Plasma surface engineering of low alloy steel

Materials Science and Engineering, A 140 ( 1991 ) 419-434

419

Plasma surface engineering of low alloy steel Y. Sun and T. Bell Wolfson Institute for Surface Engineering, School of Metallurgy and Materials, University of Birmingham, PO Box 363, Birmingham B15 2TT (U.K.)

Abstract

The surface of low alloy steel (En40B) has been engineered in the plasma of a glow discharge via plasma nitriding and ion plating of titanium nitride (TIN) coatings on the nitrided substrates with the purpose of enhancing the surface properties and fatigue strength. The nitriding response of the steel has been accessed by the evaluation of phase composition, layer thickness, hardness profile, residual stresses and nitrogen and carbon distributions. The wear and fatigue characteristics of the plasma-nitrided steel have been investigated and simple models have been developed to describe the influence of such properties as depth and strength of the nitrided case on the fatigue limit and load-bearing capacity of the nittided steel. In order to further improve the tribological properties and load-bearing capacity of the low alloy steel, a duplex plasma surface-engineering technique has been developed. This is achieved by plasma nitriding the steel first so as to produce a thick, strong subsurface and then depositing a thin, hard and wear-resistant TiN coating on the nitrided substrate by ion plating. Dry-sliding wear tests demonstrated that the duplex-treated steel, i.e. the TiN coating-nitrided steel composite, not only exhibited enhanced wear resistance over the as-nitrided steel (by a factor of 2-8) but also had much higher load-beating capacity than the TiN coating on unnitrided steel. Optimization of the coating-substrate combination can be achieved by correct control of the plasma-nitriding, surface preparation and ion-plating processes.

1. Introduction

Plasma surface-engineering technologies, owing to a number of advantages over conventional surface engineering, have found increasing applications in industry to engineer the surfaces of various steel components such as crankshafts, gears, bearings, dies, cutting tools, etc. [1-5]. Plasma nitriding and plasma-assisted physical vapour deposition (PVD) of ceramic coatings are the most widely used plasma surface-engineering techniques in industry. During the plasma-nitriding process, the nitriding reaction not only occurs at the surface but also in the subsurface owing to the long-distance diffusion of nitrogen atoms from the surface towards the core. As a result, a thin iron nitride layer is produced on the surface together with a relatively thick (about 0.5 mm) and strong (Vickers hardness about 1000 HV) diffusion zone in the subsurface of a steel component, which gradually reduces the hardness and nitrogen concentration towards the core, resulting in a diffuse case-core interface. This not 0921-5093/91/$3.50

only enhances the fatigue strength and wear resistance but also increases the load-bearing capacity of the component. On the other hand, plasma-assisted PVD of ceramic coatings, such as ion plating of thin (1-10 /zm), hard titanium nitride (TIN) coatings, owing to their unique physical, mechanical and metallurgical properties, offers superior tribological properties to the component [4, 6-8]. During the PVD process, the reaction between nitrogen and metal (titanium) atoms occurs only at or near the surface, resulting in a thin coating on the surface with an abrupt coating-substrate interface. Obviously, the remarkable properties of the thin coatings are employed only when the integrity between the coating and substrate is maintained during service [9, 10]. This requires that the substrate should be sufficiently strong to support the coating. A duplex plasma surface-engineering technique, i.e. combined plasma nitriding and PVD of ceramic coatings, has recently been developed in an attempt to further improve the tribological © Elsevier Sequoia/Printed in The Netherlands

420

performance and load-bearing capacity of low alloy steel [11]. This is achieved by plasma nitriding the low alloy steel first and then depositing a thin (2-4/~m thick) TiN coating on the nitrided surface by plasma-assisted PVD. The combination of these two processes produces a composite with a thin, hard and wear-resistant TiN coating supported by a thick nitrogen-strengthened subsurface. The purpose of the present paper is to discuss the potential of both individual and duplex plasma surface engineering in improving the fatigue and wear performance of low alloy steel.

2. Individual plasma surface engineering (nitriding) 2.1. Plasma-nitriding response A series of experiments has been carried out to investigate the nitriding response of commercial En40B steel with the following composition: 0.25 wt.% C, 3.2 wt.% Cr, 0.02 wt.% V, 0.03 wt.% AI and 0.52 wt.% Mn. In order to investigate the influence of carbon content in the substrate on the nitriding response, alloys with 3.2 wt.% Cr and various carbon contents were produced by casting in a metal mould [12]. Prior to surface engineering, all the steels were hardened and tempered. Figure 1 is typical of the plasma-nitrided En40B steel produced, comprising an outer slowetching iron nitride (compound) layer adjacent to a strengthened diffusion zone or nitrided case. The properties of a plasma-nitrided steel component are determined by both the core strength and the structural characteristics of the compound layer and the diffusion zone. Table 1 schematically illustrates the influence of process parameters and substrate carbon content on the structures of plasma-nitrided low alloy steel. 2.1.1. The compound layer It is well recognized that for each fixed process condition there exists a critical nitrogen potential, called the threshold nitriding potential, below which an iron nitride layer does not form on the surface during nitriding [13, 14]. Nitriding without iron nitride layer formation is usually called bright nitriding. Figure 2 is a threshold nitriding potential curve experimentally established for En40B steel at 480 °C [15]. The time dependence of the threshold nitriding potential is in agreement with the fact that there exists an incubation

r

l~ig. 1. Plasma-nitrided E n 4 0 B steel showing iron nitride layer and diffusion zone beneath.

TABLE 1 Schematic illustration of the influence of process parameters and substrate carbon content on the structures of plasmanitrided low alloy steel

Tt tt N2:H2~ Wt.% C in plasma t Wt.% C in substrate t

Iron nitride layer

Diffusion zone

e: y'

6~

H~

6

o~

Dec.

$ $ ~, T

t 1' ,~ .

a a ~,

t t a

~

t t ~. 1

t

t

~

~

.

. ~

{

.

--

Legend: 6, thickness; H~, Vickers hardness; o~, residual stress; Dec., decarburization; T, temperature; t, time. aSee text.

time for the formation of the iron nitride layer. Indeed, during nitriding, the surface nitrogen concentration is built up gradually with time and the incubation time for the formation of the iron nitride layer is obviously the time necessary for the build-up of the critical nitrogen concentration

421 TABLE 2

25

I

20

Z

En40B,480°C

e: y' ratio and thickness of the compound layer as a function of substrate carbon content at 500 °C for 10 h in ammonia

+

[] withC.L. b

15

0

[ withoutC, L.

0

10

20 30 40 NITRID1NG TIME, h

Fig. 2. Threshold En40B at 480 °C.

plasma-nitriding

50

Wt.% C

e: y'

Thickness (/~m)

0.19 0.25 0.34 0.40 0.50 0.57

0.58 0.71 0.83 1.03 1.19 1.29

5.5 6.5 7.4 8.1 9.O 10.2

60

potential

curve

for 8

on the surface [16]. Obviously, the threshold nitriding potential curve is temperature dependent. In the case of conventional gas nitriding it was observed that the curve moves upwards with increasing temperature [17], whilst in plasma nitriding the opposite trend was observed [11, 18]. This could be explained by the different nitriding and mass transfer mechanisms occurring in the gaseous and plasma processes. When the nitriding potential (i.e. the Nz:H z ratio) is above the threshold value, a compound layer consisting of y'-Fe4N and/or e-Fe z 3N will form on the surface of a steel component. Experiments demonstrated that even in carbon-free nitrogenous plasma a polyphase ( y ' + e ) compound layer was usually produced on the surface of low alloy steel. A monophase (y') iron nitride layer could be produced only when the nitriding temperature was very high or the nitriding potential was sufficiently low. With decreasing nitriding temperature and time or increasing nitriding potential the relative amount of e phase formed in the compound layer was increased (Table 1 ). This can be explained by the nitrogen uptake and redistribution of carbon in the nitrided case during plasma nitriding. According to the new ternary F e - N - C phase diagram [19, 20], y'-Fe4 N has only a limited solubility of 0.2 wt.% for carbon, whilst the solubility of carbon in e-Fez_3N is up to 3.6 wt.%. Accordingly, carbon facilitates the formation of e phase in the compound layer. Indeed, the carbon content in the steel had a significant influence on the phase composition of the compound layer. The e: y' ratio in the compound layer increased with increasing carbon content in the substrate (Table 2). Similarly, the introduction of carbon-containing gases into the nitriding

En4OB, 500°C, 10 h

<~ ~ 6 ~ ~ 4

zN2 0

0

5

10 15 20 NITROGEN CONTENT, %

25

Fig. 3. Iron nitride layer thickness as a function of nitriding potential.

atmosphere resulted in a compound layer composed of mainly e phase. By careful control of the carbon content in the plasma it is possible to produce a monophase (e) iron nitride layer which has superior resistance to seizure and scuffing. The development of the iron nitride layer is determined by nitriding temperature and time as well as nitriding potential. Decreasing the nitriding potential, i.e. the Nz:H 2 ratio in the nitriding atmosphere, led to a considerable reduction in the thickness of the iron nitride layer (Fig. 3). Therefore an iron nitride layer of high quality and desired thickness can be readily produced by monitoring the flow rate of N 2 and H 2 to the chamber. However, the development of the iron nitride layer only approximated to a parabolic dependence on time for the first few hours of nitriding, until a limiting thickness was reached, beyond which no further growth was observed. This phenomenon has also been observed by other investigators in plasma nitriding of a variety of steels [2, 21]. It is believed that cathode sputtering plays an important role in determining the growth of the compound layer [1, 22].

422

The carbon content in the steel also has a significant influence on the development of the compound layer. Under the same plasma process conditions the compound layer thickness increases with increasing carbon content in the substrate, e.g. increasing the carbon content from 0.19 to 0.57 wt.% increased the compound layer thickness from 5.5 to 10.2/~m at 520 °C for 10 h (Table 2). Clearly, the rate of growth of e-Fe2_3N is greater than that of ),'-FeaN owing to the wider homogeneity range of the e phase. Increasing the carbon content in steel increases the amount of e phase in the compound layer, thus giving rise to the growth rate of the layer. 21.2. The diffusion zone As a result of the diffusion of nitrogen from the surface towards the core of the steel, several reactions occur simultaneously in the diffusion zone. These include precipitation of metal nitrides, saturation of a-Fe lattices with nitrogen, residual stress generation, carbon redistribution and grain boundary phase formation [12]. Most of the nitrogen atoms in the diffusion zone (nitrided case) combine with chromium in the investigated F e - C r - C steels to precipitate CrN, which is the main cause of the hardening effect induced by nitriding [23]. The generation of compressive residual stresses in the nitrided case is the result of saturation of nitrogen in a-Fe and precipitation of metal nitrides, i.e. CrN. The compressive stresses and the uptake of nitrogen in the nitrided case cause carbon redistribution: the carbon atoms initially in the steel will diffuse to the stress-free regions, i.e. towards the surface and the nitriding front, leading to the decarburization of the nitrided case, the formation of a carbonrich zone in the nitriding front and the formation of a grain boundary phase parallel to the surface [I2]. The hardness profile produced by nitriding reflects the quantity, size and distribution of fine scale of CrN precipitates, which are dependent upon the nitriding temperature and concentrations of chromium and carbon in the steel as well as the nitriding potential. Figure 4 shows the surface hardness of nitrided En40B steel as a function of nitriding temperature. Clearly, optimum hardness can be achieved at temperatures between 480 and 520 °C. Reducing the nitriding potential (N2:H z ratio) below the threshold value (Fig. 2) leads to a significant drop in the hardening effect owing to the relatively low nitrogen

En40B nitridedfor 10 h in Nil 3 1100

~4 900

700 < e~ i

I

I

450

400

I

I

I

500

r

550

600

NITRIDING TEMPERATURE,o C Fig. 4. Variation of peak hardness of the nitrided case with nitriding temperature.

1300

520°C, 10 h, ammonia ~

0.19wt%C

700500 300 100

i

o.00

I

0.10

~

I

0.20

J

I

~

0.30

I

0.40

DISTANCE FROM SURFACE, mm Fig, 5. Effect of substrate carbon content on the hardness profile produced by plasma nitriding.

uptake in the diffusion zone. The effect of substrate carbon content on the hardness profile is shown in Fig. 5. The more carbon in the steel, the more chromium is tied up in the form of carbides; thus less chromium is immediately available to form CrN during nitriding and less hardness will develop. The development of the nitrided case is indeed a diffusion-controlled process. Figure 6 shows the nitrided case depth as a function of the square root of nitriding time at a temperature of 500 °C. The nitrided case depth can be described by the internal nitriding model based on the kinetics of internal oxidation [13, 23]. However, when the nitriding temperature was near to or higher than the tempering temperature of the steel, the case depth did not exhibit a linear relation with the square root of nitriding time, particularly when the nitriding time was long [24]. This suggests that overtempering of the substrate during nitriding, which not only causes softening of the substrate

423

0.50 E

En40B nitrided at

300

0.40

#.

0.30

~6

100 kta

r> 0.20

~

-100

< ~

-300

~

-500

¢*3

<

o.10 0.00

u.o

'

~

2.0

4.0

. . . . . . 6.0

8.0

~,

t%C t%C 0.50wt%C

SQUARE ROOT OF NITRIDING TIME, h Fig, 6. Nitrided case depth as a function of square root of nitriding time at 500 °C in cracked ammonia.

-700

0.0

i

i

l

i

0.2

0.4

0•6

0.8

1.0

DISTANCE FROM SURFACE, mm Fig. 8. Residual stress profiles measured using hole-drilling technique.

20

4500C [ 520oc 570°(7

15 cq

,

10

0

5

13 •

'

' 0 20

0 30

, 0.40

0.50

I 0.60

CARBON CONTENT, wt% Fig. 7. Diffusion coefficient of nitrogen in the steel as a function of substrate carbon content at various temperatures•

but also leads to the redistribution of chromium and carbon in the substrate, has a significant influence on the development of the nitrided case. Figure 7 shows the experimentally measured diffusion coefficient of nitrogen as a function of carbon content in the steel at various temperatures. Clearly, the diffusion coefficient is reduced as the carbon content in the substrate is increased. Increasing temperature is more effective in enhancing the diffusion of nitrogen in low carbon steels than in high carbon steels. This is because more interstitial sites are occupied by carbon in high carbon steels. During nitriding, the nitrided case tends to expand and is restrained by the unnitrided core. As a result, a compressive residual stress develops near the surface while a tensile residual stress evolves in the core [12, 25]. The level of residual stresses depends not only on the plasma process condition but also on the composition of the base material. Unalloyed carbon steels showed comparatively small stresses after nitrid-

ing. However, plasma nitriding of the chromiumbearing steels under investigation yielded much higher compressive residual stresses. Figure 8 shows the residual stress profiles measured using the hole-drilling technique [12]. Clearly, the compressive residual stress in the nitrided case decreases with increasing carbon content in the steel, corresponding to the observed reduction in hardening effect as the carbon content is increased (Fig. 5). Nevertheless, no direct relationship between microhardness and residual stress level has been observed, although the origin of residual stress generation and hardness increase is similar, i.e. nitrogen uptake and nitride precipitation. For instance, increasing the nitriding temperature from 450 to 520 °C increased the peak Vickers hardness from 780 to 940 HV but decreased the surface compressive residual stress from 2300 to 520 MPa. This arises from the fact that hardness increase is more directly related to nitrogen uptake and nitride precipitation, whilst residual stress generation is not only related to these but also related to carbon redistribution, carbide dissolution and formation of grain boundary phase in the nitrided case. Figure 9 shows the typical nitrogen and carbon concentration profiles of plasma-nitrided En40B steel. As a result of nitrogen uptake in the nitrided case, a decarburized zone near the surface and a carbon-rich zone in the nitriding front have been developed. This phenomenon has been observed in both conventional gasnitrided steel [26] and plasma-nitrided steel [15]. However, decarburization of the nitrided case is more pronounced during the plasma process than

424

z~

1.50

0.30

1100

1.25

0.25

1000

1.00

0.20

0.75

0.15 O

-.

900

,..~,,, " ~ ,--,,,~ -, ~ , . , , . , ~ ~ ~,~,,~ -,,~.'

z" r¢

700

0.50

0.10

0.25

0.05

ex~ 500

0.00

400

nitridingtime (h) -

, ~ ~'''---"-- ~ ~

~

70-,,40 - 10~

Z 0.00 O.00

0.I0

0.20

0.30

0.40

"~

..........

0.50

DISTANCE FROM SURFACE, mm Fig. 9. Typical nitrogen and carbon concentration profiles of plasma-nitrided En40B steel.

during the gaseous process Owing to the cathode sputtering involved in the plasma process. As a result, the precipitation of Fe3(N,C ) phase in the prior austenite grain boundaries is less pronounced in plasma-nitrided steel than in gasnitrided steel. This may account for the high ductility of the plasma-nitrided case observed by Edenhofer [2]. With increasing nitriding temperature and time and decreasing nitriding potential the extent of decarburization increased. This explains the observed variation of the e:y' ratio in the compound layer with nitriding temperature, time and potential (Table 1). 2.2. Fatigue and wear characteristics of plasmanitrided steel 2.2.1. Fatigue Wohler rotating-bending fatigue testing has been used to access the fatigue behaviour of plasma-nitrided steel. The Wohler fatigue specimens were machined from bars of En40B steel 10 mm in diameter. The test pieces were subsequently hardened and tempered to produce a core hardness of 270 HV. Fatigue testing was conducted employing a uniform frequency of 3000 rev min -1 [15]. In order to investigate the influence of compound layer thickness on fatigue strength, plasma nitriding of the fatigue specimens was carried out at 480 °C for 40 h at and above the threshold potential to produce a variety of iron nitride layer thicknesses with similar nitrided case depth. Fatigue testing demonstrated that all the specimens exhibited a similar fatigue limit irrespective of the iron nitride layer thickness, and it was the nitrided case which had a dominating effect on the fatigue behaviour of plasma-nitrided specimens [27]. Bright nitriding well below the thresh-

30010 4 '

10 5

10 6

10 7

108

NUMBER OF CYCLES, N Fig. 10. S - N curves of specimens (En40B) bright plasma nitrided at 480 °C for various times indicated.

old potential produced a relatively shallow case, thereby leading to a significant drop in fatigue limit. The fact that the iron nitride layer has no dominating influence on the fatigue behaviour can be explained by the fatigue fracture mechanism of plasma-nitrided specimens. All the fractured test pieces were observed to have failed by the 'fish eye" phenomenon with fatigue cracks originating from non-metallic inclusions at the case-core transition zone [27, 28]. It is envisaged, therefore, that a compound layer on the surface has no dominating influence on the initiation and propagation of fatigue cracks, and increasing the nitrided case depth increases the fatigue limit of the steel. Figure 10 shows the S - N curves for the specimens bright nitrided at 480 °C for various times. Indeed, the fatigue limit increased with increasing nitrided case depth. Improvements of 34% and 75% in fatigue limit were achieved with nitrided cases 0.12 and 0.56 mm thick respectively. The effect of increasing case depth can be viewed as effectively moving the fatigue crack initiation site further into the core, indicating that a greater applied bending stress will be required at the surface to create a sufficiently high level of stress at the case-core interface to initiate failure. 2.2.2 Wear Wear testing was carried out using an Amsler machine under combined rolling-sliding contact traction. The tests were conducted dry and were so arranged that specimens (50 mm in diameter and 10 mm thick) of identical treatment ran against each other. The upper specimen was rotating at a surface speed of 55 m min- ~and the lower specimen at 60 m min- 1 such that 9% slid-

425

ing occurred between the two contact surfaces [11]. Figure 11 shows the wear curves obtained for various specimens under an applied load of 200 kgf. No scuffing has been observed during the wear process and the iron nitride layer affected the wear behaviour only in the early stage, after which the wear rate was dominated by the nitrided case. From Fig. 11 it is evident that plasma nitriding substantially improves the wear performance of the low alloy steel. The degree of improvement depends on the plasma process conditions, i.e. temperature and time. Examination of the worn surface, subsurface and wear debris revealed that the main cause of the observed rapid wear of the unnitrided specimen was intense shear deformation occurring at the surface and in the subsurface. A shear zone approximately 200/~m thick had developed near the surface during the wear process (Fig. 12). Many voids and microcracks were induced in the subsurface of the unnitrided specimen and such cracks mostly propagated along the flow lines, resulting in the shearing away of filaments of the material of type "F" in Fig. 12, thereby leading to the generation of metallic debris [29]. When the test specimens were plasma nitrided at 500 °C for 10 and 35 h, the hard and thick case developed in these specimens prevented plastic deformation in the subsurface (Fig. 13(a)). Shearing was limited to a near-surface zone approximately 10 ~ m thick. As a result, oxidation wear dominated the wear process [30]. Clearly, a thick nitrided case is effective in preventing subsurface shearing during the wear process, so that wear occurs in a mild mode. However, when the nitriding time was short or the nitriding temperature was low, so that the nitrided case produced was relatively shallow, shearing occurred in the nitrided case and/or the

matrix beneath the nitrided case. In the case of the specimens nitrided at 420°C, where the nitrided case was only 60/~m thick and its hardness was much lower than that of the specimens nitrided at 500 °C (Fig. 4), intense cracking and shearing were observed after wear testing both in the nitrided case and in the matrix beneath the case (Fig. 13(c)). Clearly, this accounts for the increase in the wear rate after 25 000 revolutions of testing (Fig. 11), suggesting that the nitrided specimens failed to bear the applied load. In the specimens nitrided at 500°C for only 2.5 h, where the nitrided case was very hard but only 90 # m thick, most of the nitrided case was not deformed during wear but the matrix beneath the nitrided case was severely sheared (Fig. 13(b)). Such subcase shearing leads to the observed cracking in the nitrided case (Fig. 13(b)) and the increase in wear rate after 45 000 revolutions of testing (Fig. 11) via exfoliation [30]. Accordingly, the load-bearing capacity is associated with shear deformation in the contact body.

2.3. Design with nitriding It has been demonstrated that plasma nitriding can substantially improve the fatigue and wear

rolling-slidingweartest: dry, 200 kg unnitrided

60

5O 420oC P

40

//,t /

30

/f" .

0

0

~

500°C 2.5 h 35 h 10 h

10 20 30 40 50 60 70 REVOLUTIONS OF TESTING, xl000

Fig. 1 1. Rolling-sliding wear curves generated for various specimens indicated.

Fig. 12. Microsection of worn unnitrided En40B specimen showing subsurface shearing and elongation of the nonmetallic inclusions arrowed. Test condition: 200 kgf load, 20 000 revolutions of testing.

426

properties as well as the load-bearing capacity of low alloy steel. T h e - d e g r e e of improvement mainly depends on the depth and strength of the nitfided case. Under both rotating-bending fatigue and rolling-sliding contact conditions it is likely that fracture will initiate at the case-core transition zone. It is necessary, therefore, during engineering component design to determine the minimum case depth and hardness profile required to combat the applied stresses in a specified application, so that the optimum combination of plasma process parameters can be chosen. Accordingly, attempts have been made to develop simple models to quantitatively describe the influence of the properties of the nitride case on the rotating-bending fatigue limit and loadbearing capacity of plasma-nitrided low alloy steel.

2.3.1. Fatigue consideration Under the rotating-bending condition the applied stress (tensile) is distributed linearly from the surface towards the neutral axis of the specimen. Owing to the presence of a compressive residual stress in the nitrided case and a tensile residual stress in the unnitrided core, the resultant stress is decreased in the nitrided case and increased in the core (Fig. 14). Obviously, the case-core transition zone is the preferential site for fatigue crack initiation. From consideration of the physical geometry and applied stress pattern (Fig. 14) a simple model has been developed to describe the fatigue limit as a function of nitrided case depth [27], i.e. D O'lim= 2-------6 DOh

(l)

!

al~

i; residual stress

I

TM

~

i

applied stw.ss reatdtant stress

H . . . .

B

'

; !

. . . . . .

Fig. 13. Metallographic sections of worn En40B specimens plasma nitrided at 500°C for (a) 10 h, (b) 2.5 h and (c) at 420 *C for 10 h. Test condition: 200 kgf load, 65 000 revolutions of testing. Deformation (elongation) of non-metallic inclusions in the subsurface is arrowed.

Fig. 14. Schematic diagram showing the stress patterns of a loaded fatigue specimen.

427

where Olim is the fatigue limit, D is the diameter of the critical section of the fatigue specimen, 6 is the nitrided case depth and Oh is the maximum applied stress level in the case-core interface at the fatigue limit of the nitrided specimen (Fig. 14). Experiments have shown that, regardless of case depth, Oh remains constant for a given core hardness and residual stress. At the fatigue limit Olim the resultant stress (OH') at the case-core interface due to the applied stress (oh) and tensile residual stress (e c)is (Fig. 14) crw = Oh+ O~

(2)

all, can be regarded as the "subsurface" fatigue strength of the core material. This implies that the core will exhibit a fatigue limit of oH, if failures are initiated below the surface. Therefore o., is characteristic of the core strength only and is independent of the nitrided case depth and residual stress. From eqn. (2) it is clear that o h is determined by both the core strength and the tensile residual stress. When the depth of the nitrided case is much smaller than the diameter of the specimen, the tensile core residual stress generated during nitriding is very low in order to balance the compressive residual stress in the shallow nitrided case. In this case the fatigue limit Olim follows a linear relationship with D/(D - 26) (Fig. 15). It is believed that the linear relationship is valid where D / ( D - 26) ~<1.4, i.e. 6 <~D/7 [27]. Nevertheless, a very deep nitrided case (6 > D/7) can lead to a drop in the fatigue limit since a very high tensile residual stress is generated in the core during nitriding [28]. In this case Oilm no longer follows a linear relationship with D/(D - 26). 2.3.2. Contact consideration On the basis of the theory of elastic contact and the relationship between hardness and yield strength, a simple model has been developed to predict the shear behaviour of the contact bodies under rolling-sliding contact motion [31]. Modelling of the shear behaviour of the nitrided specimens comprises the determination of the shear yield strength profiles according to the measured hardness profiles of the nitrided specimens and the calculation of the stress field induced by the applied loads. The model, which was simplified by assuming uniform Young's modulus and coefficient of strain hardening, describes the influence of hardened depth and strength of the nitrided case on the initiation of

I

0.900 +1

f

800 /

7°°1

,/"

~600 ~

1.10

i i 1.20 1.30 D/(D-26)

1,40

Fig. 15. Fatigue limit vs. D / ( D - 2 6 ) curve of En40B, where D is the diameter of the fatigue specimen and ~ is the nitrided case depth.

\,0

t

k

,.

,, 00oc

",.

-'x._

:r ............. 0.05 0.10 0.15 0.20 0.25 0.30 DEPTH FROM SURFACE, mm

0.35

Fig, 16. Shear yield strength profiles of plasma-nitrided specimens and maximum shear stress profile (r) induced by the applied normal load of 200 kgf and tangential force of 0.6 x 200 kgf.

shear deformation in the contact bodies during the rolling-sliding process. Figure 16 shows the shear yield strength profiles converted from the hardness profiles for various specimens, together with the calculated maximum shear stress (rmax) profile induced under an applied normal force of 200 kgf and a tangential force of 0.6 x 200 kgf, where 0.6 is the value of the coefficient of friction of the wear couple, measured during wear tests [11]. Clearly, the introduction of a tangential force to the contact not only increases the stress level but also moves the peak stress value towards the surface. The influence of nitrided case depth and strength on the surface and subsurface shear behaviour during the wear process can be seen clearly in Fig. 16. Apparently, shear deformation or fracture is initiated at the point where the local stress exceeds the local strength. In the specimens

428

nitrided for only 2.5 h (at 500 °C) the maximum shear stress exceeds the local shear strength in the case-core transition zone (between A and B in Fig. 16), indicating that shear deformation will occur in this zone under such a loading condition. This accords with experimental observations (Fig. 13(b)). In the specimens nitrided at 420 °C, owing to the relatively low strength and shallow depth of the nitrided case, the shear stress exceeds the local shear strength not only in the case-core transition zone but also in the nitrided case (between A and C in Fig. 16). This explains the fact that intense cracking occurred in the subsurface of the test specimens during the wear process (Fig. 13(c)). Obviously, the result of increasing the nitrided case depth is that the local shear strength is greater than the local maximum shear stress everywhere in the contact body, so that shear is not possible, as in the specimens nitrided at 500 °C for 10 and 35 h (Figs. 16 and 13(a)). From the calculated rmax shear stress profile the hardness profile is obtained, which is required to resist the rmax profile so that shear does not occur during service (Fig. 17). According to the hardness profile, the minimum required case depth OD is determined. Obviously, the core strength is important: by increasing the core hardness from O to O', the case depth required to resist shear can be reduced from OD to O'D' (Fig. 17). In practice, however, the thickness of the nitrided case is reduced gradually during service owing to wear. Thus the case depth required is greater than that determined from the Zm~xprofile (Fig. 17). Assuming that the wear rate is R and the designed lifetime of the component is F, in

order that shear does not occur during service, the minimum nitrided case depth 6mi, should be 6min = O D + F R

Clearly, increasing the wear resistance of the nitrided surface can reduce the minimum case depth required to resist subcase shear. This can be achieved by optimizing the hardness of the nitrided case through controlled nitriding, and by depositing a thin, hard and wear-resistant ceramic coating such as TiN on the nitrided surface through duplex plasma surface engineering as discussed below.

3. Duplex plasma surface engineering

200

800

(3)

1"4

Experimental results presented in the previous sections demonstrated that plasma nitriding substantially enhances the load-bearing capacity of low alloy steel provided a sufficiently thick nitrided case is produced. It is thereby anticipated that if a thin, very hard and wear-resistant ceramic coating is applied on a plasma-nitrided surface, the nitrided case will serve as a strong support for the ceramic coating and bear the applied load during service, while the ceramic coating will offer superior wear resistance. Accordingly, attempts have been made to deposit thin TiN coatings by ion plating on previously plasma-nitrided substrates [11]. Prior to coating, the nitrided surface was cleaned by one of the following processes: chemical cleaning, dry blasting or slight grinding. Experimental results demonstrated that the "TIN coating-nitrided substrate" combination can be optimized by proper process control and the composite possessed superior metallurgical and mechanical properties over the nitrided steel and the "TIN coating-unnitrided steel" composite.

700 > -v

150

6OO HV

13'

D'

400O

---

100

D

<

300 200 ~"~'~'~

x max

5o

100 0

i

|

i

i

i

I

,

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0

DEPTH FROM SURFACE, mm Fig. 17. Minimum hardness profile required to resist shear, deduced from the calculated maximum shear stress profile.

3.1. TiN coating-nitrided substrate combination Figures 18(a) and 18(b) show the TiN coatings deposited on the nitrided surface with and without a compound (iron nitride) layer respectively. The TiN coatings were produced by ion plating at about 500 °C. In the case where the iron nitride layer formed during plasma nitriding was not removed prior to coating, a "black" layer about 1.5 /zm thick was observed between the TiN coating and the iron nitride layer in the composite (Fig. 18(a)). When the iron nitride layer was removed by slight grinding or the specimen was bright nitrided prior to coating, no "black" layer

429

Fig. 19. TiN coating produced at about 450°C on the nitrided En40B surface (with an iron nitride layer prior to coating). Optical micrograph of taper section.

Fig. 18. Micrographs showing ion-plated (at about 500°C) TiN coatings on nitrided surfaces. The iron nitride layer produced during plasma nitriding (a) was not removed and (b) was removed by slight grinding prior to coating.

was observed in the composite (Fig. 18(b)). It is evident that the "black" layer was transformed from the outer part of the previously present iron nitride layer during the PVD process since a reduction in the thickness of the previously present iron nitride layer has been observed after ion plating. Microhardness measurements revealed that the "black" layer had a relatively low hardness (400-500 HV) and nuclear reaction analysis of the duplex-treated specimen showed very low nitrogen content in the "black" layer [11], indicating that most of the nitrogen atoms in the outer part of the previously present iron nitride layer had been released during ion plating at 500 °C. In conjunction with the X-ray diffraction results [11], which showed that the "black" layer comprised mainly a-Fe, it can be concluded that the "black" layer was the product of the decomposition of the outer part of the iron nitride layer during the ion-plating process, which led to the transformation of Fe4N and Fe2_3N t o a-Fe. The formation and development of the "black" layer are determined by the decomposition of

iron nitride and the outward diffusion of the released nitrogen, both processes being dependent upon surface topography and PVD process temperature. Microblasting of the nitrided surface (with an iron nitride layer) prior to coating was effective in reducing the thickness of the "black" layer formed during the subsequent ionplating process. Blasting not only changed the surface topography but also effectively removed the porous outer part of the previously present iron nitride layer [11]. A study on the decomposition behaviour of the iron nitride layer in a denitriding pkasma (hydrogen) revealed that the iron nitride layer produced during plasma nitriding at 500 °C remained stable in hydrogen plasma until 450°C, above which decomposition occurred. This suggests that reducing the PVD process temperature would be effective in avoiding the formation of a soft "black" layer in the TiN coating-nitrided substrate composite. Indeed, PVD at a relatively low temperature (approximately 450°C) did not produce a "black" layer between the TiN coating and the iron nitride layer. Therefore a "TIN coating-iron nitride" combination was produced (Fig. 19). However, it should be noticed that reducing the PVD process temperature will deteriorate the coating-substrate adhesion [4]. Since the "black" layer only forms in the previously present iron nitride layer, a simple way to avoid its formation is to avoid the formation of an iron nitride (compound) layer during the plasmanitriding process via bright nitriding, or to remove the iron nitride layer before the coating process. As discussed in Section 2.1.2, in the

430

investigated F e - C r - C steel most of the nitrogen atoms in the nitrided case combine with chromium to precipitate CrN, which is much more stable than iron nitrides Fe4N and Fe2_3N [32]. At normal PVD process temperatures (about 500°C), decomposition of CrN is impossible. Thus a "TiN-nitrided case" combination was produced which exhibited the best tribological performance and highest load-bearing capacity as discussed below.

3.2. Wear behaviour of"TiN coating-nitrided steel" composite Pure sliding (ball-on-wheel) wear tests have been carried out to access the wear performance of the duplex-treated specimens, i.e. the "TiN coating-nitrided substrate" composites [11]. During the test, the specimen to be tested was running against an alumina ball (slider) 9 m m in diameter at a constant speed of 0.5 m s-L Characterization of the tested composites is given in Table 3. A comparison of the wear volume from various specimens after ball-on-wheel wear testing is given in Fig. 20. Clearly, in terms of

TABLE 3 Characterization of TiN coating-nitrided steel composites No.

Coating thickness

Coating hardness

Coating-substrate combination

2.2 2.3 2.5 2.4 1.2

1450 1900 2350 2100 1190

TiN-black layer TiN-black layer TiN-nitrided case TiN-iron nitride TiN-iron nitride

(~m) 1 2 3 4 5

(Hv)

improving wear performance, duplex plasma surface engineering (combined plasma nitriding-ion plating) is more effective than both individual plasma nitriding and individual ion plating of TiN. Among all the test specimens, including the hardened and tempered specimens, TiN coating on unnitrided substrate, as-nitrided specimens and duplex surface-engineered specimens, the duplex-treated specimens exhibited the lowest wear rate. The TiN coating on unnitrided steel surface did not offer as much wear protection as the as-nitrided steel, whilst the wear rate of the as-nitrided steel was reduced by applying a thin TiN coating by a factor of 2-8 depending on the hardness of the coating produced. Careful examination of Fig. 20 and Table 3 revealed that the rate of wear of a TiN coating on the nitrided substrate decreases with increasing coating hardness. Figure 21 is a plot of the wear resistance of the TiN coating (relative to that of the as-nitrided steel) against the hardness of the coating. The data points fall neatly along a straight line, indicating the linear relationship between relative wear resistance and hardness of the TiN coating. Examination of the worn surfaces and subsurfaces after wear testing revealed that intense shear deformation had occurred in the subsurface of the "TIN coating-unnitrided steel" composite during wear (Fig. 22), and as a result, the TiN coating failed in many areas of the wear track, leading to the observed roughening of the wear track surface and considerable material loss. Apparently, the unnitrided substrate, which had a mean hardness of about 310 HV, was insuffi-

~20N

10 103

slidingdistance:620 m

t~

Ball-on-wheeltest aluminaball, 20 N, 620 m.

8 6

10 2

4 i l01 > 10 o

101 Fig. 20. Wear volumes from various En40B specimens indicated.

<

2 . . . . . .

0 1000

. . . .

I

I

I

I

1400 1800 2200 2600 HARDNESS OF TiN COATING,HV

Fig. 21. Relative wear resistance of TiN coatings on nitrided En40B substrates as a function of coating hardness. The wear resistance of the as-nitrided specimen was designated as unity.

431

Fig. 22. Scanning electron micrographs showing the wear track produced on the TiN-unnitrided substrate composite. (a) Failure of the coating and (b) shearing in the subsurface are evident.

ciently strong to bear the applied load and support the coating during the wear process. Failure of the coating was initiated in the substrate rather than at the coating-substrate interface. On the other hand, in the "TIN coating-nitrided substrate" composite, owing to the thick and strong subsurface developed during the nitriding process, surface and subsurface shear deformation was prevented during the wear process and material loss from the coating was caused by the micropolishing action of the slider (alumina ball): the wear track surface shows a polished appearance (Fig. 23). At the onset of sliding, only the highest asperities on the coating surface were brought into contact with the slider (Fig. 23(b)). As sliding continued, the high asperities were smoothed and lower asperities were brought into contact (Fig. 23(c)). As a result, the wear track surface became very smooth compared with the as-coated surface. For the TiN coatings on plasma-nitrided substrates an approximate linear relationship has been observed between wear volume and sliding distance [11, 33].

Fig. 23. Surface appearance of the wear track produced on the TiN-nitrided steel composite under a 20 N load after (a) 0 m (as-coated surface), (b) 50 m and (c) 600 m sliding distance.

3.3. Load-bearing capacityof "TiN coating-nitrided steel" composite Figure 24 shows the wear volume-applied load curves for various specimens. The wear volume of the as-nitrided specimen increased gradually with increasing applied load. Nevertheless, for each investigated TiN coating-substrate composite system there existed some critical load

432

~

Ball-on-wheel test

50

alumina ball, dry, 320 m

--.40

b

c

d

30

© 20

0

I

0

10

40 50 20 30 APPLIED LOAD, N

!

!

60

70

Fig. 24. Wear volume v s . applied load curves for various coating-substrate combinations: a, TiN-black layer (1.5/~m thick); b, TiN-black layer (0.9 pm thick); c, TiN-iron nitride; d, TiN-nitrided case combination.

above which the coating failed catastrophically, leading to an abrupt increase in the wear rate by several orders of magnitude, whilst below the critical load the coating was worn very slowly in the mode described previously (Fig. 23) and failure of the coating did not occur. It is believed that the critical load measured in the ball-onwheel sliding test can be used as a measure of the load-bearing capacity of the TiN coating-substrate composite. It has been discussed that a variety of coating-substrate combinations could be produced depending on the plasma-nitriding, surface preparation and PVD process conditions. These include "TiN-unnitrided substrate", "TiN-black layer", 'q'iN-iron nitride" and 'q'iN-nltrided case" combinations (Figs. 18 and 19). The loadbearing capacity mainly depended upon the coating-substrate combination produced, as shown in Fig. 24. The "TiN-unnitrided substrate" combination had the lowest load-bearing capacity among all the combinations tested, obviously owing to the relatively low strength of the unnitrided substrate. When a TiN coating was deposited on a nitrided surface, the coated composite exhibited much higher load-bearing capacity than the TiN coating on the unnitrided substrate, with the "TiN-nitrided case" combination having the highest load-bearing capacity. The relatively low load-bearing capacity of the "TiN-black layer" combination was due to the low hardness of the black layer. During the wear process, shear deformation accumulated in the soft black layer, leading to cracking and premature fracture of the hard coating [11]. Figure 25 shows that cracking in the coating during sliding

Fig. 25. Micrographs showing cracking of the TiN coating due to deformation of the underneath "black layer" during the wear process. Wear test condition: 35 N load, 320 m sliding distance. (a) Optical micrograph of the cross-section. (b) Scanning electron micrograph of the worn surface.

was indeed induced by deformation of the "black" layer. It is envisaged that reducing the thickness of the black layer increases the loadbearing capacity of the composite (Fig. 24). The relatively low load-bearing capacity of the "TiN-iron nitride" combination was probably due to the highly brittle nature of the iron nitride layer, while the relatively poor coating-substrate adhesion resulted from the relatively low PVD process temperature. Clearly, slight grinding to remove the iron nitride layer produced during plasma nitriding is the suitable surface preparation process for the subsequent ion plating of TiN coatings on plasma-nitrided steel substrates. This combined process produced a composite not only with the highest wear resistance (Fig. 20) but also with the highest load-bearing capacity (Fig. 24). Although ion plating of TiN on the bright-nitrided surface also produced a "TiN-nitrided case" combination, the relatively low hardness of the nitrided case significantly worsened the load-bearing

433

capacity of the composite. In addition, microblasting of the nitrided surface prior to coating was also found to be a suitable surface preparation process for the subsequent ion plating of TiN. This is because blasting of the nitrided surface not only effectively reduces the thickness of the black layer formed during ion plating but also changes the surface topography of the substrate, providing a mechanical anchor between the coating and the substrate [34]. Similar conclusions have also been reached in rolling-sliding wear testing of various specimens [33]. 4. Conclusions

From the above results and discussion it can be seen that surface modification of low alloy steel via plasma surface engineering substantially improves the fatigue and wear behaviour as well as the load-beating capacity of the steel. The degree of improvement is dependent upon the structural characteristics of the surfaceengineered composite. A thin compound (iron nitride) layer on the nitrided surface has no dominating effect on the fatigue limit of the nitrided steel. Therefore it is anticipated that the TiN coating deposited on the nitrided surface would have no influence on the fatigue behaviour of the composite. Indeed, an evaluation of the fatigue performance of TiN-coated En19 steel revealed that the coating had no influence on the fatigue limit of either nitrided or unnitrided material [35]. Eventually, the fatigue strength only depends on such properties as depth and residual stresses of the nitrided case. Therefore it is necessary during nitriding practice to control the nitriding process so as to produce a suitable nitrided case to combat the applied stress. The minimum depth of the nitrided case required to resist a specific rotating-bending stress can be determined from eqn. (1). On the other hand, a thin ceramic coating (e.g. TiN) deposited on the nitrided surface provides superior wear resistance that cannot be obtained by either individual plasma nitriding or PVD ceramic coating on the unnitrided surface. Clearly, it is the combined effect of the thin ceramic coating and the strengthened subsurface produced by nitriding that provides the remarkable tribological properties and load-bearing capacity to the low alloy steel. During the wear process, the applied load is mainly carried by the nitrided subsurface. The load-beating capacity

depends not only on the depth of the nitrided case (Fig. 16) but also on the "TiN-substrate" combination produced (Fig. 24). It is essential to control the plasma-nitriding, surface preparation and ion-plating processes so that a composite with a TiN coating of high hardness, a sufficiently thick nitrided case and a "TiN-nitrided case" combination is produced and thus ensuring that the wear resistance and load-bearing capacity are optimized.

References 1 T. Bell and V. Korotchenko, Heat Treat. Met., 4 (1978) 88. 2 B. Edenhofer, Heat Treat. Met., 1 (1) (1974) 23; l (2) (1974) 59. 3 C. K. Jones, S. W. Martin, D. J. Sturges and M. Hudis, Heat Treatment '73, London, The Metal Society, London, 1975, p. 71. 4 A. Matthews, Surf. Eng., 1 (1985) 93. 5 K.-T. Rie, Plasma Surface Engineering, Vol. 1, DGM Informationsgesellschaft, Oberursel, 1989, p. 201. 6 R. E Bunshah and C. V. Deshpandey, Surf Coat. Technol., 27(1986) 1. 7 V. Murawa, Heat. Treat. Met., •2(2)(1985) 49. 8 J.E. Sundgren, Thin Solid Films, 128 ( 1985 ) 21. 9 B.M. Kramer, Thin Solid Films, 108(1983) 117. 10 S. Ramalingam, Thin Solid Films, 118 (1984) 335. 11 Y. Sun, Ph.D. Thesis, University of Birmingham, 1989. 12 M. A. Nosratinia, Ph.D. Thesis, University of Birmingham, 1990. 13 B. J. Lightfoot and D. H. Jack, Heat Treatment '73, London, The Metal Society, London, 1975, p. 60. 14 J. Pan, M. Hu, L. Mao, W. Tang, Z. Xu and Y. Tang, Heat Treatment '83, London, The Metal Society, London, 1984, p. 1.11. 15 N.L. Loh, Ph.D. Thesis, University of Liverpool, 1980. 16 H.C.F. Rozendaal, E. J. Mittemeijer, E F. Colijn and P. J. Van Der Schaaf, Metall. Trans. A, 14(1983) 395. 17 T. Bell, B. J. Birch, V. Korotchenko and S. P. Evans, Heat Treatment '73, London, The Metal Society, London, 1975,p. 51. 18 A.M. Staines, personal communication, 1989. 19 J. Slycke, L. Spoge and J. Agren, Scand. J. MetalL, 17 (1988) 122. 20 Z. Xu and L. Lin, Acta Metall. Sinica, 1 (1988) 1 (English edition). 21 B. Edenhofer and T. J. Bewley, Heat Treatment '76, London, The Metal Society, London, 1978, p. 7. 22 A. Marciniak, Surf. Eng., 1 (1985) 283. 23 K. H. Jack, Heat Treatment '73, London, The Metal Society, London, 1975, p. 39. 24 M.B. Karamis and A. M. Staines, Heat Treat. Met., •9(3) (1989) 79. 25 B. K. Jones and J. W. Martin, Met. Technol., 4 (1977) 52O. 26 S. Mridha and D. H. Jack, Met. Sci., 16 (1982) 398. 27 T. Bell and N. L. Loh, J. Heat. Treat., 2 (1982) 232.

434 28 B. K. Jones and J. W. Martin, Met. Technol., 5 (1978) 217. 29 Y. Sun, P. A. Dearnley and T. Bell, 1PAT '87, 6th Int. Conf. on Ion and Plasma Assisted Techniques, Brighton, May 1987, CEP Consultants, Edinburgh, 1987, p. 234. 30 Y. Sun, P. A. Dearnley and T. Bell, Plasma Surface Engineering, Vol. 2, DGM Informationsgesellschaft,

Oberursel, 1989, p. 927. 31 T. Bell and Y. Sun, Surf. Eng., 6 (1990) 133. 32 A. U. Seybolt, Trans. Metall. Soc. A1ME, 245 (1969) 769. 33 T. Bell and Y. Sun, 7th Int. Conf. for the Heat Treatment of Material, Moscow, December 1990, in the press. 34 D. M. Mattox, Thin Solid Films, 124(1985) 3. 35 T. Bell, Surf. Eng., 6(1990) 31.