Volume 18 - Friction, Lubrication, and Wear Technology Part 23 pps

The retained austenite content should be controlled by adjusting the case carbon content, not by subzero quenching after carburizing or by tempering at temperatures above 200 °C 390 °F

Fig 11 Examples of laser-clad microstructures (a) Tribaloy T-800 alloy on ASTM A 387 steel (b) Haynes

Stellite alloy No 1 on AISI 4815 steel Source: Ref 33

The Fe-Cr-Mn-C cladding of Mazumder and Singh (Ref 41) had a microstructural consisting of M7C3 and M6C type carbides in a ferritic matrix, a solid solubility extension of chromium in ferrite by 50%, and a microhardness of 550 HV The molybdenum-clad coating on gray iron of Belmondo and Castagna (Ref 42) consisted of molybdenum dendrites surrounded by a Cr-Ni matrix containing Cr2C3, which had a microhardness of 900 HV

Wear Behavior of Laser-Clad Layers. Mazumder and Singh (Ref 41) found that the Fe-Cr-Mn-C cladding on AISI

1016 steel resulted in tribological properties that were superior to Stellite 6 The width of the wear scar was reduced from 3.5 mm (0.14 in.) in the base alloy to about 0.6 mm (0.02 in.) in the laser cladding Under the same test conditions, Stellite 6, a common hardfacing alloy, developed a wear scar which of 1.4 mm (0.06 in.)

Abbas et al (Ref 43) laser clad a mild steel with Stellite 6, Alloy 4815, and their composites with SiC by using pneumatic

powder delivery Wear tests, conducted by grinding the samples against a revolving alumina disk, showed that the composite clad samples had the best abrasive wear resistance (Fig 12)

Fig 12 Abrasive wear behavior of mild steel laser clad with Stellite 6, Stellite 6/SiC, and Alloy 4815/SiC

Source: Ref 43

Belmondo and Castagna (Ref 42) performed wear tests on the Mo-Cr2C3 clad cast iron samples using a reciprocating motion testing machine When compared with plasma-sprayed coatings of similar composition, the performance of laser-clad coatings was far superior under all testing conditions, especially at the highest pressures With a load of 25 MPa (3.6 ksi) and slider roughness of 0.3 m (12 in.), the wear loss in the laser sample was 0.4 mg (1.4 × 10-5 oz), compared to 1.3 mg (4.6 × 10-5 oz) in the plasma-sprayed sample In addition, the plasma-sprayed coating developed cracks at the interface, which seriously impaired its integrity

Molian and Hualun (Ref 44) performed pin-on-block reciprocating sliding wear tests on Ti-6Al-4V alloy substrate laser clad with BN, both with and without a NiCrCoA1Y addition, and found that the wear resistance improved from 10 to 200 times more than that of age-hardened and laser-melted samples (Fig 13) The presence of solidification reaction products, TiN and TiB2, resulted in a cladding microhardness of 1600 HV, which led to improved wear resistance by preventing ploughing and reducing friction

Fig 13 Sliding wear behavior of laser-clad Ti-6Al-4V alloy Source: Ref 44

Boas and Bamberger (Ref 45) determined the abrasive wear characteristics of sprayed and laser-melted sprayed coatings of Tribaloy T-400 on AISI 4130 steel using a block-on-cylinder test The plasma-sprayed coating was prone to rapid wear by delamination, whereas laser consolidation of the plasma coating removed flaws, such as porosity and microfissures, and generated an adherent wear-resistant layer Figure 14 demonstrates the improvement in wear after laser consolidation of the plasma coating For comparison, results on a hard D2 steel are also given

plasma-Fig 14 Wear scar growth curves of plasma-coated 4130 steel before and after laser consolidation Source: Ref

45

Laser-Cladding Applications. Because one of the primary goals of laser cladding is to improve the tribological properties of components, several applications have been found for the technique Some of them are in the exploratory stage, others are in the pilot-plant stage, and a few have reached the production stage

Bruck (Ref 46) demonstrated cladding within a confined space by coating the inside surface of small-bore pipes (inside diameter of 50 to 100 mm, or 2 to 4 in., and length of 0.3 to 1.2 m, or 1 to 4 ft), as shown in Fig 15 An oscillating mirror placed in the pipe formed a 12.5 mm (0.5 in.) wide melt pass, into which the clad alloy powder was fed The pipe was rotated and translated to cover the entire inner surface Previously, chrome plating was used, but the laser technique appreciably improved galling resistance Using suitable cooling techniques, the pipe temperature was maintained below

500 °C (930 °F) to avoid deterioration of the core properties, as well as distortion

Fig 15 Laser cladding of small-bore pipes Source: Ref 46

La Rocca (Ref 19) describes laser cladding of exhaust valves with stellite at an Italian automotive manufacturer The stellite powder was fed into the laser melt pool The laser-clad material was superior to gas tungsten arc welded material

in terms of thickness uniformity and uniform microstructure and elemental distributions, and had better adhesion Powder utilization also was better (30% that of gas tungsten arc welding), and over-metal removal was reduced by 10 to 15% One British automotive manufacturer uses production-stage laser cladding of a nickel-alloy turbine blade shroud interlock

by powder feeding either Tribaloy or Nimonics (Ref 47) Previously, a manual arc melting technique was used Laser cladding reduced cladding time from 14 min to 75 s, improved productivity and quality, reduced cost by 85%, and reduced powder consumption by 50% Table 2 lists components that are laser clad for commercial applications

Table 2 Component, cladding alloy, and cladding technique

technique

Gate and seat of steel valves for oil-field, geothermal,

and nuclear energy production, and chemical

Westinghouse, Pittsburgh, PA

52

Valve stem, valve seat, aluminum block CrC 2 , Cr, Ni, Mo/cast

Fe

Preplaced powder

Fiat, Turin, Italy 52

Off-shore drilling and production parts, valve

components, boiler firewall

Stellite, Colmonoy, alloys/carbides

Powder feed Combustion Eng., OH 52

Aerospace components Stellite, Tribaloy Powder feed Rockwell Int., CA 52

Turbine blade shroud interlocks PWA 694, Nimonics Preplaced chip Pratt & Whitney, West

Gas engine turbine blades Hardfacing Powder feed General Electric, SC 54

Laser Melt/Particle Injection

Processing. Laser melt/particle injection produces an in situ, metal-matrix/particulate composite surface layer by

mixing, but not melting, the second phase with the substrate The particulate material is injected with sufficient velocity

as a spray into the melt pool formed by the laser beam If the second phase is hard, such as a carbide, the injected layer can be made to resist wear

Figure 16 depicts the laser injection process An oscillating CW CO2 laser beam produces a shallow melt pool on a substrate, into which are injected powder particles via a nearby nozzle As the sample translates, an injected layer forms

on the surface By varying the amplitude of oscillation, melt widths ranging from 3 to 20 mm (0.12 to 0.8 in.) are possible (Ref 55) Processing conditions for laser melt/particle injection are a power density from 10 to 3000 MW/m2 (6.45 to

1935 kW/in.2) and an interaction time from 0.1 to 1 s Any inert shielding gas is normally used

Fig 16 Laser melt/particle injection process

The particles are propelled by pressurized helium gas The pressure depends on the particle size and the relative densities

of the powder and the molten substrate Larger and heavier particles require lower gas pressures Typically, powder flow rates are from 0.1 to 0.5 cm3/s (0.006 to 0.03 in.3/s) and particle velocities are a few m/s Carrier gas pressures typically range from 50 to 120 kPa (7.3 to 17.4 psi) The carrier gas also serves as a cover gas and keeps the particles relatively cool, preventing them from melting or fusing Particle sizes range from 45 to 150 m (1.8 to 6 mils) Finer particles tend

to either fuse or dissolve in the melt, whereas coarser particles do not flow easily

The injection nozzle, usually made of copper, is inclined at 60° to the horizontal and is positioned from 10 to 20 mm (0.4

to 0.8 in.) away from the sample Its slotted opening is designed to produce a rectangular spray that is of the same size as the melt pool, to ensure that most of the powder is incorporated into the melt

Although minimum interaction between the carbide phase and the melt is desired, some dissolution does occur It was found that the degree of carbide dissolution, as well as other physical characteristics of the injected layer, such as penetration depth, mounding, and carbide volume, could be reasonably controlled Variations in these characteristics, as well as hardness, with laser power and powder feed rate, are determined for Ti-6Al-4V alloy injected with TiC (Fig 17) (Ref 56) The carbide volume can be varied from 15% to a limit of 60% For improved wear resistance, higher carbide volumes are desirable, but the conditions that produce increased carbide volume also increase dissolution An estimation

of the degree of carbide dissolution is given by matrix microhardness As the trends in Fig 17 show, matrix microhardness is higher when the carbide volume is greater

Fig 17 Effect of processing parameters on injected layer characteristics in Ti-6Al-4V injected with TiC (a)

Laser power (b) Powder feed rate Source: Ref 56

Initial work on laser melt/particle injection was done on low-density/high-strength aluminum and titanium alloys, for which few surface-hardening methods exist Some of the alloy/particulate systems that have been investigated are 5052 Al/TiC, Al bronze/TiC, Ti-6Al-4V/TiC, Ti-6Al-4V/WC, 304 stainless/TiC, 4340 tool/TiC, Inconel 625/TiC, and Inconel 625/WC (Ref 57)

Microstructures of Laser Melt/Particle-Injected Layers. The injected layer is a composite of particles surrounded by a metal matrix It is desirable that the only surface modification be the presence of the particulate phase, and that the surrounding metal should remain unchanged and retain most of the substrate properties, such as corrosion resistance and toughness Figure 18 shows the cross-sectional view of a WC injected layer on Inconel 625 alloy substrate The top surface is rough, but is free of deep channels The light grains are the WC particles, which are surrounded by the dark Inconel The carbides are uniformly distributed throughout the injected layer and occupy about 50% of the total volume Because the matrix phase is essentially made of the same material as the substrate, there is chemical continuity across the interface and a strong metallurgical bond

Fig 18 Cross section of Inconel 625 alloy injected with WC

Carbide dissolution products formed during solidification can influence the matrix microstructure The carbide particles sometimes develop a perturbed or scalloped, highly alloyed interface with the metal matrix Eutectic and dendritic carbides are some of the dissolution products that appear within the matrix (Ref 58) Figure 19 shows examples of some

of the dissolution products Besides influencing properties such as microhardness and friction wear, these resolidification products can cause matrix embrittlement and microcracking, although with suitable preheating, cracking can be eliminated (Ref 59)

Fig 19 Carbide morphological changes and matrix resolidification products as a result of carbide phase

dissolution in Inconel 625 alloy (a) Injected with TiC, eutectic carbides (b) Injected with TiC, dendritic carbides (c) Injected with WC, eutectic carbides (d) Injected with WC, dendritic carbides

In order to retain some of the substrate properties in the modified surface, it is important to keep carbide dissolution to a minimum Some of the matrix microstructural changes are inevitable Depending on the base alloy, melting and rapid solidification can result in microstructural refinement, formation of solid solutions, phase transformations, and precipitation These microstructural modifications, along with the dissolution products, can harden the matrix by various mechanisms Carbide dissolution products harden the matrix by a dispersion-hardening mechanism This effect was observed in stainless steel, tool steel, Inconel, and titanium alloys Transformation hardening comes into play when martensite forms in steels and titanium alloys In tool steels, precipitation hardening also plays a role, and in aluminum bronze, microstructural refinement results in modest hardening (Ref 57)

Wear Behavior of Particle-Injected Surface Layers. On a macroscopic scale, hardening of particle-injected alloys is due to plastic flow inhibition by the injected phase The relative increase in macrohardness of the composite surface is from 1.1 to 2.6 times, depending on the alloy (Ref 57) On a microscopic scale, hardening of the metal matrix varies with the type of alloy and has been found to affect friction wear behavior (Ref 60)

Friction wear was evaluated using a balt-on-surface test, which measures the coefficient of kinematic friction, k, between a hard steel ball and a polished sample surface A pressure of 1 N (0.22 lbf) and a slide velocity of 0.0001 m/s (0.02 ft/min) were used This test is sensitive to minute changes in microstructure For untreated Inconel 625, k was about 0.7 Particle-injected Inconel 625 had a k that ranged from 0.15 to 0.20 after the first slide, which increased to a range from 0.3 to 0.4 for WC and 0.4 to 0.45 for TiC after several slides (Fig 20) The wear scar was hardly evident in the optical microscope and there was very little abrasive damage Although k increased with number of slides, it was not due to continuous wear of the sample surface Rather, it was due to formation of wear debris

Fig 20 Coefficient of friction as a function of number of slides in untreated and particle-injected Inconel 625

Source: Ref 60

The wear mechanism was different for the two carbides With WC injection, the carbide and the harder (600 to 650 HV) metal matrix were strong enough to cause the slider to wear, and the wear debris was mostly that from the steel ball With TiC injection, the wear debris was a mixture of material from the steel ball and from the softer (400 to 500 HV) metal matrix (Ref 60)

From dry sand/rubber wheel tests done on aluminum- and titanium-base alloys injected with TiC, it was found that the wear volume decreased rapidly with modest volumes of the carbide phase For aluminum, the reduction is from 0.18 to 0.28 cm3 (0.01 to 0.02 in.3) to 0.02 to 0.025 cm3 (0.0012 to 0.0015 in.3) For titanium, it is from 0.06 to 0.065 cm3 (0.0037

to 0.0040 in.3) to 0.01 cm3 (0.0006 in.3) (Fig 21) (Ref 61) Examination of the worn surface showed that the softer metal matrix eroded, but that the hard carbide particles prevented further erosion of the surface The abrasive action did not uproot the carbides from the matrix, but the normally angular carbides appeared slightly rounded, which helped reduce further erosion

Fig 21 Reduction in abrasive wear rate with increase in vol% carbide in particle-injected aluminum and

titanium alloys Source: Ref 61

Ayers and Bolster (Ref 62) found that abrasive wear with 3 and 30 m (120 and 1200 in.) diamond particles of aluminum and titanium alloy samples was reduced by introducing WC or TiC into their surfaces The wear resistance of aluminum alloys improved by a factor of 30 in samples containing TiC, whereas for Ti-6Al-4V containing TiC, the improvement was less dramatic (factor of 4)

Applications of Laser Melt/Particle Injection Process. Most of the laser melt/particle injection processing work has involved process optimization, microstructure evaluation, and test sample preparation One application involved fabrication of wear-resistant surfaces on Inconel alloy shaft seal test rings, such as the one shown in Fig 22 (Ref 63) The injectant is WC and the circular pass is 127 mm (5 in.) in diameter and 10 mm (0.4 in.) wide The coverage rate was 130

mm2/s (0.2 in.2/s) Before testing, the rough surface was ground flat The seal ring was successfully tested against a graphite mating surface using water as the lubricant

Fig 22 Inconel alloy shaft seal test ring particle-injected with WC Source: Ref 63

References

1 W.M Steen, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals,

C.W Draper and P Mazzoldi, Ed., Martinus Njihoff, The Netherlands, 1986, p 369

2 O Sandven, Heat Treating, Vol 4, 9th ed., Metals Handbook, American Society for Metals, 1981, p 507

3 F.D Seaman, The Industrial Laser, Annual Handbook, D Belforte and M Levitt, Ed., PennWell Books,

1986, p 147

4 J Taylor, Metalwork Prod., Sept 1979, p 138

5 J Meijer, M Seegers, P.H Vroegop, and G.J.W Wes, Proceedings of the International Conference on Applications of Lasers and Electro-optics, C Albright, Ed., IFS Ltd., United Kingdom, 1986, p 229

6 M Yessik and R.P Scherer, Sourcebook on Applications of the Laser in Metalworking, E Metzbower, Ed.,

American Society for Metals, 1981, p 219

7 C Courtney and W.M Steen, Sourcebook on Applications of the Laser in Metalworking, E Metzbower,

Ed., American Society for Metals, 1981, p 195

8 P.A Molian, Iowa State University, private communication, 1992

9 S.L Engle, Sourcebook on Applications of the Laser in Metalworking, E Metzbower, Ed., American

Society for Metals, 1981, p 149

10 P.A Molian and M Baldwin, J Tribol., Vol 108, 1986, p 326

11 D.S Gnanamuthu, Sourcebook on Applications of the Laser in Metalworking, E Metzbower, Ed., American

Society for Metals, 1981, p 324

12 G Lu and H Zhang, Wear, Vol 138, 1990, p 1

13 P.A Molian and M Baldwin, J Tribol., Vol 110, 1988, p 462

14 W.J Tomlinson, R.F O'Connor, and T.A Spedding, Tribol Int., Vol 21 (No 6), 1988, p 302

15 D.N.H Trafford, T Bell, J.H.P.C Megaw, and A.S Bransden, Heat Treating, Proceedings of the International Conference, Metallurgical Society, 1983, p 198

16 J.A Wineman and J.E Miller, Sourcebook on Applications of the Laser in Metalworking, E Metzbower,

Ed., American Society for Metals, 1981, p 209

17 M.J Yessik, Optical Eng., Vol 17 (No 3), 1978, p 202

18 F.D Seaman and D.S Gnanamuthu, Sourcebook on Applications of the Laser in Metalworking, E

Metzbower, Ed., American Society for Metals, 1981, p 179

19 A.V La Rocca, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals,

C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986, p 521

20 General Motors, Sourcebook on Applications of the Laser in Metalworking, E Metzbower, Ed., American

Society for Metals, 1981, p 227

21 W.M Steen, Industrial Laser, Annual Handbook, D Belforte and M Levitt, Ed., PennWell Books, 1986, p

158

22 H.W Bergmann, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals,

C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986, p 351

23 C.H Chen, C.P Ju, and J.M Rigsbee, Mater Sci Technol., Vol 4 (No 2), 1988, p 161

24 M Bamberger, M Boas, and O Akin, Z Mettallkde., Vol 79, 1988, p 806

25 J Folkes, D.R.F West, and W.M Steen, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals, C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986,

p 451

26 C.P Ju, C.H Chen, and J.M Rigsbee, Mater Sci Technol., Vol 4 (No 2), 1988, p 167

27 W.J Tomlinson, J.H.P.C Megaw, and A.S Bransden, Wear, Vol 116, 1987, p 249

28 J Kusinski and G Thomas, Proc SPIE-Int Soc Opt Eng., Vol 668, Society of Photo-optical

Instrumentation Engineers, 1986, p 150

29 H De Beurs and J.Th.M De Hosson, Scr Met., Vol 21, 1987, p 627

30 H De Beurs, G Minholts, and J.Th.M De Hosson, Wear, Vol 132, 1989, p 59

31 M Hsu and P.A Molian, Wear, Vol 127, 1988, p 253

32 G Coquerelle and J.L Fachinetti, Proceedings of the 5th International Congress on Applications of Lasers and Electro-optics, C.M Banas and G.L Whitney, Ed., IFS Ltd., United Kingdom, 1987, p 19

33 J.D Ayers and D.S Gnanamuthu, Metals Handbook, 9th ed., Vol 6, Welding, American Society for Metals,

1983, p 793

34 B.L Mordike, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals,

C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986, p 389

35 C Marsden, D.R.F West, and W.M Steen, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals, C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986,

p 461

36 Z.S Yu and Z.K Quan, Mater Chem Phys., Vol 25, 1990, p 277

37 C Rieker, D.G Morris, and J Steffen, Mater Sci Technol., Vol 5, 1989, p 590

38 Y Tan and J Doong, Wear, Vol 132, 1989, p 9

39 J Xiaoping, H Zhuangqi, G Yunlong, J Ming, and S Changxu, Mater Res Soc Symp Proc., Vol 80,

Materials Research Society, 1987, p 331

40 A.M Huntz, T Puig, L Confignal, F Charpentier, and M Condat, Mater Sci Eng., Vol A121, 1989, p 555

41 J Mazumder and J Singh, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment

of Metals, C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986, p 297

42 A Belmondo and M Castagna, Sourcebook on Applications of the Laser in Metal-working, E Metzbower,

Ed., American Society for Metals, 1981, p 310

43 G Abbas, D.R.F West, and W.M Steen, Proceedings of the Materials Processing Conference, Vol 69,

Laser Institute of America, 1989, p 116

44 P.A Molian and L Hualun, Wear, Vol 130, 1989, p 337

45 M Boas and M Bamberger, Wear, Vol 126, 1988, p 197

46 G.J Bruck, J Met., Vol 39 (No 2), 1987, p 10

47 R.M Macintyre, Proceedings of the NATO Advanced Study Institute on Laser Surface Treatment of Metals,

C.W Draper and P Mazzoldi, Ed., Martinus Nijhoff, The Netherlands, 1986, p 545

48 P Koshy, Proc SPIE-Int Soc Opt Eng., Vol 527, R.R Jacobs, Ed., Society of Photo-optical

Instrumentation Engineers, 1985, p 80

49 S.J Matthews, Lasers in Materials Processing, E Metzbower, Ed., American Society for Metals, 1983, p

138

50 "Quantum Laser Technical Bulletin No 1-86," Quantum Laser Corp., 1986

51 B.F Kuvin, Weld Des Fabr., May 1987, p 35

52 G.M Eboo and A.E Lindemanis, Proc SPIE-Int Soc Opt Eng., Vol 527, R.R Jacobs, Ed., Society of

Photo-optical Instrumentation Engineers, 1985, p 86

53 D.A Belforte, Belforte Associates, private communication, 1987

54 P Mehta, General Electric, private communication, 1987

55 K.P Cooper, R Beigel, and P Slebodnick, Proceedings of the 5th International Congress on Applications

of Lasers and Electro-optics, C.M Banas and G.L Whitney, Ed., IFS Ltd., United Kingdom, 1987, p 169

56 K.P Cooper, International Conference on Surface Modifications and Coatings, R.D Sisson, Jr., Ed.,

American Society for Metals, 1986, p 409

57 K.P Cooper, Proc SPIE-Int Soc Opt Eng., Vol 957, G Sepold, Ed., Society of Photo-optical

Instrumentation Engineers, 1988, p 42

58 K.P Cooper and J.D Ayers, Proceedings of the 1988 Conference on Laser Surface Modification, American

Welding Society, 1989, p 115

59 K.P Cooper and P Slebodnick, J Laser Appl., Vol 1 (No 4), 1989, p 21

60 K.P Cooper, J Vac Sci Technol., Vol A4 (No 6), 1986, p 2857

61 J.D Ayers, Wear, Vol 97, 1984, p 249

62 J.D Ayers and R.N Bolster, Wear, Vol 93, 1984, p 193

63 K.P Cooper and J.D Ayers, Surf Eng., Vol 1 (No 4), 1985, p 263

Carburizing is used today to achieve the same combination of toughness and wear resistance (Ref 1) Power-train gears for helicopters, earth-moving equipment, heavy trucks, and passenger cars are all made of carburized steel Most roller bearings, straight and tapered, are also carburized The primary advances in carburizing technology in the last 100 years are the development of high-quality alloy steels designed for carburizing and improved processing methods to allow better control over the composition, microstructure, and properties of the carburized case

Although quenching in oils at temperatures from 50 to 75 °C (120 to 165 °F) is common, many carburized parts are martempered by quenching in oils or molten salts at temperatures from 175 to 200 °C (345 to 390 °F) Because the quench temperature for martempering is above the martensite start temperature of the high-carbon case, the case transforms to martensite during subsequent air cooling

Other forms of carburizing, described in more detail in Volume 4 of the ASM Handbook (1991), include:

• Pack carburizing, where parts are packed in a blend of coke and charcoal with "activators," and then heated in a closed container This is an old, labor-intensive process, but it is still practiced in tool rooms because facility requirements are minimal

• Vacuum and plasma carburizing, both of which utilize a vacuum chamber with a partial pressure of hydrocarbon gas as the source of carbon The gas is ionized in the plasma carburizing process The prime advantage to these processes is the absence of oxygen in the furnace atmosphere

• Salt bath carburizing, in which the baths contain cyanides, cyanates, or carbon-carbonate blends These are particularly useful for producing thin carburized cases because the carburizing time can be precisely controlled

Characteristics of Carburized Surfaces

Diffusion of Carbon. Some of the features of carburized cases are due to the fact that they are created by the diffusion

of carbon First, there is a gradual transition in carbon content, as well as a transition in microstructure and mechanical properties, between case and core As a rule, the deeper the case, the less steep is the slope of the carbon gradient (Fig 1) The absence of any sharp transition in properties assures excellent adherence of the case

Fig 1 Computed carbon concentration gradients resulting from gas carburization of SAE 8620 steel for 4, 8,

and 16 h at 927 °C (1700 °F) Carbon potential of furnace atmosphere assumed to be 1 wt% C during process

Second, carburized cases are most frequently produced in thicknesses that range from 0.5 to 1.5 mm (20 to 60 mils) At

927 °C (1700 °F), which is a typical processing temperature, this thickness range can be produced in processing times from about 2 to 15 h Cases as thin as 0.1 mm (4 mils), which require less than 10 min at 927 °C (1700 °F), are sometimes produced on small parts by salt bath carburizing Cases as deep as 3 mm (120 mils), which require more than two days of carburizing at 927 °C (1700 °F), are occasionally produced on large parts The ease of adjusting the case depth to resist the anticipated contact loading is one of the great advantages of carburizing

The microstructure of martensitic steels varies with carbon content in two important ways First, unalloyed steels with

carbon contents of less than about 0.5/0.6 wt% form dislocated lath martensite upon quenching Steels with carbon contents above 1.0% form twinned plate martensite Mixtures of lath and plate martensite are found with intermediate carbon levels (Ref 3, 4) Thus, a carburized case will have a mixture of martensite morphologies, with plate martensite dominating in the high-carbon outer layers, and lath martensite dominating toward the core Alloying elements, which strengthen austenite, promote the formation of plate martensite at lower carbon levels

Over the same range in carbon as the transition from lath to plate martensite, the amount of austenite retained in the quenched structure increases (Fig 2) However, the amount of austenite retained is a function of other variables, as well

as-It tends to increase as the quench rate decreases, and as the content of alloying elements increases Thus, at a given carbon content, the austenite content of as-quenched parts can be higher than the values shown in Fig 2

Fig 2 Percent retained austenite as function of carbon content Vertical lines show carbon ranges in which lath

and plate martensites are found in Fe-C alloys Source: Ref 5

For example, over 30% retained austenite can be produced in 52100 steel if all the carbon is dissolved during austenitization (Ref 6) In another example, a retained austenite content of 50% has been reported in SAE 8620 steel, with

a case carburized to 0.9% carbon (Ref 7)

Although the hardness of martensite increases uniformly with carbon content, the hardness of the martensite-austenite composite obtained on quenching is usually reported to exhibit a plateau above approximately 0.6 wt% carbon (Fig 3) (Ref 8, 9) In fact, most alloys, when carburized, oil quenched, and tempered at temperatures from 150 to 200 °C (300 to

390 °F), exhibit a maximum in hardness for case carbon contents in the range from 0.6 to 0.9% (Ref 10)

Fig 3 Rockwell "C" hardness as function of carbon content Upper curve, after Bain and Paxton (Ref 7), is their

estimate of hardness of martensite, with shaded region above 0.9% C representing the uncertainty that is due

to retained austenite Lower band, after Crafts and Lamont (Ref 8), is their estimate of "the maximum hardness usually attained in commercial quenching."

The range in case hardness that is possible at a given carbon level for a variety of carbon steels is much broader than suggested by Fig 3 (A highly alloyed carburizing steel, such as SAE 9310, may have a case hardness of less than 55 HRC after oil quenching.) When resistance to indentation is the primary concern, the case carbon content is usually specified to achieve the maximum in hardness

Martensitic structures subjected to alternating shear stresses from repeated contact loading for extended periods of time are found to decompose because of carbon migration aided by dislocation motion (Ref 11) However, microstructural decay does not appear to be a necessary precursor to failure by the usual contact fatigue mechanisms

Composite structures that contain primary carbides can also be produced by carburizing Three common morphologies are discussed below

Coarse primary carbides (from 1 to 10 m, or 40 to 400 in.) (Fig 4) are produced by carburizing in an atmosphere with a carbon potential that is high enough to exceed the solubility limit for carbon in austenite Coarse carbides are often found at corners and edges of parts made of alloys that are rich in strong carbide formers, such as chromium Plate martensite and high levels of retained austenite can be found with coarse carbides in parts quenched directly from the carburizing temperature Processing is sometimes designed to produce large primary carbides as a means of enhancing wear resistance (Ref 12) More often, large carbides are avoided, because they deplete the matrix in alloying elements such as Cr, thereby reducing hardenability

Fig 4 Coarse primary carbides produced by carburizing SAE 4130 steel at 950 °C (1740 °F), and then

quenching Matrix microstructure is plate martensite and retained austenite Picral etch 600×

Carbide networks (Fig 5) form in austenite grain boundaries when parts are carburized at an elevated temperature, and then slowly cooled The solubility limit in austenite is exceeded as cooling occurs, and carbon is rejected to austenite grain boundaries Because this structure tends to embrittle the case, it is usually avoided

Fig 5 Carbide networks in prior austenite grain boundaries Produced by carburizing 4130 steel at 950 °C

(1740 °F), furnace cooling to 800 °C (1470 °F), and then quenching Picral etch 600×

Fine primary carbides (from 0.1 to 0.5 m, or 4 to 20 in diam) (Fig 6) result when a part is carburized at a high temperature, such as 950 °C (1740 °F), cooled to form pearlite or bainite, and then reheated to a lower temperature, such

as 830 °C (1525 °F), for a brief time and quenched Because the carbon solubility at 950 °C (1740 °F) is on the order of

1.5 times the solubility at 830 °C (1525 °F), substantial quantities of fine carbides can be produced The carbides will not coarsen significantly if the time at the lower austenitizing temperature is on the order of 30 min or less Because the retained austenite content is a function of the carbon dissolved in austenite, the hardness will be near the maximum for the alloy Finely dispersed primary carbides can incrementally increase the hardness Their main contribution is in restricting austenite grain growth, thereby assuring fine martensite plates and finely dispersed retained austenite Many gears and bearings are heat treated in this manner

Fig 6 Fine primary carbides in lath martensite produced by carburizing at 950 °C (1740 °F), air cooling to

room temperature, then reheating to 820 °C (1510 °F) for 20 min and quenching Picral etch 600×

Residual Stress. Carburized parts usually contain compressive residual stresses in the high-carbon case The stresses arise because of the sequence of transformations upon quenching (Ref 13) The higher the carbon content, the lower the temperature at which martensite begins to form The lower the transformation temperature, the greater the volume increase upon transformation Therefore, upon cooling, the core transforms first and the volume change is accommodated

by plastic deformation of the austenitic case When the case transforms, the martensitic core accommodates the volume increase in the case by elastic, rather than plastic, deformation Consequently, compressive residual stresses in the case tend to be balanced by tensile residual stresses in the core

Actual stress distributions can be more complex than the scenario above would suggest, because:

• Temperature gradients within the part cause transformation to begin below the case/core interface, and then progress inward and outward simultaneously

• Austenite that is retained upon quenching does not contribute to the case expansion

From a consideration of force balances, it follows that when the case is thin relative to the core, compressive residual stresses in the case will be high If the case is thick relative to the core, compressive stresses in the case will be low and tensile residual stresses in the interior will be high Compressive surface residual stresses are beneficial in applications involving torsional and bending loads, because the tensile stresses caused by these loads are a maximum at the surface Compressive residual stresses are also beneficial in resisting rolling contact fatigue because they offset the tensile stresses responsible for initiating and propagating subsurface fatigue cracks (Ref 14, 15) However, in some parts, high compressive surface residual stresses can cause problems, such as case/core interface cracks at tooth tips in gears

Residual stresses can be significantly reduced by tempering at temperatures up to 200 °C (390 °F), without greatly affecting either the hardness or retained austenite content (Ref 16) Changes in residual stress commonly occur during service (Ref 17)

Transformation of Retained Austenite. Under the action of elastic stresses or plastic deformation, retained austenite can transform to martensite (Ref 18, 19, 20) Transformation often occurs in service in parts that experience high contact loads The surface hardness increases, and the resulting volume expansion increases the compressive residual stress at the surface Carburized roller bearings typically contain from 30 to 40% retained austenite in the case Lowering the retained austenite content usually has an adverse effect on rolling contact fatigue life (Ref 7, 21)

On the other hand, manufacturers of instrument bearings, who strive to maximize dimensional stability, make an effort to eliminate retained austenite (Ref 22) Similarly, in applications requiring the best resistance to indentation, the retained austenite content is kept low to maximize hardness Ball bearings, for example, experience "point loading," whereas gears and roller bearings experience "line loading." Therefore, manufacturers of ball bearings, attaching more importance to indentation resistance, choose a through-hardened steel, SAE 52100, which is readily processed to yield from 5 to 10% retained austenite and hardness values greater than 62 HRC

Surface Oxidation. Most carburizing is done in environments that contain some oxygen An atmosphere capable of carburizing steel will also reduce iron oxides, but it often contains sufficient oxygen to oxidize alloying elements, such as

Si, Mn, Cr, and V During carburizing cycles of several hours duration, oxides of these elements will form at the surface and/or in grain boundaries intersecting the surface (Fig 7)

Fig 7 Oxides formed at grain boundaries near the surface Same sample as Fig 5 Unetched 600×

Depletion of these elements from the matrix can lower the hardenability enough to allow nonmartensitic products to form upon quenching, causing low hardness in a surface layer that is one to two grain diameters in thickness (Ref 23) Nonmartensitic products at the surfaces of gear teeth are known to reduce the tooth bending fatigue strength and would be expected to have an adverse effect on tribological properties, as well If the matrix microstructure is unaffected, the effect

of oxides on tribological properties is probably similar to that produced by oxide inclusions in the steel Manufacturers of fine-pitch gears often do not attempt finishing operations to remove grain boundary oxides after heat treatment, arguing that the risk of surface damage caused by overheating during grinding outweighs any possible gain

The thin layer of iron oxide that forms on carburized parts tempered in air at temperatures from 150 to 200 °C (300 to 390

°F) is usually benign Thicker oxide layers, formed at higher temperatures, are sometimes used on steels as "break-in" coatings

Carburizing Steels

Three factors are usually considered when selecting carburizing steels: cleanliness, hardenability, and fabricability, each

of which is discussed below

Cleanliness. For the best resistance to rolling contact fatigue (spalling), the content of aluminate, silicate, and globular oxide inclusions (Types B, C, and D, respectively, in the Jernkontoret system, ASTM E 45) must be as low as possible Manganese sulfide inclusions (Type A) are generally not regarded as detrimental to rolling contact fatigue life (Ref 24)

The inclusion standards specified in ASTM A 534, "Carburizing Steels for Anti-Friction Bearings," have been steadily tightened since the late 1960s, reflecting improvements in steelmaking Some individual steel suppliers claim to be able to furnish premium-quality carburizing steels with oxygen contents below 15 ppm, titanium contents below 30 ppm, and inclusion ratings considerably better than ASTM A 534 Calcium treatment of bearing-quality steels to modify aluminates, thereby improving machinability, is usually avoided, because of the possibility of forming large inclusions

When sliding is combined with rolling contact, near-surface inclusions (including oxides formed in grain boundaries during heat treatment) promote pitting (Ref 25) The role of inclusions in most forms of sliding wear is not as well defined as their role in rolling contact fatigue, probably because the conditions that are possible at a sliding interface are more diverse and more difficult to characterize, than at a rolling contact interface Some insight into possible effects of

inclusions on sliding wear comes from the machining literature (Ref 26) It is known that some inclusions in a steel workpiece can promote tool wear during machining, whereas others can reduce wear

Hardenability. The alloy content of carburizing steels is usually selected on the basis of hardenability If the application involves high contact loads (roller bearings, for example), the uncarburized core must be martensitic to prevent the subcase from yielding An alloy that allows the part to attain full hardness (through-harden) in whatever quenchant is employed will be selected The selection of an alloy with sufficient core hardenability almost always assures sufficient case hardenability When contact loads are well within the capability of the case to support them, it is often not necessary nor desirable for the core microstructure to be martensitic Shape distortion during quenching, for example, is usually reduced if the core transforms at a relatively high temperature to a nonmartensitic structure For such parts, the alloying need only be sufficient to ensure case hardenability

Jominy hardenability bands for alloy carburizing steels are described in the article "Hardenability Curves," in Properties and Selection: Irons, Steels, and High-Performance Alloys, Volume 1 of ASM Handbook Narrower bands for restricted

hardenability steels (Ref 27) have been devised in response to user demands for more precisely defined hardenability Modern steelmaking practice provides much better control over hardenability, through composition control, than is reflected in the older standards

The Jominy bands for carburizing steels provide a means of comparing the core hardenability of steels To ensure adequate case hardenability, one needs a Jominy band for the same base composition, with the carbon content raised to about 0.4 wt% (This presumes that one wishes all of the case with a carbon content of 0.4% or greater to be fully hardened.) For common alloys, such as 1524, 4023, 5120, 8620, and 8720, a "low-side estimate" of the case hardenability

is given by the Jominy bands for 1541, 4042, 5140, 8640, and 8740, respectively For other carburizing steels, case hardenability must be estimated from the composition or determined experimentally (Ref 28)

Fabricability. Manufacturing considerations also influence alloy selection SAE 8620 steel readily transforms after hot rolling and annealing to a ferrite-pearlite microstructure (150-180 Brinell hardness) that has good machinability If fabrication is by cold forming, rather than machining, then SAE 4118 might be selected, because it can be softened to a lower hardness with better ductility than 8620

In manufacturing small automotive gears, steel rod is often annealed to facilitate blank-making by cold forming Blanks are then normalized to a higher hardness to facilitate subsequent gear hobbing

Unfortunately, with increasing hardenability comes increasing difficulty in softening This is particularly true in alloys such as 4815 and 9310, in which the transformation to pearlite is very sluggish, requiring about one day at 600 °C (1110

°F) for completion Because microstructures are usually less than optimum, machining of these alloys is more difficult

Other Considerations. Several secondary hardening carburizing alloys have been developed for applications that

require resistance to elevated temperatures, such as helicopter gearing and rock drill bits (Ref 29, 30, 31, 32) These alloys make use of the precipitation of copper and/or M2C and MC carbides to provide resistance to softening for temperatures

up to 550 °C (1020 °F) Because they contain substantial amounts of Mo and V, these alloys resemble low-carbon versions of tool steels Some of the alloys are difficult to carburize because of high Si and Cr contents; preoxidation prior

to carburizing is necessary to permit carbon penetration (Ref 33) Secondary hardening alloys could also be useful in ambient-temperature applications in which lubrication is marginal, because the heat generated by intermittent metal-to-metal contact would not readily soften the underlying metal (Ref 34) The article "Microstructures and Properties of

Carburized Steels" in Heat Treating, Volume 4 of ASM Handbook (1991) expands upon the discussion in this section

Process/Materials Selection for Wear Resistance

The independent variables available for controlling the microstructure/properties of carburized cases are those that define the carburizing alloy (composition, cleanliness), and those that define the carburizing process (time/temperature/carbon-potential carburizing history, time/temperature quenching history, time/temperature tempering history) These tools provide a considerable degree of control over these microstructural features:

Martensite

• (a) Carbon content of source austenite

• (b) Plate size (austenite grain size)

• Residual stress distribution

which determine the tribological properties of the case The combination of properties that is best for each application must then be decided

The necessary case depth and case hardness can be estimated from a Hertzian stress calculation, but other microstructural objectives can be specified only qualitatively For many applications, the following "rules of thumb" apply:

• Sufficient case depth and case hardness must be provided to prevent indentation or case crushing under the anticipated contact loads For gears and bearings loaded in "line contact," a minimum case hardness

of 58 HRC frequently is specified When high contact loads are accompanied by sliding, the surface hardness (to a depth of about 50 m, or 2 mils) may have to be raised to prevent shearing of surface layers

near-• The retained austenite content should be as high as possible, consistent with the requirements of the rule above The retained austenite content should be controlled by adjusting the case carbon content, not by subzero quenching after carburizing or by tempering at temperatures above 200 °C (390 °F)

• The tempering temperature chosen should be as low as possible, but above the surface temperatures anticipated in finishing operations and in service

• The content of nonmetallic inclusions should be no higher than that needed for economical machining

• Coarse primary carbides can be helpful in resisting abrasive wear Fine primary carbides can permit more retained austenite at the same hardness level Experiments should be conducted to verify any benefits presumed to be associated with primary carbides

For certain types of gearing, the selection of materials and processes is dominated by concerns with resistance to tooth bending fatigue, rather than the tribological properties of the surfaces in contact The best resistance to high-cycle bending fatigue appears to be associated with lower case carbon contents, higher hardness, and less retained austenite than is optimum for contact fatigue (Ref 10) Fortunately, in the roots of gear teeth (where bending strength is important), the case carbon content tends to be lower than in the tooth area (where tribological properties are important), because of curvature effects

References

1 C.M Kim, Case Hardened Steels: Microstructural and Residual Stress Effects, D.E Diesburg, Ed.,

TMS-AIME, 1984, p 59-87

2 C.A Stickels, ASM Handbook, Vol 4, ASM International, 1991, p 312-324

3 G Krauss, Steels: Heat Treatment and Processing Principles, ASM International, 1990

4 R.W.K Honeycombe, Steels: Microstructure and Properties, American Society for Metals, 1981

5 K.-E Thelning, Steel and its Heat Treatment, 2nd ed., Butterworths, London, 1984

6 C.A Stickels, Metall Trans., Vol 5, 1974, p 865-874

7 J.A Erickson, Met Prog., Vol 92, 1967, p 69-73

8 E.C Bain and H.W Paxton, Alloying Elements in Steel, American Society for Metals, 1966

9 W Crafts and J.L Lammont, Hardenability and Steel Selection, Pitman & Sons, London, 1949

10 C Razim, Alloys for the Eighties, R.Q Barr, Ed., Climax Molybdenum Co., 1980, p 9-23

11 H Swahn, P.C Becker, and O Vingsbo, Metall Trans A, Vol 7A, 1976, p 1099-1110

12 R.F Kern, Heat Treat., Oct 1986, p 36-38

13 J.A Burnett, Residual Stress for Designers and Metallurgists, American Society for Metals, 1981, p 51-69

14 G.T Hahn, V Bhargava, C.A Rubin, and X.G Leng, Carburizing: Processing and Performance, ASM

International, 1989, p 101-113

15 D Brooksbank and K.W Andrews, J Iron Steel Inst., Vol 210, 1972, p 246-255

16 R.L Brown, H.J Rack, and M Cohen, Mater Sci Eng., Vol 21, 1975, p 25-34

17 H Muro, N Tsushima, and K Nunome, Wear, Vol 25, 1973, p 345-356

18 F.T Krotine, M.F McGuire, L.J Ebert, and A.R Troiano, ASM Trans Quart., Vol 62, 1969, p 829-838

19 R.H Richman and R.W Landgraf, Metall Trans A, Vol 6A, 1975, p 955-964

20 C.A Stickels, Metall Trans A, Vol 8A, 1977, p 63-70

21 H Muro, Y Sadaoka, S Ito, and N Tsushima, Proceedings of the 12th Japanese Congress of Materials Research, Mar 1969, p 74-77

22 T.J Hughel, ASM Trans Quart., Vol 62, 1969, p 18-23

23 R Chatterjee-Fischer, Metall Trans A, Vol 9A, 1978, p 1553-1560

24 P Tardy, Archiv Eisenhuttenw., Vol 43, 1972, p 583-587

25 T.M Clarke, G.R Miller, L.M Keer, and H.S Cheng, ASLE Trans., Vol 28, 1985, p 111-116

26 F.W Boulger, Properties and Selection: Irons, Steels, and High-Performance Alloys, Vol 1, 10th ed., Metals Handbook, ASM International, 1990, p 591-602

27 Standard J1868, SAE Handbook, Vol 1, Society of Automotive Engineers, 1990

28 Standard J406, SAE Handbook, Vol 1, Society of Automotive Engineers, 1990

29 C.F Jatczak, Met Prog., Vol 113, 1978, p 70-78

30 W.E Burd, Met Prog., Vol 127, 1985, p 33-35

31 R.A Cutler and W.C Leslie, J Test Eval., Vol 11, 1983, p 3-15

32 J.D Saulnier, Carburizing: Processing and Performance, ASM International, 1989, p 211-219

33 M.L Schmidt, J Heat Treat., Vol 8, 1990, p 5-19

34 C.A Stickels and C.M Mack, J Heat Treat., Vol 4, 1986, p 223-236

Nitriding and Nitrocarburizing

F.T Hoffmann and P Mayr, Institute of Material Science (Germany)

Introduction

THE NITRIDING AND NITROCARBURIZING PROCESSES are, aside from carburizing, the most important thermochemical processes in heat treating industrial parts for the production of case-hardened surface layers Layer of high hardness, which provide high resistance to corrosion and wear in addition to high fatigue strength, are produced by the diffusion of atomic nitrogen into the surfaces Other advantages include very low distortion, because of the low temperatures involved and the absence of phase transformations, and high tempering resistance associated with the high hardness property at service temperatures below the nitriding temperature

The original nitriding process used gaseous ammonia as the nitrogen source Within the past 60 years, numerous nitriding and nitrocarburizing processes have been developed, some of them for specialized applications Today, parts can be nitrided by using powders, salt baths, gaseous mixtures, amd plasma-assisted processes as nitrogen sources Detailed

information on the many forms of nitriding and nitrocarburizing are given in Heat Treating, Volume 4 of the ASM Handbook (1991)

Characteristics of Nitrided Surfaces

The structure of a nitrided layer comprises:

• A compound layer on the surface, consisting primarily of nitrides and carbonitrides of types Fe4N and Fe2-3 (N,C) In addition, special nitrides are formed in steel that contain nitride-forming elements, such

as aluminum, chromium, molybdenum, vanadium, and titanium If specified by the user, the compound layer can be suppressed by a special nitriding practice

• A diffusion layer below the compound layer, composed of nitrogen in interstitial solution in the ferritic matrix in combination with nitride dispersions

The composition and microstructure of the nitride layer can be greatly affected by process selection and specific process parameters The composition of the compound layer can be affected by the nitriding atmosphere Higher temperatures or a longer nitriding time produce increased growth of the compound layer The structure of the diffusion layer (that is, type, size, and distribution of the nitrides) can be affected by:

• The nitriding temperature

• The cooling rate from nitriding temperature to room temperature

• Low-temperature annealing after cooling to room temperature

• The type of steel being nitrided

High nitriding temperatures promote the formation of large, coarse distributed nitrides; low temperatures, promote the formation of finely dispersed nitrides Slow cooling rates from nitriding temperature to room temperature, annealing treatments, or heating by stressing the parts favors dissolution, growth, and change in the nitride structure The nitriding depth can be increased either by increasing the process temperature or by increasing the nitriding time Increasing the content of nitride-forming elements in the material increases the hardness of the surface layer ad decreases the nitriding depth Both the hardness and the depth parameters depend on the extent of nonferrous nitride formation

In addition, the atmosphere composition and nitriding temperature (that is, the nitriding potential and the carbon potential

of the atmosphere) affect the phase composition and the distribution of nitrogen and carbon in the compound layer Investigations (Ref 1) have shown that if the carbon potential of a nitriding atmosphere is low, every nitriding process starts with decarburization of the matrix Increasing carbon potential (when nitrocarburizing) and increasing nitriding potential (when nitriding) decreases decarburizing This effect is present even when Armco iron is used This leads to zones below the compound layer with a lower carbon content than in the matrix The diffusion coefficient of carbon in the iron matrix is lowe than the diffusion coefficient of carbon in the nitride layer, but the solubility of carbon in the nitride layer is higher than the solubility of carbon in the iron matrix Thus, after the formation of the compound layer, carbon diffusing to the surface is enriched in the compound layer at the interface with the iron matrix Figure 1 is a plot of the compound layer composition as a function of nitriding time The carbon content in this enriched zone depends on the type

of forming nitride (that is, the nitriding potential of the atmosphere) and on the carbon content of the steel itself

Fig 1 Plot of compound layer concentration versus distance from the surface for Ck 45N nonresulfurized

carbon steel as a function of gas nitriding duration (a) Nitrogen distribution (b) Carbon distribution Nitrided at

540 °C (1005 °F) Nitriding number (KN), 0.5

Wear Resistance of Nitrided and Nitrocarburized Materials

There are hundreds of articles in the literature that describe wear problems and their solution by different nitriding or nitrocarburizing processes However, few of them describe the investigated nitride layers sufficiently Most fail to specify the nitriding process, nitriding parameters, phase composition, compound layer thickness, and chemical composition and microstructure of the nitrided material Therefore, the wear behavior of nitrided parts is somewhat difficult to ascertain from the literature In this article, trends were deduced from recent, albeit contradictory, results (Ref 1, 2, 3, 4, 5, 6, 7, 8,

9, 10, 11, 12, 13, 14, 15, 16, 17, 18, 19, 20, 21, 22, 23, 24, 25, 26, 27, 28, 29, 30, 31, 32, 33) in order to make generalized statements on wear behavior

The wear resistance of a nitride layer changes with distance from the surface (Fig 2) Wear resistance is reduced in the porous zone of the compound layer because of the lower fatigue strength, reduced density, and the notch effect of the pores The wear resistance of the poreless zone of the compound layer is significantly higher than that of the diffusion zone and the core material If the compound layer has a homogeneous structure, wear resistance is constant throughout In the diffusion layer, wear resistance decreases to the value of the core material with increasing distance from the surface;

this is because of the reduction of the density of the nitride precipitates and the supersaturation of the matrix with nitrogen

Fig 2 Plot of adhesive and abrasive wear resistance versus distance from the surface for a nitride layer

The wear behavior of a nitride layer is often assumed to be an integral reaction of the nitriding layer to wear loading This view does not explain wear mechanisms and their interaction during loading Therefore, wear resistance and mechanisms must be discussed in conjunction with the structure of the nitriding layer Table 1 gives the wear resistance of the compound layer and the diffusion layer in terms of the four major wear mechanisms discussed below

Table 1 Effect of nitride layer on the wear resistance and the interaction parameter of metals as a function

of wear mechanism

Interaction parameters Relative wear resistance(b)

Wear

mechanism Compound layer (a)

Diffusion layer (a) Compound layer Diffusion layer Adhesion Structure Formation of a protection layer (pl) XXX X

vpl,form > vpl,sol vpl,form > vpl,sol X X

Tribo-oxidation

vpl,form < vpl,sol vpl,form < vpl,sol (c) (c)

Abrasion Structure Solid solution and precipitation hardening XX X

(a) vpl,form, formation rate of friction-induced protective layer; vpl,sol , solution rate of friction-induced

protective layer

(b) Increase relative to nonnitrided surface

(c) Decrease in wear resistance

Adhesion. In adhesive wear, the wear resistance of the compound layer is very high compared with nonnitrided parts This is because the inclination to microwelding decreases due to the change in the electron configuration in the outer region of the part In the diffusion layer, wear resistance can also increase because of the formation of protective nitrogen-containing layers

Tribo-oxidation. The wear resistance of the compound layer against tribo-oxidation depends on the formation of

protective layers in the contact zone If the formation rate, vpl,form, of the friction-induced protective layer is greater than the concurrent solution rate, vpl,sol, wear resistance increases If vpl,sol, > vpl,form, wear resistance decreases The wear

resistance of the diffusion layer seems to be behave similarly

Abrasion. The compound layer is very resistant to abrasive wear, because the structure of the compound layer allows only very low plastic deformation Compared with nonnitrided steel, the wear resistance of the diffusion layer is higher a result of the higher fatigue strength obtained by solid-solution strengthening and precipitation hardening

Surface Fatigue. It is assumed that the wear resistance of nitrided parts to surface fatigue is higher than that of nonnitrided parts, as the changed lattice structure of the compound layer prevents plastic deformation The diffusion zone has very high resistance to surface fatigue wear; here plastic deformations are very small because of solid-solution strengthening and precipitation hardening

Influence of Variables on Wear Resistance of Nitrided Parts

Compound Layer. Variables that influence the wear resistance of the compound layer of nitrided parts are illustrated in Fig 3, 4, 5, 6, 7, 8, and 9 A porous zone will raise the initial wear in the case of adhesive and abrasive wear (Fig 3 and 4) If tribo-oxidation is the main wear mechanism, a large increase in wear resistance can occur (Fig 5) This may be due

to absorption of the lubrication medium and subsequent formation of a lubrication layer In the case of surface fatigue (Fig 6), compound layer thickness has no influence on wear resistance if the porous zone of the compound layer is small and the maximum of the true stress ( v) lies in the deeper regions of the compound layer or below it If the thickness of the porous zone increases the distance from the surface of the maximum stress, wear resistance to surface fatigue will rapidly decrease

Fig 3 Effect of porous zone thickness on adhesive wear resistance of the compound layer of a nitrided part

Fig 4 Effect of porous zone thickness on abrasive wear resistance of the compound layer of a nitrided part

Fig 5 Effect of lubrication on resistance of the compound layer of a nitrided part to wear caused by

tribo-oxidation

Fig 6 Effect of porous zone thickness and maximum stress on resistance of the compound layer of a nitrided

part to wear caused by surface fatigue

The structure and composition of the compound layer also influence its wear resistance Investigations have shown that the adhesive wear resistance of the compound layer is strongly affected by the volume of -nitrides In most cases, resistance increases with increasing -nitride content (Fig 7) Similar behavior exists when nitrided parts are abrasively stressed (Fig 8a) However, there is a difference between layers consisting only of nitrides and those containing carbonitrides At constant volume of -nitride, carbonitride layers show significantly higher wear resistance than nitride layers (Fig 8b) For surface fatigue, a higher content of -nitride in the compound layer will result increased wear resistance if the surface pressure is held constant (Fig 9)

Fig 7 Effect of -nitride content on adhesive wear resistance of the compound layer of a nitrided part

Fig 8 Effect of -nitride content on abrasive wear resistance of the compound layer of a nitrided part (a)

Resistance increases with increasing content (b) At constant volume of , carbonitride layers show significantly higher resistance than nitride layers

Fig 9 Effect of -nitride content on resistance of the compound layer of a nitrided part to wear caused by

surface fatigue at constant pressure

Diffusion Layer. Variables that affect the wear resistance of the diffusion layer of nitrided parts are illustrated in Fig

10, 11, 12, 13, 14, 15, 16, and 17 Hardening of the alloy by precipitation methods such as solid-solution strengthening will raise the adhesive, abrasive, and surface fatigue wear resistance (Fig 10 11 12) Greater nitriding depths also increase wear resistance (Fig 13 14 15) Increases in the nitriding depth allow initial high wear resistance to abrasion to be upheld for a longer time, thus increasing component life (Fig 14) Behavior is similar for surface fatigue stressing: increases in nitriding depth allow higher surface pressures and constant wear resistance (Fig 15)

Fig 10 Effect of material strength on adhesive wear resistance of the diffusion layer of a nitrided part

Fig 11 Effect of material strength on abrasive wear resistance of the diffusion layer of a nitrided part

Fig 12 Effect of material strength on resistance of the diffusion layer of a nitrided part to wear caused by

surface fatigue

Fig 13 Effect of nitriding depth on adhesive wear resistance of the diffusion layer of a nitrided part

Fig 14 Effect of nitriding depth on abrasive wear resistance of the diffusion layer of a nitrided part Increases

in nitriding depth increase component life Nitriding depth, d: d1 < d2 < d3 < d4

Fig 15 Effect of nitriding depth on resistance of the diffusion layer of a nitrided part to wear caused by surface

fatigue Increases in nitriding depth allow the part to withstand higher surface pressures without sacrificing

wear resistance p1 < p2 < p3

Nitride type and distribution in the diffusion layer can be changed by a postnitriding aging, strongly influencing wear resistance For example, when nitrided parts are age hardened at constant temperature for different times, a characteristic hardness profile with a distinct maximum will result Adhesive or abrasive wear stressing of such parts also results in a maximum wear resistance; however, the maximum shifts to higher aging times Thus, both types of wear are sensitive to the type and distribution of the nitride precipitation in the diffusion layer: a matrix wit finely dispersed Fe16N2 nitrides has

a lower wear resistance than one with a coarse distribution of Fe4N nitrides (Fig 16) The same is true for parts subjected

to wear caused by surface fatigue (Fig 17)

Fig 16 Hardness profile showing effect of annealing time on abrasive and adhesive wear resistance of the

diffusion layer of a nitrided part See text for discussion of the effect of type and distribution of nitrides

Fig 17 Hardness profile showing effect of annealing time on resistance of the diffusion layer of a nitrided part

to wear caused by surface fatigue See text for discussion of the effect of type and distribution of nitrides

Influence of the Nitriding/Nitrocarburizing Process on Wear Behavior

Investigations of the influence of the nitriding process on wear behavior can be divided into two groups: process oriented and materials science oriented The objective of process-oriented investigations is often to show the clear advantages of a certain process Materials-science-oriented investigations are usually more varied Evaluation of published results of experimental work indicates that nitriding process variables do not exert much influence if nitriding layers are compared and are optimized to the actual wear stress

Process variables do have an influence, however, if changes in geometry, which depend on the process, are considered Typical process-dependent changes in geometry are shown in Table 2 Parts with elevations at their corners will have

shorter lifetimes because of the higher surface pressures at the corners than parts with constant surface pressures over the entire contact zone

Table 2 Effect of nitriding process on surface topography and surface finish

Optimization of Wear by Process Technology

The influence of nitriding process parameters on wear resistance is not very high, but, with limitations, specific wear properties can be fitted to loading conditions by process optimizing Table 3 summarizes possible enhancements of wear resistance for various wear mechanisms

Table 3 General guidelines for improving the wear resistance of nitrided and carbonitrided steels as a function of wear mechanism and surface layer type

• Convert from nitriding to carbonitriding to add additional carbon donors

• Minimize porous zone region by polishing surface

Modify nitriding atmosphere:

• Salt bath: change content of base from CN- to CNO-

• Gaseous: change nitriding number, flow rate, and carbon donors

Abrasion, adhesion

Formation of reactive layers Add oxidation process after nitriding process is

completed Tribo-

Increase strength of material Select alternate materials, add heat treatment prior to

nitriding (use hardening and tempering operations instead

oxidation

Tribo-Modification of tribosystem Optimization of hardness profile Increase temperature and duration of nitriding process;

select alternate materials; add heat treatment prior to nitriding; increase cooling rate after nitriding; add annealing treatment after nitriding

• Changing '-nitride to -nitride by higher nitriding potentials (nitriding numbers)

• Changing nitrides to carbonitrides by adding carbon donors such as Endogas and CO2 or by using a material with a higher carbon content

• Minimizing the porous zone by polishing

Porosity can also be influenced by the atmosphere (for example, lower nitriding potential or lower content of carbon donors), but this may worsen the properties of the compound layer However, a small porous zone in the compound layer may be advantageous when the shape of the surface can be optimized by running-in (initial) wear

In gas nitriding, the structure of the compound layer can be changed by a definite nitriding atmosphere that is, a definite nitriding number [ratio p(NH3)/p(H2)1.5] or by additional carbon-containing gases such as CO2 or Endogas In plasma nitriding, additional parameters such as pressure and glow discharge conditions are beneath the atmosphere In bath nitriding, the concentration of the cyanate and cyanide content of the bath can be changed

A reaction layer can be formed by an oxidation process after nitriding Usually, these processes are performed in salt baths or in gaseous atmospheres

The adhesive and abrasive wear resistance of the diffusion layer can be enhanced by strengthening the diffusion layer The most effective means of this are by proper material selection (for example, high content of nitride-forming elements) and by heat treatment before nitriding (for example, hardening and tempering instead of normalizing)

Tribo-oxidation. If tribo-oxidation is the primary wear mechanism, only limited improvements can be made by changing the nitriding conditions In this case, it is necessary, if possible, to change the conditions of the tribosystem

Surface Fatigue. If controlling surface fatigue is the objective, measures to optimize the diffusion layer can be effective in particular, fitting the nitriding hardness profile to the stressing profile and influencing the nitride precipitates (type, form, size, and distribution) The choice of material and heat treatment before nitriding depend on the required core strength Nitriding of alloyed steels often results in thick deposits of carbides and/or nitrides in the grain boundaries, which may promote crack initiation If the parts are to be under high load, such deposits should be avoided by changing the heat treatment process

The nitriding depth can be increased by increasing process temperature and/or treating time If merely prolonging the nitriding time does not result in sufficient nitriding depths, a two-step process can be performed In the first step, nitriding

is done at low temperatures to generate a finely dispersed nitride distribution The second nitriding step is done at higher temperatures to raise the nitriding depth This procedure avoids coarsening of the nitride distribution and thus a decrease

in hardness The cooling rate from the nitriding temperature and postnitriding treatments can be used to influence the type, size, and distribution of the nitrides

Corrosive Wear. Nitriding improves resistance to corrosive wear Resistance can be further improved by oxidation of the compound layer after nitriding To achieve a very smooth layer capable of sustaining high loads, it is recommended that the oxidized layer be polished and then oxidized again

References

1 H Klümper-Westkamp, F Hoffmann, and P Mayr, Härt.-Tech Mitt., Vol 44, 1989, p 346-355

2 K.H Zum Gahr, Reibungs- und Verschleissmodelle, Reibung und Verschleiss, Werkstoffeigen-schaften, K.H Zum Gahr, Ed., DGM, 1983, p 53-78

Mechanismen-Prüftechnik-3 "Verschleiss, Begriffe, Systemanalyse von Verschleissvorgängen, Gliederung des Verschleissgebietes," DIN 50 320, 1979

4 K.H Habig, Systemtechnik tribologischer Vorgänge, Reibung und Verschleiss, Werkstoffeigen-schaften, K.H Zum Gahr, Ed., DGM, 1983, p 13-28

Mechanismen-Prüftechnik-5 A Begelinger and A.W.S de Gee, TNO Delft Report, 1967

6 A Begelinger and A.W.S de Gee, Lubr Eng., Vol 16, 1970, p 56-63

7 H Mittmann and H Czichos, Materialprüfung, Vol 17, 1975, p 366-372

8 J Föhl, Reibung und Verschleiss, Mechanismen-Prüftechnik-Werkstoffeigen-schaften, K.H Zum Gahr, Ed.,

DGM, 1983, p 29-51

9 K.-H Habig, Verschleiss und Härte von Werkstoffen, Carl Hanser Verlag, 1980

10 K.-H Habig, Möglichkeiten der Model-Iverschleissprüfung, Werkstoffe und ihre Veredlung, Vol 2, 1980, p

229-232

11 W Schröter, W Uhlig, and G Alisch, Schmierungstechnik, Vol 11, 1980, p 9-15

12 A.E Miller, J Met., 1983, p 56-62

13 K.H Zum Gahr, Reibung und Verschleiss, Mechanismen-Prüftechnik-Werkstoffeigen-schaften, K.H Zum

Gahr, Ed., DGM, 1983, p 135-156

14 D Liedtke, Z Wirstsch Fertigung, Vol 65, 1970, p 234-237

15 J.C Gregory, Tribology Int., Vol 3, 1970, p 73-83

16 K.-H Habig, R Chatterjee-Fischer, and F Hoffmann, Härt.-Tech Mitt., Vol 33, 1978, p 28-35

17 K.-H Habig, W Evers, and R Chatterjee-Fischer, Härt.-Tech Mitt., Vol 33, 1978, p 272-280

18 C Gleave and M Farrow, Heat Treatment '81, The Metals Society, London, 1981, p 123-129

19 K.-H Habig, "Verschleissuntersuchungen an gas-, bad- und ionitriertem Stahl 42 CrMo 4," BAM-Berichte

No 38, June 1976

20 A Dubus and J.-P Peyre, Trait Therm., 1984, p 27-31

21 F Hoffmann, Lecture 38, Härtereikolloquium, 6-8 Octo 1982

22 D Römpler, Härt.-Tech Mitt., Vol 34, 1979, p 219-226

23 P Hammer and G Polzer, Fertigungstech Betr., Vol 15, 1965, p 498-500

24 J Zysk, Härt.-Tech Mitt., Vol 31, 1976, p 137-144

25 Fr.W Eysell, Durferr Hausmitt., Vol 28, 1958, p 7-18

26 H Tauscher and E Stecher, IfL-Mitt., Vol 4, 1965, p 220-224

27 H Tauscher and H Fleicher, Maschinenbautechnik, Vol 19, 1979, p 201-205

28 T.M Norn and L Kindbom, Stahl Eisen, Vol 78, 1958, p 1881-1891

29 D Edenhofer, Fachber Oberflächentech., Vol 12, 1974, p 97-102

30 C.S Nanjunda Ram and A Ramamohana, Materialprüf., Vol 1981, p 125-128

31 K Keller, Ind Anzeig., Vol 8.09, 1967, p 1580-1581

32 R Woska, Härt.-Tech Mitt., Vol 38, 1983, p 10-17

33 H Döpke, Z Wirtsch Fertigung, Vol 71, 1981, p 273-274

Glossary of Terms

Compiled by Peter J Blau, Metals and Ceramics Division, Oak Ridge National Laboratory

WORDS AND PHRASES are the essential tools of technical communication Without clearly defined terminology, technological progress is impossible The written and spoken word must be precise and mutually understood by all communicating individuals in the same context Unfortunately, the terminology used in friction, lubrication, and wear science and technology ("tribology") has evolved from sources in many different scientific and engineering fields Some terms (especially hyphenated terms) are coined during wear failure analysis and do not necessarily conform to existing terminology Consequently, the terminology varies greatly in origin and breadth of accepted usage Terms also vary in specificity; that is, many terms with slightly different connotations are sometimes erroneously substituted for one another For example, in certain instances the surface of a plain bearing might be described alternately as being polished, burnished, or scuffed Strictly speaking, these three terms are not equivalent The use of tribology terminology also varies

between different English-speaking countries A term such as density can be rather clearly and unambiguously defined,

but scuffing cannot because too many different interpretations of the term are in common use To avoid confusion and misunderstanding, terms which are known to have several possible interpretations should not be used

Any compilation of a glossary of terms in tribology becomes a difficult task because, to be truly successful, it should remove ambiguity and form a basis for future usage However, the glossary must also allow users to find definitions for terms that are deeply entrenched in the literature of the field even though they might be vague or inconsistently used Avoiding the pit-falls of providing too narrow a series of definitions and descriptions of terms necessitated using a variety

of references in assembling this glossary The definitions obtained from these references are supplemented by additional ones written specifically for this Volume in order to provide more complete coverage Sometimes the definitions presented in the cited references were inconsistent with the style or intent of this glossary and were edited When such editing was necessary, the reference number is followed by the letter "m." More than one entry was provided for certain terms in order to show differences in common usage In addition to the definitions given here, the reader is referred to articles dealing with specific fields where additional terminology is provided (for example, the parts of a bearing or the types of lubrication)

The basic form of an entry in the glossary contains the word or phrase in boldface type, one or more definitions and/or notes, a reference number (no reference number is given if the definition was written for this glossary), and additional information about synonyms and/or related terms as follows:

• term in boldface type

• (1) First definition Note: a note contains supplementary clarifying material not required in the formal definition (Ref X) (2) Another definition, possibly from another source (Ref Y) Synonym

or see also another term

Not all terms in the cited references were included here for various reasons For example, a given term might occasionally

be used in tribology, but is not principally associated with friction, lubrication, and wear (for example, belleville spring),

or the definition from one source was nearly identical to that from another and need to be repeated

While the editors have attempted to coordinate the use of terminology throughout this Volume, we can offer no guarantee that the terminology used in the other Sections of this Handbook is entirely consistent with the definitions presented in this glossary As with friction, lubrication, and wear technology as a whole, this Volume was compiled by individuals with backgrounds in many different fields

To further clarify certain terms in this glossary, illustrations or photographs are included Photographs, like written language, leave some room for precision, and examples were chosen to represent the general case of the phenomenon, not necessarily the only way in which a certain tribological phenomenon can appear

• a-spot

• One of many small contact areas through which electrical current can pass when two rough,

conductive, solid surfaces are touching Note: This term derives from the analysis of R Holm in

modeling electrical contact behavior A-spots may represent an area smaller than the total asperity contact area referred to in other surface contact models

• Abbott-Firestone curve

• Introduced in 1933 by E.J Abbott and F.A Firestone and sometimes called a "bearing area curve," the Abbott-Firestone curve is a plot of depth below a reference level parallel to a solid surface (ordinate) versus the percent bearing area intercepted by a horizontal line at that depth

(abscissa) Note: The Abbott-Firestone curve is sometimes used to compare the load-supporting

behavior of bearings with different surface roughnesses See also bearing area

• abrasion

• A process in which hard particles or protuberances are forced against and moving along a solid

surface Note: Sometimes this term is used to refer to abrasive wear See also abrasive erosion

Abrasive wear of the surface of 1020 steel abraded by 220 grit SiC paper showing characteristic grooves and attached, tiny cutting chips Courtesy of L.K Ives, NIST

• abrasive wear factor

• An empirical factor ( v) that expresses the influence of various service environments on the

increase in the radial clearance of a bearing (V) and, hence, the extent of life reduction as:

v = (V/e0)

• where e0 is a geometrical factor (Ref 3, p 699)

• abrasivity

• The extent to which a surface, particle, or collection of particles will tend to cause abrasive wear

when forced against a solid surface under relative motion and under prescribed conditions Note:

Specifically, an ASTM standard for the abrasivity of slurries ASTM G 75 has been developed

• absolute impact velocity

• See impact velocity

• accumulation period

• See preferred term acceleration period

• actual contact area

• The total area of contact formed by summing the localized asperity contact areas within the apparent area of contact Also known as real area of contact

• actual slip

• See macroslip

• additive

• In lubrication, a material added to a lubricant for the purpose of imparting new properties or of

enhancing existing properties Note: Main classes of additives include anti-corrosive , antifoam ,

antioxidant , antiwear , detergent , dispersant , extreme-pressure , and VI improver additives (Ref 1m)

• adherence

• In tribology, the physical attachment of material to a surface (either by adhesion or by other means of attachment) that results from the contact of two solid surfaces undergoing relative

motion Note: Adhesive bonding is not a requirement for adherence because mechanisms such as

mechanical interlocking of asperities can also provide a means for adherence See also adhesion (adhesive force) and adhesion, mechanical

• adhesion (adhesive force)

• In frictional contacts, the attractive force between adjacent surfaces Notes: In physical chemistry,

adhesion denotes the attraction between a solid surface and a second (liquid or solid) phase This definition is based on the assumption of a reversible equilibrium In mechanical technology, adhesion is generally irreversible In railway engineering, adhesion often means friction (Ref 1)

• (1) Wear by transference of material from one surface to another during relative motion due to a

process of solid-phase welding Note: Particles that are removed from one surface are either

permanently or temporarily attached to the other surface (Ref 1) (2) Wear due to localized bonding between contacting solid surfaces leading to material transfer between the two surfaces

or loss from either surface (Ref 2)

• AFS 50-70 test sand

• A rounded quartz sand specified for use as an abrasive in the dry sand-rubber wheel abrasive

wear test (ASTM G 65) Note: Using the U.S Sieve Series, none of this sand will be retained on

Sieve No 40, 5% maximum will be retained on Sieve No 50, 95% minimum will be retained on Sieve No 70, and none will pass Sieve No 100 This places all the particle diameters between

425 and 150 m

AFS 50-70 test sand particles used in the ASTM dry sand-rubber wheel abrasion test Courtesy of L.K Ives, NIST

• age hardening (of grease)

• The increasing consistency of a lubricating grease with time of storage (Ref 1)

ring that is rotating on a shaft Note: This term is based on manufacturer's designation, but has

become common terminology for referring to the block-on-ring geometry Further information can be found in ASTM Standard Test Method G 77

• Amontons' laws

• Two laws propounded in 1699 that state that (1) friction force is proportional to normal force, and (2) friction force is independent of the size of the contact area (Ref 1)

• Amsler wear machine

• A wear and traction-testing machine consisting of two disk-shaped specimens oriented such that their axes are parallel and whose circumferential, cylindrical surfaces are caused to roll or roll

and slide against one another Note: The rotation rates of each disk may be varied so as to

produce varying degrees of sliding and rolling motion

• angle of attack

• The angle between the direction of motion of an impinging liquid or solid particle and the tangent

to the surface at the point of impact (Ref 2)

• angle of contact

• In a ball race, the angle between a diametral plane perpendicular to a ball-bearing axis and a line drawn between points of tangency of the balls to the inner and outer rings (Ref 1)