tribological behaviour of shot peened cu-ni austempered ductile iron

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Tribological behaviour of shot peened Cu–Niaustempered ductile iron

ARTICLE in WEAR · APRIL 2013Impact Factor: 1.91 · DOI: 10.1016/j.wear.2012.12.027

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5 AUTH O RS , INCLUDING:

Ann Zammit

University of Malta

2 PUBLICATIONS 7 CITATIONS

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Stephen C Abela

University of Malta


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Lothar Wagner

Technische Universität Clausthal


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Mansour Mhaede

Technische Universität Clausthal


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Available from: Mansour MhaedeRetrieved on: 19 December 2015

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initiation and propagation which in turn prolongs the lifetime of components. However, as the spherical shots hit the surface,dimples are formed which roughen the surface. The surface nishof mating surfaces affects the wear resistance of components. Thiscombination of a high friction coefcient at the beginning of thetest and a work hardened surface was also observed by Ohba et al.[12] when studying rolling contact fatigue properties of ADI.Similar ndings were also reported by Ho et al. [13] who showed

that shot peening did not improve the sliding wear resistance of annealed 1018 steel, but it did decrease the wear rate of hardened4340 steel. On the other hand, work by Vaxevanidis et al. [14]showed that shot peening had a benecial effect on the tribolo-gical behaviour of steel.

A large number of studies have shown that shot peeningimproves the bending fatigue strength [15 –19 ]. However, veryfew works have been conducted on the tribological behaviourafter shot peening, and it is thus not clear whether it actuallyimproves the wear resistance. This paper compares the slidingwear behaviour of ground ADI with shot peened ADI specimenstested under dry conditions.

2. Experimental procedure

2.1. Material and processing

Test pins of 5 mm diameter used for the pin-on-disk weartesting were machined from ductile iron keel blocks, having thecomposition shown in Table 1 . After machining, samples wereaustenitised at 900 7 2 1 C, and then rapidly quenched in a saltbath at 360 7 5 1 C and held for 1 h. The samples were then aircooled to room temperature. These austempering parameterswere optimised in a previous study [20] . Samples were coatedusing a dedicated paint (SEMCO Zir H) to prevent decarburisationduring the austempering process.

Test disks of 90 mm diameter were made out of D2 tool steelof chemical composition shown in Table 2 and heat treated to ahardness of 61 HRC. The heat treatment cycle consisted of pre-heating, followed by austenitising at a temperature of 1025 1 Cand then quenching using nitrogen at a pressure of 5 bar. Afterthat, the disks were tempered at a temperature of 190 1 C for 3 h.

2.2. Surface treatment

After heat treatment of both pins and disks, the surfaces wereground up to a mean surface roughness Ra of 0.2 mm. Aftergrinding, half of the pins were shot peened. Shot peening wascarried out using S330 shots, with an Almen intensity of 0.38 mmA up to 100% coverage. The stand-off distance was90 mm while the angle of impingement was set at 90 1 . The

surface roughness Ra of the shot peened pins was measured tobe 3.7 mm.

2.3. Characterisation

The microhardness measurements were taken using aMitutoyo MVK-H12 microhardness tester. Three measurementsfor the hardness were taken. The coefcient of variation (the ratio

of standard deviation to the mean) of the measurements was inall cases below 5%.

Phase analysis before and after shot peening was carried outusing the X-ray diffraction method and a Bruker D8 AdvanceX-Ray diffractometer (Mo- K a radiation). The scanning step was0.01 1 , the dwell time 0.2 s and 2 y values between 15 and 40 1 . Thetube acceleration voltage and current used were 45 kV and35 mA, respectively. The XRD patterns obtained were subjected

to the Savitzky–Golay smoothing lter which performs a localpolynomial regression to the raw data reducing the signal noise[21] . A 3rd order regression was found to preserve features of thepattern including relative maxima and width of peaks. Theretained austenite content ( gret ) in the ADI was measured withthe X-ray diffractometer using the simplied method describedby Miller [22] .

2.4. Test equipment and conditions

Dry sliding wear tests were carried out using a conventionalpin-on-disk tribometer capable of maintaining a constant unidir-ectional, sliding velocity between the pin and disk. The machine

used was an Italdesign TR-20 which allowed control of the load,velocity, duration of test and radius at which the pin acts on thedisk. In this study, a cantilever loaded at ended cylindrical ADIpin was made to slide over a rotating hardened steel disk. The pinwas xed to one end of the cantilever arm, and the other end wasattached to displacement and force transducers. The tribometerwas connected to a computer which monitored and recorded thedisplacement of the pin and disk and the frictional force.

Tests followed ASTM G99-05 ( Standard test method for wear testing with a pin-on-disk apparatus ) procedures and were carriedout in ambient air held at room temperature. Tests wereperformed at two values of pressure acting between the cylind-rical surface of the pin and the horizontal rotating disk namely at2.5 and 10 MPa, while the sliding distance, sliding velocity and

radius at which the pin acted upon the disk were kept constant at3.6 km, 4 m/s and 26 mm, respectively.Before and after tribological tests, both the pins and the disks

were cleaned for 10 min in an acetone ultrasonic bath, and thenrinsed in isopropanol and dried in a jet of hot air. For eachexperiment, a new pin and new counter body were used. Themass of both the pins and the disks was measured before andafter each test using a digital balance having an accuracy of 7 0.1 mg. The mass lost was converted into wear volume W ,taking the density r of the ADI material as 6890 kg/m 3 . The wear factor K was then calculated using the relation K ¼ W/F s, where W represents the wear volume in mm 3 , F is the applied load in N ands is the sliding distance in m [23] . Two surface conditions of ADIwere tested namely: ground ADI (G) and shot peened ADI (SP)(Table 3 ). At least ve sliding tests were carried out for eachcondition and applied load, and the average data is reported.

Table 1Chemical composition of the keel blocks.

Element C Si Cu Ni Mn P Mg Al S Fe

Composition (wt %) 3.26 2.36 1.63 1.58 0.24 0.011 0.057 0.024 0.006 Bal

Table 2Chemical composition of the steel reference disks.

Element Cr C Mo V Mn Si Fe

Composition (wt %) 11.8 1.55 0.8 0.8 0.4 0.3 Bal

A. Zammit et al. / Wear 302 (2013) 829 –836 830

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The tests were carried out at an ambient temperature of 18 7 3 1 Cand relative humidity of 42 7 5%.

2.5. Post-wear characterisation

The worn surfaces of the pins were examined using opticalmicroscopy and scanning electron microscopy. Changes in thesubsurface region were also investigated by using metallographictechniques and by taking microhardness measurements on thecross-section wear test samples.

After the tests, the debris was collected in order to investigatethe products generated during sliding of the specimens. SEM andEDX were used to obtain information regarding the morphologyand chemical composition of the wear debris.

3. Results and discussion

3.1. ADI microstructure

The microstructure of the austempered material ( Fig. 1) con-

tains 200 graphite nodules per mm 2 with graphite nodularitygreater than 95% in a matrix consisting of acicular ferrite andcarbon enriched austenite. The macrohardness of the ADI struc-ture is 336 7 15 HV, while the microhardness of the ausferriticstructure is 370 7 10 HV.

3.2. Characterisation of shot peened specimens

Shot peening of austempered ductile iron results in a strain-induced phase transformation. In fact, the face centred cubic (FCC)retained austenite present in the ausferrite matrix transforms to bodycentred tetragonal (BCT) martensite by cold-working caused by shotsimpinging on the surface of the component. It is evident from themicrograph in Fig. 2 that the microstructure at the surface is distinct

from that of the base matrix.

Optical micrograph observations are supported by the X-raydiffraction patterns shown in Fig. 3. The amount of retained austenitein as-austempered polished specimens was calculated to be about44%. In comparison, none of the austenite peaks ( g111 , g200 , g220 , g311 )were observed in the shot peened ADI specimen, but only ferrite andmartensite peaks. This indicates that the austenite has transformed tomartensite as a result of shot peening.

The work hardening induced by shot peening and the phasetransformation to martensite results in an increase of the surfacehardness from 370 to 535 HV. The microhardness depth-proleshown in Fig. 4 also indicates that the hardness falls continuouslyas the distance from the surface increases.

The residual compressive stress of shot peened specimens wasmeasured in a previous study [15] . The maximum occurred at thesurface and was about 975 MPa, 67% higher than the yield point of the material.

3.3. Study of the surfaces of wear test samples

In tribological tests, the wear process changes as the proper-ties of the surface alter. At a low applied pressure of 2.5 MPa, the

microhardness of the worn surface is about 19% higher than that

Table 3Tribological test conditions.

Specimenidentication

Surfacecondition

Nominal appliedpressure (MPa)

G2.5 Ground ADI 2.5SP2.5 Shot peened ADIG10 Ground ADI 10SP10 Shot peened ADI

Fig. 1. Austempered ductile iron microstructure.

Fig. 2. SP microstructure.

Fig. 3. X-ray diffraction patterns of as-treated and shot peened ADI.

A. Zammit et al. / Wear 302 (2013) 829–836 831

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of the bulk; 440 and 370 HV, respectively ( Fig. 5) for tested Gspecimens. The thickness of this hardened layer is around100 mm. This was also observed by Refaey et al. and Islam et al.[24 ,25 ] who state that this increase in hardness is due to strainhardening of the ausferritic matrix at the surface region whichpredominates over any frictional heating effect. As a result of thisplastic deformation, the material is stronger and causes surfaceow and the microstructure to distort. Fig. 6 shows the distortedmicrostructure of a section taken across the worn surface aftertesting with the lower load.

It is noted that the hardness of SP specimens decreases afterwear testing, from 535 to 450 HV. This is probably due to theremoval of part of the shot peened layer during the wear test, ortempering of the martensite which was formed during the shotpeening process ( Fig. 3).

On the other hand, the microhardness at the surface of speci-mens tested at the higher load is over 600 HV ( Fig. 5). Thisindicates a phase transformation to a high hardness phase.Micrographs show that a white non-etchable phase is present atthe surface of the specimens tested with the higher load ( Fig. 7).When the two surfaces slide over each other, most of the work

done against friction is converted into heat, causing a general rise

in temperature, as well as localised temperature spikes where anasperity makes contact with the mating surface. The resulting risein temperature may modify the mechanical and metallurgicalproperties of the sliding surfaces, causing them to oxidise, orpossibly melt. This high temperature transforms the ausferrite toaustenite and can result in carbon diffusion into the austenite,making it stable and hence increasing the hardenability of the pin.As a consequence, the critical cooling rate is lowered, resulting inthe formation of untempered martensite at a slow cooling rateupon cooling of the pin and disk after the test is stopped. Anotherplausible reason of martensite formation is that the austenitebeing produced due to the high temperatures being formed at theasperities is rapidly cooled due to heat being conducted into theunderlying bulk material.

This was also observed by Fordyce et al. [26] who reported awhite non-etching layer formed during the unlubricated slidingwear of austempered spheroidal cast iron. However, these whitelayers were not found on the worn surfaces of the as-castspheroidal iron. Straffelini et al. [4] explained how the wear rateof ADI at high sliding speeds (1.5–2.6 m/s) was dominated by theformation and cracking of this white layer formed on the sliding

surface. Sharma [27] has also shown that high loads applied

Fig. 4. Microhardness prole of specimens before wear tests.

Fig. 5. Microhardness proles of a cross section of the worn surface.

Fig. 6. Microstructure just below the worn surface shows distortion and surfaceow.

Fig. 7. Martensite formed on worn surfaces, tested at an applied pressure of 10 MPa.


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alloying or shot peening, increased surface hardness resulting inlower frictional resistance [29] . He attributed this to the fact thathardening and shot peening introduce residual stresses and phasetransformations which lower the strength of the cold-welded

junctions. This, he maintains, lowers adhesion of the surfaces and

hence lowers the frictional resistance [30] . This contrasts with theresults shown in the present work. One may argue that thenegative impact of the rougher surface, was in this study moredominant than the benecial inuence of any phase transforma-tion ( Figs. 2 and 3 ), and compressive residual stresses present atthe surface of SP specimens [15] . Analysis of the surface rough-ness of the shot peened wear test samples used in the two studiesmay throw more light on this apparent inconsistency.

Fig. 14 also shows that a lower CoF was recorded when testingspecimens at a higher applied load. High applied pressurepromotes the phase transformation of retained austenite tomartensite, a harder and load bearing phase. It is known thatfriction properties are generally improved when the hardness of the surface is increased. As explained by Mokhtar et al. [31] , coldweld junctions formed when hard phases like martensite arepresent are relatively easy to break, hence lowering the adhesionof the surfaces and the frictional resistance. Also, martensite hasbetter thermal properties, providing better heat dissipation lead-ing to a reduction in the CoF [32 ,33 ].

On the other hand, under the action of low applied loads, nomartensite formation is observed. The asperities of the harderdisk indent into the softer ausferritic structure of the pin causingplastic deformation and strong cold-welded junctions are formed.Frictional sliding resistance to motion is thus higher due to thelarger force required to shear these welded junctions [31] .

3.6. Wear mechanism

Different wear mechanisms are believed to have occurredduring the tests. As observed by Straffelini et al. [4] , the wearfactors shown in Fig. 9 are typical of mild oxidational wear [7 ,34 ].During dry sliding, the surfaces interact with the atmosphereresulting in a mild form of corrosive wear, known as oxidationalwear, which primarily occurs during unlubricated conditions of sliding [23] . Due to the high frictional heating during sliding, fastoxidation occurs. The oxide layers formed usually appear asislands on the sliding surfaces ( Fig. 15 ). The separation of thesurfaces due to these layers results in mild wear [35] .

Debris is formed due to the fatigue of the oxide lm producedand the generation of frictional heat which raises the temperatureof the sliding surfaces. This weakens the bonding strengthbetween the oxide lm and the substrate, resulting in delamina-tion of the oxide layer. The worn surfaces and the wear debris

collected after wear tests had a red–brown tinge and EDX analysis

Fig. 11. Pitting of worn SP surfaces tested at an applied pressure of 10 MPa.

Fig. 12. Crack propagation through graphite nodules (G10).

Fig. 13. Crack arrested by the graphite nodule.

Fig. 14. Friction coefcient evolution.


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identied large amounts of oxygen ( Fig. 16 ). This conrms theoxidational wear as the main wear mechanism which occurredduring the tests. The other identied peaks in Fig. 16 shows thepresence of iron, carbon, silicon and copper, all of which areelements present in the material being studied.

SEM observation of wear debris generated by sliding showedthat small particles were agglomerated to combine into largerwear particles ( Fig. 17 ). This was also observed by Stachowiak

where the wear particles were seen to agglomerate into clustersduring the unlubricated sliding wear of steels and cast iron [36] .

Fig. 18 shows the wear scar of a pin where sliding marksparallel to the direction of sliding were observed. The ne groovesshow that several plateaus are formed at the beginning of thewear process when the contacting surfaces achieved conformity[23] . During sliding, these plateaus become unstable, they breakup to form debris, and other wear grooves tracks are formed. Thegeneration of wear particles can change the wear mechanism tothree-body abrasive wear, which leads to microploughing of the

surface. The dimpled surface of the SP specimens can trap thesewear particles, thus increasing the wear rate [37] .

4. Conclusions

In this study, unlubricated sliding wear tests were carried outto determine the effect of shot peening (SP) on the tribologicalbehaviour of Cu–Ni alloyed austempered ductile iron. Pin-on-disktests were conducted on ground ADI, and shot peened ADIspecimens using two nominal applied loads. The results of thepresent work are summarised as follows:

(1) Shot peening of ADI results in an increase in surface rough-ness and hardness, together with austenite to martensitetransformation.

(2) Specimens tested at the lower load showed a distortedmicrostructure just below the wear scar indicative of surfaceow. An increase in surface hardness of 19% was noted onthese surfaces after wear tests. On the other hand, untem-pered martensite was observed on the surface of specimensFig. 15. Oxide layer on worn surface of G2.5 pin.

Fig. 16. EDX analysis of the wear debris.

Fig. 17. SEM micrographs of debris collected from wear tests; (a) ADI 10 MPa, (b) SP 10 MPa.

Fig. 18. Ploughing marks on the pin surface.

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tested at the higher load. This was attributed either to thehigh ash temperatures and/or to the high stresses reachedduring sliding.

(3) The main wear mechanism was mild oxidative wear wherea reddish-brown layer was seen on the worn surfaces. Inaddition, material loss was due to the result of three-bodyabrasive wear.

(4) No improvement was noticed on the wear factor after shot

peening, despite the increase in surface microhardness, theintroduction of residual compressive stresses, work hardeningof the surface and stress-induced austenite to martensitetransformation. The roughened surface induced by shot peening,and possibly the entrapment of wear particles between thedimpled surface of the SP specimens and the disk may have hadan inuence on the wear rate. In addition, no differences werenoticed on the microstructures just below the worn surfaces,debris analysis and friction coefcients of ground and shotpeened samples.

(5) This work shows that shot peening does not result in animprovement in the sliding wear resistance. This indicatesthat shot peening can be applied to components whichrequire a higher bending fatigue resistance, without loweringsignicantly the wear resistance.

Acknowledgements

The authors would like to thank Mr. Mark Joseph Zerafa(B.Eng.(Hons.)) for his contribution in machining and testing of specimens.

In addition, the authors would like to acknowledge thepositive impact of ERDF funding and the purchase of the testingequipment through the project: ‘‘ Developing an InterdisciplinaryMaterial Testing and Rapid Prototyping R &D Facility (Ref. no. 012)’’.

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tribological behaviour of shot peened cu-ni austempered ductile iron

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