 560°; an increase in the tempering temperature leads to a decrease in the alloying of austenite because the diffusion paths are shortened, owing to the penetration of the austenite grain by martensitic crystals. The presence of martensite accelerates the decomposition of austenite, which is manifested especially during repeated tempering.

When the tempering temperature is raised above 560°, the decomposition of retained austenite depends not only on the tempering temperature but also on the amount of martensite formed during the preceding temperings or quenching. The higher the temperature and the greater the number of temperings, the more intensive is the decomposition of austenite.

This influence is not taken into account when a single specimen is used in studying all tempering temperatures. The intensity of decomposition is considered to depend only on the tempering temperature. The preceding temperings are not taken into account.

Figure 9 shows the change in hardness after tempering at increasing temperatures for one and the same specimen and for “fresh” specimens.

During tempering in the temperature interval \(100\text{--}560^\circ\), the hardness of a “fresh” specimen is lower than the hardness of a specimen that has already been tempered. At a tempering temperature above \(580^\circ\), the hardness of “fresh” specimens exceeds the hardness obtained when tempering a single specimen. This is explained by the influence of the preceding tempers when one and the same specimen is used.

Fig. 9. Change in hardness \(H_{Rc}\) of high-speed steel (after surfacing and cooling in air) as a function of tempering temperature (surfacing No. 7). Solid line—tempering of one specimen. Dashed line—tempering of fresh specimens for each tempering temperature.

The actual temperature of intensive decomposition of austenite during tempering of “fresh” specimens is \(50^\circ\) higher than the optimum temperature.

The investigations of P. M. Yushkevich [11] show that depletion of austenite, its transformation into the \(\alpha\)-phase, and the decomposition of martensite into a ferrite-carbide mixture proceed isothermally during tempering.

The rate of transformation of austenite into the \(\alpha\)-phase increases with multiple tempering owing to the influence of martensite.

A lowered “optimum” tempering temperature leads to an increase in the number of tempers required to obtain maximum secondary hardness after the decomposition of austenite. Increasing the number of tempers is uneconomical and, in addition, leads to intensified decomposition of martensite into a ferrite-carbide mixture with each subsequent temper.

The foregoing shows that hardening to secondary hardness with the smallest content of “primary” martensite is the most expedient.

These requirements are satisfied by hardening the deposited metal by cooling in air from the solidification temperature. Tempering may be no more than double. The temperature and holding time during the first temper must ensure the most complete decomposition of retained austenite through the martensitic reaction.

The second temper is necessary for relieving stresses from the martensitic reaction of the first temper and for completing the decomposition of the work-hardened austenite enclosed in the spaces between the martensite crystals. The lowered temperature of the second temper must not cause further decomposition of the martensite of the first temper into a ferrite-carbide mixture.

The work-hardened austenite in the spaces between the martensite crystals will decompose during the second temper despite its low temperature. This is promoted by the depletion of austenite, the presence of precipitated carbides—additional crystallization centers, and the presence of new interfaces—martensite crystals.

The correctness of the choice of the second-temper regime is checked by a third temper at the same temperature. The hardness and the readings of the austenitometer galvanometer should remain unchanged.

To confirm these assumptions, 8 specimens were made from forged annealed R9 steel, and 4 plates of size [[unclear: continuation cut off at page bottom]] from surfacing No. 7.

$2 \times 20 \times 25$ mm, with identical deviations. Plates from the deposited metal (without tempering) were cut out with an abrasive wheel under abundant cooling. The absence of burns during cutting of the deposits and grinding was checked by measuring hardness and by visual inspection. Seven plates of wrought steel were quenched in a salt bath according to the plant procedures (quenching temperature 1240–1260°, holding time—1 min.).

The various tempering regimes for wrought and deposited high-speed steel, and their influence on red hardness and cutting properties, are shown in Table 2.

Tempering of the plates was carried out in a salt bath. Temperature fluctuations reached $\pm 5^\circ$. The plant tempering was threefold, at a temperature of 550°.

A plate of annealed wrought P9 steel served as the standard for the austenitometer.

The optimum tempering temperature for all deposits, with the exception of the seventh group, is the same and reaches 580°. This makes it possible to elucidate the influence of the steel composition on red hardness. Increasing the carbon content (group 1) from 0.72 to 0.96–1.14% leads to an increase in hardness and to its retention with a larger number of tempering operations. At a content of 1.38% C, apparently, saturation of the solid solution occurs, which leads to an increase in the amount of retained austenite. The hardness at a tempering temperature of 580° does not rise above 62 $H_{\mathrm{RC}}$ and is retained for 4 temperings.

Apparently, the retarding effect of the increased amount of eutectic makes itself felt. During 2 additional temperings at a temperature of 600°, the hardness rises to 63.5 $H_{\mathrm{RC}}$.

It is characteristic to note that if tempering is begun immediately at a temperature of 600°, then the hardness after the first tempering rises to 63 $H_{\mathrm{RC}}$, and after the second decreases to 59 $H_{\mathrm{RC}}$.

Increasing the carbon content in the deposits of group 1 leads to an increase in red hardness.

The introduction of 0.15% titanium at a content of 1.04% C and of 0.3% Ti at 1.18% C causes an increase in red hardness and spheroidization of carbides in comparison with deposits Nos. 19 and 2 of the same composition but without titanium. The introduction of titanium into deposited steel at an increased carbon content is expedient. This is confirmed by testing the cutting properties of these deposits. Increasing the titanium content to 0.5% at 1.40% C (deposit No. 13) showed a decrease in red hardness to 600°.

Apparently, the content of carbon, carbide-forming elements (W, Cr, V, Mo, Mn), and titanium in deposits Nos. 3 and 4 does not lead to depletion of the solid solution. In deposit No. 13, the content of 0.5% Ti leads simultaneously to an increase in the amount of the carbide phase and enrichment of the solid solution. The latter causes a decrease in hardness after deposition to 58 $H_{\mathrm{RC}}$ and an increase in it, with repeated tempering, to 63 $H_{\mathrm{RC}}$.

The introduction of titanium leads to a certain stabilization and decrease in the hardness of high-speed steel after deposition. However, during tempering, deposits alloyed with titanium increase in hardness because of the subsequent decomposition of the solid solution. In deposits alloyed with vanadium, no increase in hardness during tempering is observed.

Increasing the vanadium content from 2.56 to 5.98% (group III) does not cause an increase in red hardness or cutting properties.

Apparently, under deposition conditions, the alloying of austenite with vanadium (in deposits Nos. 2, 12: 2.6–2.8%) is sufficient to raise the cutting properties of the deposits to the level of P18 steel. A further increase in the vanadium content leads to excessive stability of the austenite, which does not produce an increase in cutting properties. Increasing the vanadium content is favorable for wrought [14] and sand-cast [5, 15] high-speed steels. Vanadium prevents growth of the austenite grain at elevated quenching temperatures. In this case,

the alloying of the austenite, its red hardness, and its cutting properties increase.

Under conditions of deposition, the above-listed requirements are satisfied by the conditions of crystallization and hardening during cooling in air after deposition. Therefore, increasing the vanadium content to more than 3% in medium-tungsten deposited steels is excessive.

The introduction of molybdenum (group IV): 0.48% Mo at 0.92% C; 1.73% Mo and 1.16% C; 1.42% C; and 3.02% Mo—increases red hardness. The greatest increase in red hardness is observed in deposits six and seven. The increase in red hardness is accompanied by a corresponding increase in cutting properties. The highest durability was shown by deposit No. 7. In deposit No. 15, chipping of the cutting edge was observed.

Deposits of group V are analogous to deposits of group III and are additionally alloyed with molybdenum. The high vanadium content in deposits of group V is the cause of their low red hardness, which is characteristic of deposits of group III. Additional alloying with molybdenum apparently does not change the composition of the solid solution. The decomposition of austenite during tempering of these deposits does not cause an increase in hardness.

The results of simultaneous alloying with molybdenum and titanium are represented by deposits Nos. 10, 11, 17, and 18. The aim here was to increase red hardness by introducing molybdenum and to spheroidize the carbides with titanium.

The red hardness of deposits No. 10 (0.94% C, 0.63% Mo, and 0.2% Ti) and No. 11 (1.15% C, 2.64% Mo, and 0.3% Ti) decreases in comparison with deposits without titanium: No. 6 (0.92% C and 0.48% Mo) and No. 7 (1.16% C and 1.73% Mo). The introduction of titanium without an increase in carbon content reduces the content of the latter in the austenite.

At contents of 1.39% C and 0.4% Ti (deposit No. 17); 1.59% C and 0.5% Ti (deposit No. 18), red hardness increases, but brittleness also increases (especially in deposit No. 18).

Deposit No. 20 corresponds in composition to steel of the R18 type. Increasing the carbon and tungsten contents (deposit No. 21, 1.15% C and 21.46% W) improves their cutting properties.

The introduction of vanadium up to 3.5% at a carbon content up to 1.24—1.40% leads to a significant increase in red hardness and cutting properties. Toughness remains satisfactory.

Determination of the Cutting Properties of Deposited High-Speed Steels

Cutting tests were carried out on all deposits, with the exception of Nos. 8, 9, and 16, whose hardness and red hardness were not sufficiently high.

The tests were carried out in the experimental shop of State Bearing Plant No. 5 under operations and conditions established for cutters made of rolled R18 steel.

For the test, 2 straight turning tools and 4 cut-off tools were made from each deposit. The turning tools were obtained after cutting at an angle of 45° (in plan) in the middle of the deposited blank. The blank for deposition for cut-off tools was slotted to a depth of 10–12 mm along its entire length. This facilitated the subsequent manufacture of cut-off tools.

The conditions and results of testing the cut-off tools are presented in Table 3. The tools were sharpened in the shop. The tools were not subjected to honing.

The same durability ratio was obtained in comparative tests of straight turning deposited tools made of rolled R18 steel.

Tests of cut-off tools made from deposits Nos. 19, 7, and 21 on ShKh-15 (hardness \(H_B = 190—205\)) showed increased durability, respectively,

130–195–180% of the life of cutters made of rolled R18 steel. Cutting conditions: \(S = 0.10\ \text{mm/rev}\), \(V = 30\ \text{m/min}\) with cooling. Reducing the blade width to 1–1.5 mm resulted in a saving of 10–12% of ShKh-15 steel.

Table 3

Tests of cut-off cutters \((B = 5\ \text{mm})\) made from surfacings 19, 4, 7, and 21 were carried out on steel 45 in making a rectangular thread for a lead screw (Vakhruchev Electromechanical Plant, MUGP). The life of the cutters proved to be twice as high as the life of cutters made of forged R18 steel of factory manufacture. Cutting conditions: \(S = 0.12\ \text{mm/rev}\), \(V = 21\ \text{m/min}\) without cooling. Increasing the cutting conditions was hampered by the insufficient rigidity of the screw.

The testing of three-sided milling cutters with surfaced knives was carried out at the Tomsk Order of the Red Banner of Labor Cutting-Tool Plant.

Cutting conditions: \(V = 44\ \text{m/min}\), \(t = 11\ \text{mm}\), \(S_z = 0.13\ \text{mm/tooth}\) with cooling. Work material—hammered blanks \(60 \times 60 \times 500\ \text{mm}\) of steel 45 \((H_B = 195—224)\). Blunting criterion—wear on the flank face \(0.6—0.8\ \text{mm}\). Comparative tests showed that the life of cutters with knives, surfac—

deposited with steel No. 7 is 95–124% of the life of cutters with knives made entirely of P18 steel.

The deposited knives were subjected only to triple tempering at a temperature of 580°. The sharpening of cutters with deposited and factory-made knives was identical.

Conclusions

1. Deposits of medium-tungsten steel (9–10% W) with varying contents of carbon, vanadium, molybdenum, and titanium, as well as high-tungsten steels (17–24% W) with varying contents of carbon and vanadium, were obtained and investigated.

2. During tempering of deposits that harden on cooling in air, the absence of, or only a slight, hardness drop in the temperature range 200–500° was observed. This is the result of the presence of tensile stresses in the deposited layer.

3. The existing method for determining the optimum tempering temperature by using a single specimen over the entire tempering-temperature range (from 100 to 600–650°) leads to intensive decomposition of the solid solution at lower temperatures. The actual temperature of intensive decomposition lies higher, which is explained by the absence of the influence of martensitic crystals from previous temperings. The presence of martensite accelerates the decomposition of austenite.

4. Investigation of various tempering regimes after high-temperature hardening of forged P9 steel and of a deposited layer hardened from solidus-line temperatures showed the effectiveness of double tempering. The first—high-temperature—tempering ensures intensive decomposition of the solid solution by a martensitic reaction. The second—low-temperature—tempering serves to relieve stresses and complete the decomposition of stressed austenite located in the spaces between martensitic crystals.

5. The red hardness of medium-tungsten steel is not increased by adding more than 3% vanadium. The greatest increase in red hardness and cutting properties is observed when 2–3% molybdenum is added. The addition of 0.3–0.4% Ti has a favorable effect on the structure and cutting properties, while the carbon content is simultaneously increased to 1.1–1.2%.

Saturation of austenite at a content of 3% V in the deposit is sufficient for a considerable increase in cutting properties. Apparently, it is advisable to alloy the steel with the largest possible number of elements. This promotes development and the fullest use of the dissolving capacity of the γ-iron lattice. At the same time, it provides the possibility of forming a complex carbide that is resistant to coagulation. As a result, the alloying of austenite, its red hardness, cutting properties, and toughness are increased.

1. Increasing the manganese content (to 1.1%) and silicon content (to 0.9%) did not affect the cutting properties of deposited medium- and high-tungsten steel. The identical manganese content in the carbides of the specimen after deposition and after tempering shows that manganese does not form an independent carbide, whose presence would be undesirable.

References

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