An investigation of the tribological behaviour of a high-nitrogen CrMn austenitic stainless steel

WEAR ELSEVIER

Wear 215 { 1~)8) 83-th)

An investigation of the tribological behaviour of a high-nitrogen Cr-Mn austenitic stainless steel D.J. M i l l s t R . D . Knutsen * D('partmcm tft*M~teriaL~"Engtnecru;.~. fJ#ltrt,r::lly qt ('ope Towtl. R, md~'hosch. Cape Pr*~vince. .~tmth A[rica

Received I Novemtx,r It~-graccepled 14 N~Jven,hcr !t~.,~7

Abstract

The tribt)logical bchaviour of a novel high-nitrogen Cr-Mn austenitic stainless steel ( referred to as CrMnN-steel ) has been investigated using dry abrasion, abraskm-corrosion, and cavitauon erosion te~t techniques. The pcrfon,ance of CrMnN-s(¢ct has ~ ¢ n compared to AISI 304 stainless steel and Hadtield manganese steel. Hadfleld steel demonstrate~ th~ highest abrasion resistance, but CrMnN-steeF is far superior when a corrosive environment is introduced. The cavitation erosion rcsi:,lance of CrMuN-stcel in distilled water is greater than bo4h AIS! 3(M and Hadfield steel. © 1998 Elsevier Science S.A. Keyu'ord.,: High-nitrogensteel.,Abrasion:Cavhalion oo~itm

!. I n t r o d u c t i o n

The ongoing research and development of stainless steel continues to produce new ideas for improving the mechanical properties and corrosion properties of this important class of engineering materials. Considerable emphasis has been placed on improving steelmaking practice with the intcntitm o f improving both the cost and pertbrmauce of the final product. Of course, not all classes o f stainless steel have advanced at the same rate. and emphasis has been divided over the years. Ferritic stainless steel received .'nuch attention during the F970s with the introduction of the argon-oxygen decarburisation ( A O D ) process thai enabled a reduction in the alloy carbon content. This technology led to renewed interest in the t'erritics as a result of the marked improvement in toughness and weldability properties. The development of austenitic stainless st::el with improved properties was initiated in the 1960s and became widespread in the 1980s. Earlier improvements related to increases in Cr, Mo and Ni contents and. more recently, much interest has been expressed in raising the level of dissolved nitrogen in steel. The latter development of the so-called highnitrogen austenitic stainless steel ( HNS ), with nitrogen levels sometimes exceeding 0.5 wt.¢~, has resulted in austenitic steel w ith an exceptional combination of strength, toughness, cold

work capacily and co;'rosion resistance ! ! l- The production of such steel i!~not trivial in view of the limited solubility o f nitr,-~gen in liquid metal, and initial success wa~ only achieved using the high-pressure electroslag remeiting process. However, re.,~arch on the influence of alloying elements on nitrogen solubility has made it possible to produce stainless steel with nitrob,en contents up to I wt.f/~• by melting m atmospheric pressure on condition that a well-balauccd combination o f alloying elcmems is umd [ 2 [. The :;uccess of this operation depends largely on the o p t i m i ~ t i o n of the Ni and Mn levels, and it is shown that the nitrogen solubility in liquid steel i n c r e a ~ s with increasing Mn level for a fixed casting temperature ] 2 ]. This approach has been successfully employed on a steel plant, and HNS alloys have been produced at ~he 30t level with nitrogen contents between 0.45 and 0.65 wt,~.2 The composition o f this alloy, designated CrMnN-steei. is shown in Table I. This recently established production route for HNS alloys otters attractively priced stainless steel for applicationswbere a combination of good strength and toughness, high workhardening capacity, and corrosion resistance is required. The.,~ attributes suggest that steel could perform well in various tribological situations [3 I. particularly where hostile environments are involved, which would normally lead to extensive corrosive attack of conventional wear-resistant

* Corres|~mdin~authur. Universityuf Caua:Town. Dept.of MaterialsEng.. University Private Ba~. Rondebosch 77(}0, South Afri~:a. = Presentlywith Boart-LongycarResearchCentre, Gttuleng.South Africa.

2C.A.P_Rennhard. High-nitrogenstainles~steel in South Africa: Prt~iuclion. properties and applications. Paper presented at the SAIMM Annual Minerals ProcessingSymposium.Cape Town. Aug. 1~)5 [ unpublislled).

[X~.3-1048/98/$I9.(~) ~b ItJ98 Ebevicr Science S.A. All rights remrvcd PIIS()043- 164.8t 97 )OO273- t

84

D.J. Mills. R.D. K , IItselt / Wear 215 ¢1998) 83-90

2. E x p e r i m e n t a l a p p r o a c h

In order to simulate the consequence of a liquid corrosive medium flowing over an abraded surface, dry abrasion intervals of only I m were followed by a 22-1, immersion period of the abraded surface in synthetic mine water [41. Thus, each cycle of abrasion-corrosion represents I m of dry abrasion and 22 h of corrosion. Volume loss after each of four cycles was determined in the usual way, and the relative abrasion-corrosion wear rate (RWR) was measured against mild steel:

2. L Mawrials

RWR =

The CrMnN-steel alloy was produced as a plant trial and was supplied for investigation in both the annealed and several cold-worked conditions. Cold work reductions were carried out at 16%. 30%, 48% and 66% levels, and specimens were tested in each of these conditions. The average tensile yield and ultimate tensile strength va!ucs of the annealed CrMnN-steel were measured as 570 MPa and 880 MPa, n,~pectively. Hadfield steel and AISI 304 were tested in the as-received annealed condition, and the compositions of these alloys are included in Table I.

Cavitation erosion was carried out in distilled water at a stabilised temperature (25°C). The equipment consists of a high-frequency generator, commercial ultrasonic drill, amplifying horn, and drill tip. The experimental conditions were standardised according to the procedures described by Heathcock 15 I: vibration amplitude of 70 p.m at a frequency of 20 kHz, and a separation between the sample and drill tip of 0.35 mm. Specimens were mechanically polished prior to testing ( final polish using 1 /xm diamond paste) and I cm-" of polished surface was exposed to cavitation erosion. Weight loss measurements were determined at hourly intervals up to a maximum exposure period of ! 2 h. Weight loss was measured to the nearest 0.01 mg. and the results were subsequently converted to volume loss values.

steel. To this end, labmatory scale wear tests, including dry abrasion, corrosion-abrasion, and cavitation erosion, were carded oat on the novel CrMnN-steel alloy, and the results were compared to conventional steel such as AiSi 304 austenitic stainless steel and Hadfield steel, which were tested under the same conditions.

2.2. Wear testing Two principal types of wear tests were employed in this study: dry abrasion testing on a modified belt sander, and cavitation erosion using an ultrasonic laboratory cavitation rig. A corrosion component based on South African gold mine water was added to the abrasion test procedure in order to introduce a synergistic corrosion-abrasion environment. Abrasion behaviour was characterised by the pin-on-abrasive method according to the procedures described by Barker and Ball [ 4 I- Alumina abrasive belts ( 80 grit) were employed, and the specimen ( 10× 10 mm) was loaded onto the belt surface at a constant force of 32. I N. During testing, the belt speed was maintained at 260 m m s - t, and the specimen was traversed across the belt to continuously wear against virgin abrasive. Specimens were abraded at 3-m intervals up to a total abrasive path length of 12 m. Weight loss was measured nearest 0.01 mg ) alter each 3-m interval and subsequently converted to volume loss. The relative abrasion resistance cRAR) of the specimens was ranked against mild steel according to ',he following relationship: RAR =

Total volume loss of mild steel Total volume loss of test materials

2.3. Mi('ro.',copy Specimens for light microscopy examination were polished by conventional metallographic techniques and etched in 10c/~oxalic acid at 10 V for 30 s. Imaging was carried out using a Reichert MeF3A microscope in the Nomarski interference contrast mode. Thin foils were prepared for transmission electron microscopy (TEM) in a Struers Tenupoljet polisher. The polishing solution comprised 5% perchloric acid in ethanol at - 15°C, and polishing was carried out at 35 V. Wear surfaces were examined using secondary electron imaging in a Cambridge $200 scanning electron microscope (SEM). In order to investigate the mechanism of material removal duff ng cavitation erosion, specimens were examined alter 30-min intervals of exposure, particularly during the initial stages of material loss.

Volume loss per metre of abrasion of mild ~teel V t d u m e loss p e r m e t r e tff ahr',t,,ion o f test m a t e r i a l s

Table I C ~ ' m i c a l conqx),,ilion of CrMnN-sl¢¢l. AISI 3tlJ, and Hadlicld Steel ! ia

3. Results

3. i. Mi,*roslr:,~ture Alloy

('~

Mn

N

Ni

Mo

C

Si

CrMnN-~leel

I*,~

In

().63

0.8

().11~

A|SI ~n4

18

H a d l i d d sit'el

I+ I

1,2 12

(I.(12

O.03

I),n3

0.3

--

0.05

0.02

O. ]

-

I. 17

I).gl

The microslructure liar CrMnN-stcel in the annealed condition is presented in Fig. I. The austenitic grain structure is more or Icss equiaxed, and the grain size varies from 10/~m to al'amt 50/xm. The microstructure is not entirely austenitic. and approximately 5 vol.Ch, &ferrite occurs randomly dis-

D.J. Mill.~.R.D. grease,#/Wear 215 f 19~8) 83-90

85

microstmclures, although AISi 304 contains less than 2 vol.% &ferrite. Examination of thin foils of cold-worked CrMnN~.'teel was principally aimed at understanding the work-hardening behaviour of the alloy. Focus was directed towards the arrangement o f dislocations as a function o f cold work in order to establish the mode o f slip. All the different cold work conditions were examined, and planar slip was noted to dominate in all ca~s. An example of planar slip is demonstrated for the 16oh cold-worked condition in Fig. 3. 3.2. Dry abrasion Fig. I. Predominantly austenit¢ micrnstructure in CrMnN-ste¢l. The darker area~ arc &fertile (approximately 5 vol.~ ).

The dry abrasion wear results are presented for C r M n N steel. AISI 304. Hadfield steel and mild steel in Table 2. The conventional wear-resistant Hadlield steel performs 1.94 times better than mild steel, whereas CrMnN-steel has an RAR value of 1.53. Tim initial Vickers hardness (30 kgf) values prior to abrasion testing are included in Table 2 for comparison. The abrasion wear rate behaviour of C r M n N steel as a function of prior cold work is presented in Fig. 4. Volume loss increases steadily as a function of prior cold work. The hardness increase due to cold work is also shown in Fig. 4. 3.3. Abnrvion-cnrrosion

s,=! Fig. 2. Banded austenile/ferdlc microsteucltw¢ developed in CrMnN-slcel after 16r/¢ ¢otd work.

Volume loss and relative abrasion-corrosion wear ram (RWR) measurements for the test alloys and mild steel are presented in Table 3. It is obvious from Table 3 that when a corrosion component is added, the wear behavio~r is changed Table 2

Dry abrasion results for the test materials and mild steel AIIo~/

Hardness ( HV,,,k~, )

Volume loss per rl~lr¢

RAR

L'rMnN~;tcel H;Kltield ~,1¢cl A I S I 3(H M i l d ~;~eci

321 .~2X 2d0 I~X)

(}.86 l).68 1.14 1.32

1.53 |.(M. |.15 I.(~)

1

50O volume loss A Hardness

E 0.95

Fig. 3. Brigh|-ficld TEM micrograph hldicating the occurrence of planar slip in CrMnN-stt:el alter 16<;~cold work.

tributed. When CrMnN-steel is cold-worked, the ,5--fertile becomes elongated parallel to the rolling direction to produce a banded structure of alternating regions of austenite and &ferrite. The microstructure alter 16% cold work is exhibited in Fig. 2. Hadfield steel and AISI 31)4 both possessaustenitic

", ==

g 0.9 = o, 0.85

" •

I A i

450 •

A

>

40o~

•

"6 0.8 A

0.75

300 16 30 48 66 % Cold Deformalion Flu. 4. Abra.~ion wear rate and hardness as a fnnclkm o f prior cldd work fnr CrMnN-.~teel. 0

D.J, MilL~. R.D. Km~r~'e.qi W,,ar 215 (199,'q~,~J'-q0

86

Table 4 (.'avitalioneros.ionresi..,tanceof the ile.,¢,lItlateriab

Table 3 Volume Io~,,,,and RWR value~,after abrasiem~,~rrosionte.,,fing Alloy

CrMnN-~,ted Hadlield ,.,1¢cI AISI 3(14 Mild ~teel

12

Tolal ",'tdume loss per c y c l e I ahFdsi~l, li + ¢(z,n'(i, sillll ) I i)1115~

RWR

11.,.12 2.(12 1,15 2.8q"

3.14 1.43 2.73

Alh)y

CrMnN-sled Hadlield ,,.levi AISI 3(14

Incul~alkql

period

Steady-slate t:rosilm rate

( I1 }

t 11191~ h

4.8 2.2 4.3

n,05 1). I1)

I)

O.I)U

1.0(I

Table 5 Inthtencetff prior cold w,wk,m the cavilalitmem~,ilmrc,,b,tance of ('rMnN~,tecl

. . . . .

~o ~, ~ a ~ s =

• al~,,sio,n/eorrosionI

on~

E ~8

Pri~,rcold w.ttrk

os

82

V ,dk,

,dk.

IF i Ifl~.~ 31Y,;

.,lit,

lncl=batitm period

lh)

!

4.8 4,~

Sleud~-.qalc erosion ralc (rmllth ut

5,1)

11.05 11.1)3 11.()2

4S~i

5.g

11.112

h6+~

h.O

0.11[

Meters=~'. (m) I Hours Corr. (hrs) Fi;_ 5_ Malerial ht,.sdue to abrasi~anand corm.,,ion1br Hadlield .,,h:el.

A

12

~0

.A abrasion only

v abrasion/corrosion I

E

o, 6

0

~ra2

~

316o

41i38

Metem Abr. (m) I Hours Corr, (hrs) Fig. h. Material I,~ ~. due t~, ahra~,i*m a n d ct,rn,,,ion for ('rMnN-~,lct:].

in favour of the more c~wrosion-resislant alloys. An indication of the relative comi'a,nents of abrasion and corrosion fi)r Hadlield steel and C r M n N - s t e e l are presented in Figs. 5 :rod 6. respectively. 3.4. C , l | ' i l a l i o n e / ~ s i o J l

When the cumulative volume loss due Io cavitation erosion is plotted as a function of cavitation unze. a distinctive curve emerges. The curve can be divided inlo two stages: the lirst stage is the incubation period during which there is very little material loss. and the second stage represents a steady-state erosion condition, where there is a c~,nstant rate of nlalerial remt,~'al from the surface of the er~Kled specimen. The stead),state erosion rate is obtained from the slope of the curve in the steady-state erosion regime, whereas the incubation time (/.,) is given by an exlralm)lation of the linear steady-state curve OlUO the tinl~ axis. The incubation period and the sleady-slate erosion 1ate lbr each test material are presented in Table 4, The superior cavitation erosion resistance of the C r M n N - s t e e l alloy is dt:monslrated by tvath the longest incubation time and the towest steady-state erosion rate.

Fig. 7. ('rMnN-~teel sur1:aceafter cavitalion exposure for I(I mln. Grain bmmdarics;=redelinealed by pla~li¢th,',v ( aniicak:dcondilk,,l 1. The ¢ffe¢! of prior cold work on the cavitation erosion rcsi.,,tance of CrMnN-steei was also considered, mid the results ;ire summarised it] Table 5. Unlike the i n f l u e n c e o f prior cold work on tile dry abrasion resistance of C r M n N .steel. tile cavitati,sn erosion rate decreases with increasing cold v..'ork. In order It5 investigate tile mechanism of material removal during cavitation erosion of the CrMnN-steel alloy, the erodeu surlaces were examined in the SEM alter lixed inter',';=Is. The iqvestigatioq was carried out for CrMnN-.',teel in both the annealed and cold-worked conditions. Bearing in mind that the initial specimen surface was polished, the development t51 surfi=ce topography during erosion gives an indication of the mode of deformation and material removal. Fig, 7 indicates the surface of annealed CrMnN-steel alter only I(i rain ot'erosion, The surlilce is generally featureless except for the emergence t)f fain!ly delined grain boundaries. In Fig.

OJ. Mill.~. R.D. Kmasea / Wear 215 t 1098) ~3-90

87

¢

Fig. g. CrMnN-steel surface after | h exposure to cavitation. Evidem:e is provided fi~rmalerial exlrusitm at the grain boundaries (annealed cnndition ).

Fig. 10. Slriatkm.~ indieale maleriat removal by a fatigue l~tt~er~.

~.

Fig. % CrMnN-xtcel surface after 5 h expi~sure In cavitation. Extensive material removed ha~ (~'curred {annealed condition 1.

8. which demonstrates the surface appearance alter I h of erosion, the grain boundaries are much more clearly delincd. and in some cases minor extrusion of material has occurred al the grain boundaries. The surface after 5 h of erosion is exhibited in Fig. 9 and shows extensive material removal that has mostly occurred in the groin boundary regions. A slightly higher magnilication view in Fig. I0 of the eroded sadace alter 3 h gives an indication of the mode of material removal. A flake of material has been removed after crack initiation at the grain boundary, and several striations are visible in the area from which the flake was removed. Similar mierographs were obtained Ik~rCrMnN-steel eroded in the prior cold-work conditkm and the initial delineation of grain tmundaries and material removal was observed to ~x:cur in the same way. In view of the strong preferred ~,rientation of the ferrite stringers after cold working, the in fluence of ferrite on material fracture and removal was easier to identify during cavitation erosion of the cold-worked specimens. A distinct preference was noted for material removal at fcrrite-auslenite boundaries as opposed to austenitc/austenite boundaries. Fig. I!. which

Fig. I I. Material loss after 2 h exposure to cavkalkm is ~t,~.,~,t~i',~cdwith a t'errile stringer ( 66t~ " pti,r timid-work condition ).

repre~nts the surface of a specimen after 2 h exposure to erosion, shows material loss associated with a ferrite stringer. Again. a flake of material has been removed, and striations are visible on the underlying surface. Fig. 12 ( 3 h of erosion) shows the partial removal of a flake of material, and the propagation of a sub-surface crack can he envisaged, which leads Io the eventual removal of the flake.

4. Di~ussion 4.1. Micro.so-ttctto'e All three test materials, namely CrMnN-steel. AISI 304 and Hadfield steel have es~ntially austenitic micro.~:~uctures. Approximately 5 vol.% residual ~-ferrite occurs in CrMnNsteel, which is randomly distributed. During cold rolling, the residual &fenite becomes stringered parallel to the rolling direction, and imparts a banded effect in the microstructure. Although the banded structure is expected to pr~xluce anisotropic mechanical properties, the anisolropy is not likely to

88

D.J. Mills, R.D. gmttse, / Wear 215 ( 1998t ,~3-~)

Fig. 12. lXataial removal of a flake of material after 3 h of cavitation (66~d prior cold-~,'ork condition I.

influence wear behaviour, Examination of the dislocation structure of cold-worked material indicates delbrmation by planar slip, which produces a triangular dislocation cell structure (Fig. 3). The sharpness of the planarity of glide is reflected by the remarkable regularity of the spacing between the criss-crossing active planes. This mode of dislocation movement gives rise to high work-hardening rates and is related to the influence of nitrogen on slacking fault energy 16 I. Since high nitrogen levels have a marked effect on lowering stacking fault energy, the partial disloeations are widely ~pa,-ated. This means that a glide plane of given length will accept fewer dislocations, and the amount of strain accommodated per glide plane is reduced. Hence, more glide planes have to be activated, and the required activation stress lbr more dislocation sources has an influence on the work-hardening behaviour. Work-hardening is further influenced by the dislocation barriers, which are created by the intersecting active planes, and the widely spaced partial dislocations demand high stresses to pass these barriers. 4.2. Dry ahntshm

Hadlield steel performs approximately 2(Y;~ better than CrMnN-stee[ in dry abntsion conditions. The mtxte of nmtetrial removal, however, seems to be the same fiw both materials, and is manifested by the formation of plastically deformed lips. and the separation of work-hardened particles in the form of microchips. Ball 171 has demonstrated that there is no simple and direct relationship between wear resislance of a material and mechanical properties such as hardness. yield stress, ultimate tensile strength, toughness and ehmgation, hut that work-hardening bchaviour is important in determining wear resistance. A material of moderate yield strength, but high work-hardeni ng capacity is able to respond plastically to the impacz loading associated with abrasion. This plasticity dissipates energy and obviates the possibility of brittle fracture and spalling. Thus. the indications are thai Hadlield steel has a work-hardening capacity slightly more

favourable for abrasion resistance than CrMnN-steel. A comparison of the work-hardening capacity of these two alloys was made possible by performing compression tests under identical conditions up to a maximum equivalent tensile strain of 0.30. The :,. srk-hardening exponent (n) was calculated over the strain interval 0.12-0.26 and was determined as 0.38 and 0.47 for CrMnN-steel and Hadfleld steel, respectively. Evaluation of the influence of prior cold work on the abrasion resistance of CrMnN-steel demonstrates an increase in wear rate with increasing cold work. Although the surface hardness of the material increase:; as a function of cold work, the ability to accommodate plastic strain during abrasion decreases, and the material is removed more easily. The increase in yield strength due to prior cold work is insufficient to have an influence on the threshold impact of the abrasive strike: consequently, the loss in work-hardenability reduces the resistance to material removal. This behaviour is in contrast to what happens during cavitation erosion in the present investigation, and is also contrary to the common belief thai wear resistance increases with increasing hardness. Since wear rate in this case is determined by the rate of volume toss, the wear is ultimately attributed to micro-fracture that causes the material to separate from the surface, in considering energy-to-fracture, a higher energy-to-fracture would be recorded tot a material that is initially in the annealed condition as opposed to a previously worked material, whose energy-to-fracture has now been lowered by an amount equivalent to the prior work-hardening proeess. Thus, in a situation where the stress level of the abrasive action is considcrably higher than the yield stress of the nmterial, then resistance to yield is not of prime importance, but rather the energy consumed in causing micro-fracture to oecur becomes the dominant factor. If resistance to yield were a determining factor, then it would be expected that the wear resistance should increase with increasin,~ hardness and hence increasing prior cold work. However. the results indicate that this is certainly not the case. Other factors such as changing ferrite nlorphology due to rolling deformation may contribute to material remowfl, but this situation is not clear. The important observation is that the dry abrasion resistance does not increase with increasing prior cold work (and concomitant increase in hardness ). but rather the indications are that there is in fitct a small steady increase in the rate of material relllOVa~.

4.3. Abra.~ioll-com'oshm

The introduction of a corrosion component to the abrasion test has highlighted the signilicance of corrosion resistance in evahmting wear resistant materials for applications that operate in tlostile elwironments. Despite the excellent abrasion resistance of Hadlield steel, the poor corrosion resistance results in it umch inferior wear perl~rmance compared to CrMnN-steel. The superimposition of the abrasion loss and abrasion/corrosion loss curves for CrMnN-steel indicate no loss as a result of corrosion when exposed to the synthetic

D,J, Mills. R.D. Knut.~en / Wear 215 ( l~Jgl¢) &J-90

mine water environment (Fig. 6). On the other hand, the corrosion loss suffered by Hadfield steel during the .same cycle is very significant, as demonstrated by the separation of the abrasion loss and abrasion/corrosion loss curves in Fig. 5. there is very little difference in the inlluence of the corrosion environment on the material loss suffered by AISI 304 and CrMnN-steel. and the latter alloy performs better overall due to its superior abrasion resistance. 4.4. C a v i t a t i o n e r o s i o n

In a study carried out on the cavitation erosion of stainless steel, Heathcock et al. [ 81 concluded thai the erosion resistance of austenitic stainless and mangane~ steel is primarily dependent on their work-hardening rate. Since it is well known that the work-hardening rate of Hadfield steel is higher than AIS1304 stainless steel, the cavitation erosion resistance of Hadlield steel should be significantly different to AlSl 304. However, Table 4 indicates similar steady-state erosion rates for these two alloys, but the possibility of some material loss due to corrosion during testing has been neglected. Unfortunately, the test method does not allow corrosion effects to be determined. The consideration of the role of corrosion becomes more serious when comparing the erosion resistance of Hadfield steel and CrMnN-steel. It is suggested that Hadfield steel has a higher work-hardening rate than CrMnN-steel in view of its superior dry abrasion resistance and higher measured work-hardening exponent, but Table 4 illustrates a much better cavitation erosion resistance for CrMnN-steel. The incubation period is also decidedly longer for CrMnN-steel. The result further emphasises the fact that some material loss may be suffered by Hadlicld steel due to corrosion in the distilled water medium. Since corrosion conditions are always likely to be prevalent in cavitation situations, the ranking produced by the laboratory tests is significant. A similar influence of corrosion was determined when comparing the cavitation erosion behaviour of Hadfield steel, and a manganese steel containing 13"~ chromium. where it is reported that the Hadfield steel suffered significant material loss due to corrosion {9 lInvestigation of the mechanism of material removal during the cavitation erosion of CrMnN-sleel indicates that as cavitation commences, the stress is accommodated by plastic deformation. Slip is discontinuous at the grain boundaries and results in 'raised" areas occurring at the grain boundaries ( Fig. 8 ). Eventually, the level of extrusion at the grain boundaries progresses to the extent where material is removed Fig. 9). The striations exhibited in Fig. 10 suggest material removal by a fatigue process that can be understood from the following explanation. The imploding bubbles result in stress waves, with most of the stress being directed parallel to the surface. The raised grain boundary areas are subjected to a greater stress than the fiat surface, and cracks are initiated and propagated by cyclic implosion action. Sub-surface propagation of cracks more or less parallel to the specimen surface results in eventual flaking of the material, L'Esperance et al.

89

[ I01 have demonstrated that the fatigue crack initially propagates away from the original surface, but curves back to rejoin this surface, resulting in the removal of a chip of metal. The observations in the pre~nt study demonstrate that there are three main events in the initial stages of the cavitation erosion process, namely ( i ) plastic deformation of individual grains, which is discontinuous at the grain boundaries, i it) formation of 'rated" areas as material is extruded at the grain boundaries, and ( iii ) fatigue crock initiation and growth from the extruded regions. Examination in the SEM also reveals that the pre~nce of fi-ferrite adversely affects the erosion resistance, becau~ material is more likely to he lost from 8ferrite grains than from austenite grains during the early part of cavitation erosion, Less strain can be accommodated by the ~-ferrite due to its much lower work-hardening capacity. with the result that a more brittle failure mode occurs in this region. Thus. a further improvement on the CrMnN-steel alloy could be obtained by modifying the composition slightly to decrea~ Ibe &ferrite content, The improved cavitation erosion resistance of CrMnNsteel specimens that were previously cold-worked indicates that resistance to plastic flow is important during cavitation erosion. This is manifested by both an increase in incubation time and a decrease in steady-state erosion rate a.s the amount of prior cold work increases. Cold work raises the yield stress that affects all three stages of the cavitation erosion process previously mentioned. Since the higher yield stress resiststhe amount of plastic deformation in the individual grains, the onset of grain boundary extrusion is delayed. Furthermore, fatigue crack initiation and growth in the extruded regions is made more dit~cult by the higher flow stress. Experiments that were carried out to determine the cavitation behaviour of metaslable Cr-Mn austcnitic steel demonstrate that prior cold work reduces the erosion resistai,ce due to the formation of martensite during the pre-detbrmation process 19 l- The preexisting martensite accelerates further martensite formation during cavitation: as a result, the volume fraction of retained austenite is too low tbr a relaxation of high local stres~s. In the case of CrMnN-steel. the austenite is found to be stable. and martensite does not form during either cold working or during actual cavitation erosion cycles ! I I ]. In view of the stability having been determined by X-ray diffraction, where the resolution is within 2-3 vol.~, very small amounts of martensite formation on interacting glide planes cannot he discounted. However. this would not afi~.ct the ability of the austenite to accommodate high stresses. Thus, the indications are that the pre-deformed austenite in CrMnN-steel can still accommodate the high stresses by further pl:tstic detbrmation. as well as resisting fatigue crack growth.

$. Conclusions The evaluation of the tribological behaviour of CrMnNsteel in comparison to AISI 304 stainless steel and Hadfield steel has demonstrated the following findings: ( a ) CrMnN-

~)

D.d. Mills. R.D. Knulsen / Wear 215 ( 199818.,1-90

steel performs better than AISI 304 during dry abrasion, but not as good as Hadfield steel; (b) when a corrosion environment is considered in the tribological system, CrMnN-steel out performs Hadfield steel; (c) CrMnN-steel demonstrates superior cavitation erosion resistance in distilled water:. (d) the cavitation erosion resistance of CrMnN-sleel improves with increasing prior cold work.

Acknowledgements The authors gratefully acknowledge the linancial and technical support provided by Columbus Stainless ( Middleburg, South Africa) and the financial support provided by the Foundation for Research Development ( FRD, Pretoria). The assistance of the technical and secretarial staff in the Department of Materials Engineering. UCT, is greatly appreciated.

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Biographies Robert Knutsen is an Associate Professor in the Materials Engineering Department at the University of Cape Town. He obtained his PhD degree in Materials Engineering at UCT in 1989 where he specialised in the physical metallurgy of duplex ferrite-martensite stainless steels, Hiscurrent research interests include stainless steel metallurgy and thermomechanical processing of sheet metals, Duncan Mills obtained a Masters degree in Materials Science at the University of Cape Town in 1995. This paper is based on his Masters Thesis.