**3. Shot peening of ADI**

The high strains applied by the high-pressure impact of the shots on the surface of ADI are higher than the yield strength of the material, and both the ferrite and the austenite undergo plastic deformation. The ferrite work hardens with an increase in the dislocation density, while the austenite has the ability to cold work and to locally transform into martensite at high plastic deformation. The transformation induced plasticity (TRIP) phenomenon has been studied by various researchers [12, 13] and shown to depend on the carbon content in the austenite, its size and morphology and distribution within the structure.

As a result of work hardening and phase transformation, the hardness at the surface following SP of ADI increases by approximately 40–60% [11, 14–19]. Apart from an increase in the surface hardness, SP also results in the creation of compressive residual stresses of around 700–1000 MPa [15–19]. The residual compressive stresses caused by SP increases the dislocation density hindering dislocation motion. The stress-induced austenite to martensite transformation also results in volume expansion further creating local compressive stresses.

### **3.1. Influence of shot peening on bending fatigue strength of ADI**

Shot peening is known to improve the bending fatigue properties of ADI [14, 15, 20, 21]. The improvement of bending fatigue resistance following SP is due to the formation of a compressed layer beneath the surface of the SP component. Cracks tend not to initiate or propagate in surfaces upon which a compressive force is acting. The induced compressive stresses shift crack nucleation to the sub-surface and hinders fatigue crack propagation at the surface. Apart from hardening the surface, SP also eliminates microscopic defects, machining marks and grinding defects. This also increases the bending fatigue strength. However, Uematsu et al. [22] showed that SP cast iron containing spheroidal vanadium carbides (VCs), dispersed in a martensitic matrix, did not eliminate large casting voids. These voids and clusters of VCs served as sources of crack initiation, and hence SP did not improve the bending fatigue strength in this particular case.

## **3.2. Influence of shot peening on tribological characteristics of ADI**

The induced compressive layer and work hardening increase the resistance to crack initiation and propagation, which in turn prolongs components' lifetime. In fact, SP has been used for years to extend the bending fatigue life of various engineering components in the transport industry, mainly for automobile and aircraft parts. Such components include gears, axles, springs, connecting rods, crankshafts, I-beams in heavy duty applications, compressors, turbine rotors and shafts. SP can be applied to the whole component, or just confined to parts of the component that are expected to be highly stressed. For example, in gears, one can SP the entire gear tooth or alternatively focus the SP only at the tooth root fillets, which are exposed

Although SP is not a new process, it is still very popular and has evolved considerably in the late twentieth and early twenty-first centuries. A great deal of research has been carried out to study the effects of SP on the materials being treated, including metallurgical, mechanical, and geometrical effects and also the effects of SP on the mechanical properties, tribological characteristics and corrosion resistance of a wide range of metals and alloys [11]. Recently, advanced surface modification technologies show the use of variants of the traditional SP process, four of which are shown in this issue, namely ultrasonic SP, severe SP (SSP), surface mechanical attrition treatment (SMAT) and duplex SP. Other variants of SP can be found in literature, including laser shock peening (LSP), micro-peening, cavitation shotless peening

The high strains applied by the high-pressure impact of the shots on the surface of ADI are higher than the yield strength of the material, and both the ferrite and the austenite undergo plastic deformation. The ferrite work hardens with an increase in the dislocation density, while the austenite has the ability to cold work and to locally transform into martensite at high plastic deformation. The transformation induced plasticity (TRIP) phenomenon has been studied by various researchers [12, 13] and shown to depend on the carbon content in

As a result of work hardening and phase transformation, the hardness at the surface following SP of ADI increases by approximately 40–60% [11, 14–19]. Apart from an increase in the surface hardness, SP also results in the creation of compressive residual stresses of around 700–1000 MPa [15–19]. The residual compressive stresses caused by SP increases the dislocation density hindering dislocation motion. The stress-induced austenite to martensite transformation also results in volume expansion further creating local com-

Shot peening is known to improve the bending fatigue properties of ADI [14, 15, 20, 21]. The improvement of bending fatigue resistance following SP is due to the formation of a compressed layer beneath the surface of the SP component. Cracks tend not to initiate or

the austenite, its size and morphology and distribution within the structure.

**3.1. Influence of shot peening on bending fatigue strength of ADI**

to the highest stress.

28 Advanced Surface Engineering Research

and ball or roller burnishing.

**3. Shot peening of ADI**

pressive stresses.

Wear of materials is a complex phenomenon, and depends on the running conditions and the properties of the tribopair materials. SP should be beneficial in reducing wear rates because of the high hardness due to work hardening, stress-induced austenite to martensite transformation, and residual compressive stresses at the surface [10, 23, 24]. Compressive stresses prevent micro-cracks from forming, thus inhibiting pitting or spalling. Kobayashi and Hasegawa [25] showed this to be true for carburised steel gears. This was attributed to the compressive stress suppressing cracking and delaying crack growth. Champaigne [10] and Townsend and Zaretsky [26] reported an improvement in the contact fatigue life of SP steel gears of around 1.5 times. In another study, Townsend [27] reported that a higher peening intensity (0.38–0.43 mmA) resulted in higher compressive stresses, and hence lead to a 10% rolling contact fatigue life (*L10*) of 2.15 times that of the carburised gears peened with an intensity of 0.18–0.23 mmA. Adamović et al. [28] investigated the sliding wear characteristics of ground and SP steel under boundary lubrication, and reported a slight decrease in coefficient of friction and a 30% increase in the wear resistance after SP.

The effect of the rough SP surfaces on the tribological behaviour has been reported in a number of articles. The dimpled surface is sometimes considered to be favourable in lubricated contact, which acts as reservoirs that aid in the retention of the lubricant and maintaining a full film thickness between meshing teeth. Better lubrication reduces fretting, noise, spalling, scuffing and the operating temperature by reducing friction. That said, Vaxevanidis et al. [29] still reported improved sliding wear resistance of SP tool steel tested under dry conditions. A higher coefficient of friction was reported at the beginning of the test, but as the test progressed, this decreased to a lower value than that of specimens, which were not SP. This can be attributed to flattening of the rough surfaces during the wear test.

In contrast, studies by other researchers show no improvement in the tribological characteristics of surfaces after SP [30–32]. The reason given is that surface roughening counterbalances the positive effects of compressive stresses and hardening caused by SP.

Very few works have been conducted to study the tribological behaviour of SP ADI [30, 31]. Work by Sharma [31] showed that at a given load, the contact fatigue life of SP Mo-Ni ADI austempered at 230°C is 35–45% lower than that of carburised steel. The author attributed this to the rougher surface and a lower hardness of the SP ADI when compared to carburised steel. Lubricated rolling contact fatigue tests carried out by Ohba et al. [30] showed that SP Cu-alloyed ADI exhibited similar wear rates to corresponding as-austempered specimens. This was attributed to surface roughening, which counterbalances the positive effects of compressive stresses and hardening caused by SP.

## **3.3. Case study: shot peening of Cu-Ni-alloyed ADI**

This section is related to a study, which was conducted to address the inconsistencies related to the tribological characteristics of SP ADI. A Cu-Ni-alloyed ADI having the composition shown in **Table 1** was used for this study [16–19]. Ductile iron samples were first austenitised at a temperature of 900°C for 2 hours and subsequently austempered at 360°C for 1 hour. Following the austempering process, this material had an upper ausferrite matrix with an yield strength of 737 MPa, a tensile strength of 1012 MPa and an elongation of 7%. SP was done up to full coverage with S330 steel shots and with an Almen intensity of 0.38 mmA. The stand-off distance was 90 mm, while the angle of impingement was set at 90°. The surface roughness *Ra* following SP was measured to be 3.1 μm.

As indicated in Section 2, SP of austempered ductile iron results in strain-induced phase transformation from the face-centred cubic (FCC) austenite to body centred tetragonal (BCT) martensite. This can be observed in the X-ray diffraction pattern presented in **Figure 3**. The top pattern, which is for the SP specimen, does not show any of the austenite peaks present in the as-austempered ductile iron (bottom pattern in **Figure 3**), suggesting that only ferrite and martensite peaks are present and that the strain induced during SP has transformed the austenite to martensite. Similar findings are commonly reported in ferrous alloys having retained austenite present in the initial micro-structure before SP [14, 15, 23, 33].

The SP-induced work hardening and phase transformation result in an increase of the surface hardness of 43% from 370 to 535 HV (**Figure 4(a)**) and decreases steadily towards the interior of the specimen. This is in agreement with results reported in literature, where the typical hardness increase following SP of ADI is in the range of 40–60% [14, 15]. The depth of the SP layer was measured from the hardness-depth profile (**Figure 4(a)**) and is approximately 400 μm. The maximum compressive stress occurring at the SP surface has a value of 975 MPa (**Figure 4(b)**), which is 67% greater than the yield strength of the material. Similar values of compressive stress for the Cu-Mn ductile iron austempered at 380°C were reported by Ebenau et al. [15].

**Figure 3.** X-ray diffraction patterns of polished as-treated and shot-peened ADI (S330, intensity = 0.38 mmA) [18].

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 31

**Figure 4.** (a) Micro-hardness-depth profile of shot-peened ADI (S330, intensity = 0.38 mmA) and (b) residual stress-depth

profile of shot-peened ADI (S330, intensity = 0.38 mmA) [18].

#### *3.3.1. Bending fatigue resistance of shot-peened Cu-Ni ADI*

As a result of SP, both the mean bending fatigue strength and the fatigue life of the Cu-Ni ADI were increased [18]. **Figure 5** shows the S-N curve for both the as-austempered and SP condition. It can be noted that the fatigue strength increased by approximately 60%, that is from 250 to 390 MPa. The improvement in fatigue life of around 35% was noted at all stress levels. On


**Table 1.** Chemical composition of the Cu-Ni-alloyed ductile iron.

This was attributed to surface roughening, which counterbalances the positive effects of com-

This section is related to a study, which was conducted to address the inconsistencies related to the tribological characteristics of SP ADI. A Cu-Ni-alloyed ADI having the composition shown in **Table 1** was used for this study [16–19]. Ductile iron samples were first austenitised at a temperature of 900°C for 2 hours and subsequently austempered at 360°C for 1 hour. Following the austempering process, this material had an upper ausferrite matrix with an yield strength of 737 MPa, a tensile strength of 1012 MPa and an elongation of 7%. SP was done up to full coverage with S330 steel shots and with an Almen intensity of 0.38 mmA. The stand-off distance was 90 mm, while the angle of impingement was set at 90°. The surface

As indicated in Section 2, SP of austempered ductile iron results in strain-induced phase transformation from the face-centred cubic (FCC) austenite to body centred tetragonal (BCT) martensite. This can be observed in the X-ray diffraction pattern presented in **Figure 3**. The top pattern, which is for the SP specimen, does not show any of the austenite peaks present in the as-austempered ductile iron (bottom pattern in **Figure 3**), suggesting that only ferrite and martensite peaks are present and that the strain induced during SP has transformed the austenite to martensite. Similar findings are commonly reported in ferrous alloys having retained

The SP-induced work hardening and phase transformation result in an increase of the surface hardness of 43% from 370 to 535 HV (**Figure 4(a)**) and decreases steadily towards the interior of the specimen. This is in agreement with results reported in literature, where the typical hardness increase following SP of ADI is in the range of 40–60% [14, 15]. The depth of the SP layer was measured from the hardness-depth profile (**Figure 4(a)**) and is approximately 400 μm. The maximum compressive stress occurring at the SP surface has a value of 975 MPa (**Figure 4(b)**), which is 67% greater than the yield strength of the material. Similar values of compressive stress for the Cu-Mn ductile iron austempered at 380°C were reported by Ebenau et al. [15].

As a result of SP, both the mean bending fatigue strength and the fatigue life of the Cu-Ni ADI were increased [18]. **Figure 5** shows the S-N curve for both the as-austempered and SP condition. It can be noted that the fatigue strength increased by approximately 60%, that is from 250 to 390 MPa. The improvement in fatigue life of around 35% was noted at all stress levels. On

**Element C Si Cu Ni Mn P Mg Al S Fe** Wt.% 3.26 2.36 1.63 1.58 0.24 0.011 0.057 0.024 0.006 Bal.

pressive stresses and hardening caused by SP.

30 Advanced Surface Engineering Research

roughness *Ra*

**3.3. Case study: shot peening of Cu-Ni-alloyed ADI**

following SP was measured to be 3.1 μm.

austenite present in the initial micro-structure before SP [14, 15, 23, 33].

*3.3.1. Bending fatigue resistance of shot-peened Cu-Ni ADI*

**Table 1.** Chemical composition of the Cu-Ni-alloyed ductile iron.

**Figure 3.** X-ray diffraction patterns of polished as-treated and shot-peened ADI (S330, intensity = 0.38 mmA) [18].

**Figure 4.** (a) Micro-hardness-depth profile of shot-peened ADI (S330, intensity = 0.38 mmA) and (b) residual stress-depth profile of shot-peened ADI (S330, intensity = 0.38 mmA) [18].

**Figure 5.** S-N curves for polished as-austempered ductile iron and SP ADI specimens [18].

the one hand, the results obtained for SP ADI (390 MPa) are similar to those obtained in other studies conducted by Mhaede et al. [34] and Ochi et al. [35] for unalloyed ADI, and by Benam et al. [14] for Cu-Ni-alloyed ADI. On the other hand, however, Ebenau et al. [15] reported values of 560 MPa when carrying out bending fatigue tests on SP Cu-Mn ADI austempered at 380°C. Surprisingly, these high values were obtained using the same S330 shots used in the present investigation: similar peening pressure of 3 bar and with recorded maximum residual compressive stress of 700 MPa. These results could potentially be explained in terms of differences in chemical compositions or heterogeneity in micro-structures.

surfaces. When traversing between graphite nodules, cracks were seen to pass through the lath of austenite and ferrite or through the austenite/ferrite interface, these being paths of least resistance. Similar results were also reported by Voigt [38] and Tayanç et al. [39]. Therefore, if nodules are sometimes considered as defects, the high toughness of stable austenite makes up for this negative effect. Crack propagation is also affected by the high toughness and ductility of the ausferritic micro-structure, making it a strong crack arrester. This structure absorbs the energy of the advancing crack during fracture, arresting the crack or deflecting it. In addition, ausferrite can strain harden during cyclic loading, providing high plastic deformation, which further hinders crack growth. Tanaka et al. [40] measured the hardness of upper ausferritic matrix after being subjected to different number of cycles during bending fatigue tests. The hardness of the matrix increased with longer number of cycles. Apart from strengthening by plastic deformation, unstable austenite can transform to martensite as the crack advances. The accompanying volume change, resulting in a compressed zone ahead of the crack tip, may

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 33

**Figure 6.** Crack propagation along graphite-matrix interface for a shot-peened ADI specimen [18].

The extent to which the nodules can affect the fatigue behaviour of ADI depends on the nodularity where a high nodularity provides a better continuity of the matrix with less stress raisers. Furthermore, the nodule size, count and distribution are of equal importance. A high count, a smaller diameter and evenly distributed nodules decrease intercellular micro-segregation of elements. This results in an increase in the fracture toughness and fatigue resistance of ADI.

Dry sliding wear tests were carried out on the Cu-Ni ADI using a conventional pin-on-disk tribometer using two different applied pressures, 2.5 and 10 MPa [16]. **Figure 8** shows the wear factor *K* of the as-austempered DI and SP ADI specimens as a function of sliding distance. It can be noted that there is no difference between the wear rates for the two surface

retard crack growth, if not arrest it completely.

*3.3.2. Dry sliding wear resistance of shot-peened Cu-Ni ADI*

The fatigue ratio, which is the ratio of the fatigue limit to the tensile strength, is 0.31 and 0.39 for the polished and SP specimens, respectively. This ratio is in agreement with British Cast Iron Research Association (BCIRA) reports that quote a fatigue ratio of 0.37 for austempered irons with tensile strengths in the range of 900–1000 N/mm<sup>2</sup> [36]. Johansson et al. [37] obtained higher endurance ratios of around 0.44, which were obtained in irons with tensile strengths in the range of 1000–1200 N/mm<sup>2</sup> produced by austempering between 350 and 375°C.

When comparing the bending fatigue strength obtained for the ADI in the present study to that for the carburised steel, it can be noted that the performance of carburised steel is far superior. Bending fatigue strengths of between 850 and 1500 MPa have been reported in a number of studies. This can be attributed to the harder surfaces (approximately 700 HV) and the deeper carburised layers (approximately 1.2 mm).

The crack propagation deflected along the graphite-matrix interface is shown in a cross section of a fractured specimen in **Figure 6**, which indicates a relatively weak interface between graphite nodules and the matrix. **Figure 7** also shows nodule debonding from one of the fractured

**Figure 6.** Crack propagation along graphite-matrix interface for a shot-peened ADI specimen [18].

surfaces. When traversing between graphite nodules, cracks were seen to pass through the lath of austenite and ferrite or through the austenite/ferrite interface, these being paths of least resistance. Similar results were also reported by Voigt [38] and Tayanç et al. [39]. Therefore, if nodules are sometimes considered as defects, the high toughness of stable austenite makes up for this negative effect. Crack propagation is also affected by the high toughness and ductility of the ausferritic micro-structure, making it a strong crack arrester. This structure absorbs the energy of the advancing crack during fracture, arresting the crack or deflecting it. In addition, ausferrite can strain harden during cyclic loading, providing high plastic deformation, which further hinders crack growth. Tanaka et al. [40] measured the hardness of upper ausferritic matrix after being subjected to different number of cycles during bending fatigue tests. The hardness of the matrix increased with longer number of cycles. Apart from strengthening by plastic deformation, unstable austenite can transform to martensite as the crack advances. The accompanying volume change, resulting in a compressed zone ahead of the crack tip, may retard crack growth, if not arrest it completely.

The extent to which the nodules can affect the fatigue behaviour of ADI depends on the nodularity where a high nodularity provides a better continuity of the matrix with less stress raisers. Furthermore, the nodule size, count and distribution are of equal importance. A high count, a smaller diameter and evenly distributed nodules decrease intercellular micro-segregation of elements. This results in an increase in the fracture toughness and fatigue resistance of ADI.

#### *3.3.2. Dry sliding wear resistance of shot-peened Cu-Ni ADI*

the one hand, the results obtained for SP ADI (390 MPa) are similar to those obtained in other studies conducted by Mhaede et al. [34] and Ochi et al. [35] for unalloyed ADI, and by Benam et al. [14] for Cu-Ni-alloyed ADI. On the other hand, however, Ebenau et al. [15] reported values of 560 MPa when carrying out bending fatigue tests on SP Cu-Mn ADI austempered at 380°C. Surprisingly, these high values were obtained using the same S330 shots used in the present investigation: similar peening pressure of 3 bar and with recorded maximum residual compressive stress of 700 MPa. These results could potentially be explained in terms of differ-

The fatigue ratio, which is the ratio of the fatigue limit to the tensile strength, is 0.31 and 0.39 for the polished and SP specimens, respectively. This ratio is in agreement with British Cast Iron Research Association (BCIRA) reports that quote a fatigue ratio of 0.37 for austempered

higher endurance ratios of around 0.44, which were obtained in irons with tensile strengths in

When comparing the bending fatigue strength obtained for the ADI in the present study to that for the carburised steel, it can be noted that the performance of carburised steel is far superior. Bending fatigue strengths of between 850 and 1500 MPa have been reported in a number of studies. This can be attributed to the harder surfaces (approximately 700 HV) and

The crack propagation deflected along the graphite-matrix interface is shown in a cross section of a fractured specimen in **Figure 6**, which indicates a relatively weak interface between graphite nodules and the matrix. **Figure 7** also shows nodule debonding from one of the fractured

produced by austempering between 350 and 375°C.

[36]. Johansson et al. [37] obtained

ences in chemical compositions or heterogeneity in micro-structures.

**Figure 5.** S-N curves for polished as-austempered ductile iron and SP ADI specimens [18].

irons with tensile strengths in the range of 900–1000 N/mm<sup>2</sup>

the deeper carburised layers (approximately 1.2 mm).

the range of 1000–1200 N/mm<sup>2</sup>

32 Advanced Surface Engineering Research

Dry sliding wear tests were carried out on the Cu-Ni ADI using a conventional pin-on-disk tribometer using two different applied pressures, 2.5 and 10 MPa [16]. **Figure 8** shows the wear factor *K* of the as-austempered DI and SP ADI specimens as a function of sliding distance. It can be noted that there is no difference between the wear rates for the two surface

The micro-hardness-depth profiles taken on cross sections of the worn specimens are shown in **Figure 9(a)**, while **Figure 9(b)** compares the surface hardness at both applied pressures of specimens before and after the wear tests. At a low applied pressure of 2.5 MPa, the microhardness of the worn surface of as-austempered specimens was measured to be around 19% higher than that of the bulk. The thickness of this hardened layer is around 100 μm and is due to strain hardening of the ausferritic matrix at the surface region, which predominates over any frictional heating effect. As a result of this plastic deformation, the material is stronger

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 35

It can be noted that the hardness of SP specimens decreases from 535 to 450 HV (**Figure 9**) after testing at the lower applied pressure of 2.5 MPa. This is probably due to the removal of part of the SP layer during the wear test, or tempering of the martensite, which was formed

On the other hand, the surface micro-hardness of specimens tested at the higher load is over 600 HV (**Figure 9**). This indicates a phase transformation to a high hardness phase. Micrographs show that a white non-etchable phase is present at the surface of the specimens tested with the higher load (**Figure 10(b)**). 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 an asperity makes contact with the mating surface. The resulting rise in temperature may modify the mechanical and metallurgical properties of the sliding surfaces, causing them to oxidise, or possibly melt. This high temperature transforms the ausferrite to austenite and can result in carbon diffusion from the nodules into the austenite and hence increasing the hardenability of the pin. Consequently, the critical cooling rate is lowered, resulting in the formation of untempered martensite at a slow cooling rate upon cooling of the pin and disk after the test is stopped. It is also possible that the austenite being produced due to the high temperatures reached at the asperities is rapidly cooled as heat is conducted into the underlying bulk material when the tip of the asperity breaks during

**Figure 9.** (a) Micro-hardness-depth profiles of cross sections of worn specimens and (b) surface hardness before and after

and causes surface flow and the micro-structure to distort, as shown in **Figure 10(a)**.

sliding. As a result, the austenite transforms to martensite during testing.

during the SP process.

dry sliding wear tests [16].

**Figure 7.** Graphite nodules attached to one of the fractured surfaces of an as-austempered DI specimen [18].

**Figure 8.** Wear factors of as-austempered ductile iron (A) and shot-peened ADI (SP) samples [16].

conditions: as-austempered DI and SP ADI specimens at the two applied loads. This indicates that SP did not result in an improvement in the dry sliding wear resistance. This suggests that a higher original hardness (535 HV for SP specimens compared to 370 HV for as-austempered specimens) does not necessarily result in better wear resistance. It may be argued, however, that the potential improvement resulting from an increase in hardness is being counterbalanced by the increased surface roughness caused by SP.

The micro-hardness-depth profiles taken on cross sections of the worn specimens are shown in **Figure 9(a)**, while **Figure 9(b)** compares the surface hardness at both applied pressures of specimens before and after the wear tests. At a low applied pressure of 2.5 MPa, the microhardness of the worn surface of as-austempered specimens was measured to be around 19% higher than that of the bulk. The thickness of this hardened layer is around 100 μm and is due to strain hardening of the ausferritic matrix at the surface region, which predominates over any frictional heating effect. As a result of this plastic deformation, the material is stronger and causes surface flow and the micro-structure to distort, as shown in **Figure 10(a)**.

It can be noted that the hardness of SP specimens decreases from 535 to 450 HV (**Figure 9**) after testing at the lower applied pressure of 2.5 MPa. This is probably due to the removal of part of the SP layer during the wear test, or tempering of the martensite, which was formed during the SP process.

On the other hand, the surface micro-hardness of specimens tested at the higher load is over 600 HV (**Figure 9**). This indicates a phase transformation to a high hardness phase. Micrographs show that a white non-etchable phase is present at the surface of the specimens tested with the higher load (**Figure 10(b)**). 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 an asperity makes contact with the mating surface. The resulting rise in temperature may modify the mechanical and metallurgical properties of the sliding surfaces, causing them to oxidise, or possibly melt. This high temperature transforms the ausferrite to austenite and can result in carbon diffusion from the nodules into the austenite and hence increasing the hardenability of the pin. Consequently, the critical cooling rate is lowered, resulting in the formation of untempered martensite at a slow cooling rate upon cooling of the pin and disk after the test is stopped. It is also possible that the austenite being produced due to the high temperatures reached at the asperities is rapidly cooled as heat is conducted into the underlying bulk material when the tip of the asperity breaks during sliding. As a result, the austenite transforms to martensite during testing.

conditions: as-austempered DI and SP ADI specimens at the two applied loads. This indicates that SP did not result in an improvement in the dry sliding wear resistance. This suggests that a higher original hardness (535 HV for SP specimens compared to 370 HV for as-austempered specimens) does not necessarily result in better wear resistance. It may be argued, however, that the potential improvement resulting from an increase in hardness is being counterbal-

**Figure 8.** Wear factors of as-austempered ductile iron (A) and shot-peened ADI (SP) samples [16].

**Figure 7.** Graphite nodules attached to one of the fractured surfaces of an as-austempered DI specimen [18].

34 Advanced Surface Engineering Research

anced by the increased surface roughness caused by SP.

**Figure 9.** (a) Micro-hardness-depth profiles of cross sections of worn specimens and (b) surface hardness before and after dry sliding wear tests [16].

**Figure 10.** Micro-structure just below the worn surface of a specimen tested at an applied pressure of (a) 2.5 and (b) 10 MPa [16].

Fordyce et al. [41] also observed this white non-etchable layer during the unlubricated sliding wear of austempered spheroidal cast iron but not of as-cast spheroidal iron. Straffelini et al. [42] explained how the wear rate of ADI at high sliding speeds of 1.5–2.6 m/s was dominated by the formation and cracking of this white layer formed on the sliding surface. Sharma [31] has also shown that high loads applied during wear testing may transform the metastable austenite to martensite.

SP specimens. On the other hand, the larger number of smaller asperities for the ground asaustempered specimens leads to a larger real area of contact upon application of the normal force and the presence of a very thin oil film. As a result, the number of sliding cycles to failure decreases due to plastic flow of the softer ADI specimens and micro-fracture of the asperities. This highlights the importance of the surface topography in asperity-asperity contact during

**Figure 11.** (a) Number of sliding cycles to scuffing failure for as-austempered (A) ADI and shot-peened (SP) ADI specimens and (b) friction coefficient evolution for lubricated sliding wear tests involving the as-austempered ADI and

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 37

The higher scuffing resistance of the SP specimens could also be attributed to the high compressive stresses present in these specimens. It has been explained in previous sections that compressive stresses create a resistance to crack propagation and flaking of the surface. This was also attested by Adamović et al. [28], who reported a 30% improvement in the wear

Additionally, the SP specimens exhibit a lower value of coefficient friction at 0.08, when compared to that measured for their counterpart as-austempered specimens. This can be seen from the friction coefficient data of the lubricated sliding wear tests presented in **Figure 11(b)**. The coefficient of friction for the as-austempered specimens is seen to increase progressively as the test progresses. This is attributable to the larger number of small asperities of the asaustempered pin in contact with the disk. In contrast, the local traction at asperity contacts is reduced in the dimpled SP specimens, again as a result of the oil pockets on the surface. It is known that at lower levels of friction, the surface traction forces and sub-surface shear forces between two interacting bodies might not be sufficient to initiate crack growth and delamination. Hence, the lower friction for the SP specimens leads to longer number of cycles before the onset of failure. A lower coefficient of friction than that produced by ground surfaces was also reported for SP steel specimens [28] and for dented steel surfaces [44] under starved

The higher hardness of the SP specimens (~535 HV) should have also contributed to the improved scuffing resistance. The benefit of a high hardness in wear tests carried out under

sliding of components.

shot-peened ADI specimens [43].

resistance of steels after SP.

lubrication conditions.

The presence of graphite nodules has a major influence on the wear rate of ADI as superficial graphite is smeared over the surface and aids in lubricating the surfaces in sliding contact. Graphite naturally has an inherent lubricating ability, being able to smear over the contacting surfaces. This lowers the friction coefficient and prevents metallic contact, hence reducing the adhesive bonding between the surfaces and material loss due to wear. Due to this, ADI components, for example railcar wheels, are sometimes run dry without the need of lubrication. Cracks have a tendency of passing through the graphite-matrix interface, this being the path of least resistance. On the other hand, a nodule may arrest crack propagation. Whether or not a crack is arrested or assisted to propagate as it reaches a graphite nodule would depend on the angle of approach. It follows that graphite nodules influence the propagation path.

#### *3.3.3. Scuffing resistance of shot-peened Cu-Ni ADI*

**Figure 11(a)** shows that during starved lubricated sliding wear tests, the SP Cu-Ni ADI (SP) specimens exhibited a higher scuffing wear resistance than corresponding as-austempered (A) specimens [43]. SP specimens survived 21 × 10<sup>3</sup> cycles, while the as-austempered specimens endured 2.3 × 10<sup>3</sup> cycles before failure. The improved scuffing performance due to SP might seem anomalous, since rough surfaces generally create low values of the specific film thickness λ and induce scuffing. However, the superposition of indentations arising from the SP process can be considered as an advantage in starved lubricated moving parts. These act as oil reservoirs by dragging the oil into them and generating a load-carrying hydrodynamic pressure. This decreases the pressure from the sliding surfaces, leading to longer lives for the

**Figure 11.** (a) Number of sliding cycles to scuffing failure for as-austempered (A) ADI and shot-peened (SP) ADI specimens and (b) friction coefficient evolution for lubricated sliding wear tests involving the as-austempered ADI and shot-peened ADI specimens [43].

Fordyce et al. [41] also observed this white non-etchable layer during the unlubricated sliding wear of austempered spheroidal cast iron but not of as-cast spheroidal iron. Straffelini et al. [42] explained how the wear rate of ADI at high sliding speeds of 1.5–2.6 m/s was dominated by the formation and cracking of this white layer formed on the sliding surface. Sharma [31] has also shown that high loads applied during wear testing may transform the metastable

**Figure 10.** Micro-structure just below the worn surface of a specimen tested at an applied pressure of (a) 2.5 and (b)

The presence of graphite nodules has a major influence on the wear rate of ADI as superficial graphite is smeared over the surface and aids in lubricating the surfaces in sliding contact. Graphite naturally has an inherent lubricating ability, being able to smear over the contacting surfaces. This lowers the friction coefficient and prevents metallic contact, hence reducing the adhesive bonding between the surfaces and material loss due to wear. Due to this, ADI components, for example railcar wheels, are sometimes run dry without the need of lubrication. Cracks have a tendency of passing through the graphite-matrix interface, this being the path of least resistance. On the other hand, a nodule may arrest crack propagation. Whether or not a crack is arrested or assisted to propagate as it reaches a graphite nodule would depend on the angle of approach. It follows that graphite nodules influence the propagation path.

**Figure 11(a)** shows that during starved lubricated sliding wear tests, the SP Cu-Ni ADI (SP) specimens exhibited a higher scuffing wear resistance than corresponding as-austempered (A) specimens [43]. SP specimens survived 21 × 10<sup>3</sup> cycles, while the as-austempered specimens endured 2.3 × 10<sup>3</sup> cycles before failure. The improved scuffing performance due to SP might seem anomalous, since rough surfaces generally create low values of the specific film thickness λ and induce scuffing. However, the superposition of indentations arising from the SP process can be considered as an advantage in starved lubricated moving parts. These act as oil reservoirs by dragging the oil into them and generating a load-carrying hydrodynamic pressure. This decreases the pressure from the sliding surfaces, leading to longer lives for the

austenite to martensite.

36 Advanced Surface Engineering Research

10 MPa [16].

*3.3.3. Scuffing resistance of shot-peened Cu-Ni ADI*

SP specimens. On the other hand, the larger number of smaller asperities for the ground asaustempered specimens leads to a larger real area of contact upon application of the normal force and the presence of a very thin oil film. As a result, the number of sliding cycles to failure decreases due to plastic flow of the softer ADI specimens and micro-fracture of the asperities. This highlights the importance of the surface topography in asperity-asperity contact during sliding of components.

The higher scuffing resistance of the SP specimens could also be attributed to the high compressive stresses present in these specimens. It has been explained in previous sections that compressive stresses create a resistance to crack propagation and flaking of the surface. This was also attested by Adamović et al. [28], who reported a 30% improvement in the wear resistance of steels after SP.

Additionally, the SP specimens exhibit a lower value of coefficient friction at 0.08, when compared to that measured for their counterpart as-austempered specimens. This can be seen from the friction coefficient data of the lubricated sliding wear tests presented in **Figure 11(b)**. The coefficient of friction for the as-austempered specimens is seen to increase progressively as the test progresses. This is attributable to the larger number of small asperities of the asaustempered pin in contact with the disk. In contrast, the local traction at asperity contacts is reduced in the dimpled SP specimens, again as a result of the oil pockets on the surface. It is known that at lower levels of friction, the surface traction forces and sub-surface shear forces between two interacting bodies might not be sufficient to initiate crack growth and delamination. Hence, the lower friction for the SP specimens leads to longer number of cycles before the onset of failure. A lower coefficient of friction than that produced by ground surfaces was also reported for SP steel specimens [28] and for dented steel surfaces [44] under starved lubrication conditions.

The higher hardness of the SP specimens (~535 HV) should have also contributed to the improved scuffing resistance. The benefit of a high hardness in wear tests carried out under starved lubrication conditions was also mentioned by Adamović et al. [28]. In fact, the hardness of mating surfaces is a crucial factor in avoiding scuffing.

It was also noted that the graphite nodules play a part in determining the scuffing resistance. In fact, as seen in **Figure 12** cracks emanate or stop at the nodules. This observation is similar to that mentioned earlier during bending fatigue [18] and dry sliding wear of ADI [16].

## *3.3.4. Rolling contact fatigue resistance of shot-peened Cu-Ni ADI*

As-austempered and SP Cu-Ni ADI specimens were tested up to pitting failure using a conethree ball tribosystem at an applied stress of 2.56 GPa [19]. **Figure 13** presents the Weibull probability plot for the tests, showing the percentage of specimens that will fail up to a specific number of cycles. The cumulative distribution function (CDF) is shown on the y-axis, while the number of rolling cycles to contact fatigue is shown on the x-axis. In this plot, the data points represent the number of cycles to failure of the ADI specimens.

**Figure 13(b)** shows the average contact fatigue lives for the as-austempered DI and SP specimens. Results show that the average contact fatigue life of SP specimens decreased by 72% when compared to the performance of the as-austempered specimens. Similarly, Sharma et al. [31] reported that SP lowered the contact fatigue life of ADI by 60%. Also, Vrbka et al. [46] report an 82% decrease in life of steel specimens following rolling contact fatigue tests. In the current study, austempering at 360°C resulted in a hardness of 370 HV, while SP increased the surface hardness to approximately 530 HV. Based only on hardness, one would expect an improvement in the contact fatigue resistance. Also, the residual compressive stresses present in SP layers (**Figure 4(b)**) should effectively reduce the maximum shear stress inside the Hertzian contact field, delaying crack nucleation and propagation.

value of the specific film thickness *λ* of 0.05, which contrast with a corresponding value of 2.81 for the as-austempered specimens. A value of *λ* smaller than 0.4 denotes that a boundary lubricated condition exists, implying that peaks of the asperities of the SP specimens are penetrating the lubricant film. Contact, therefore, occurs between the asperities of the SP ADI cones and balls. Metal-to-metal contact is unavoidable and nearly all the load is supported by the asperities, defying the principal purpose of lubrication. Contact fatigue failure starts from

**Figure 13.** (a) Weibull probability plot for data obtained from rolling contact fatigue tests and (b) L10, L50, characteristic

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 39

This might seem anomalous after having shown that under starved lubricated sliding wear, the dimples forming the rough surface of the SP specimens act as lubricant reservoirs, aiding in keeping the lubricant between the surfaces and serving to delay failure. One may have expected this positive attribute to apply also for the rolling tests. However, with the application of larger quantity of lubricant, smoother surfaces of the as-austempered samples proved to be more beneficial as it resulted in the formation of a full lubricating film. One should note that in contrast, all the specimens in the starved lubricated sliding wear tests were tested under boundary lubricated conditions. These results are in agreement with Zhai et al. [44] who report that the effect of surface dents is favourable under poorly lubricated conditions, but adverse under a well-lubricated environment. This means that the influence of dimpled surfaces on the tribological characteristics of the material depends on the lubrication regime. For rough SP surfaces to operate in a fully lubricated condition, the minimum film thickness should be thicker than the combined surface roughness of the two interacting components. A thicker film is obtained by increasing the operating speed, the oil viscosity or the relative radius of curvature of the surfaces in contact. The elastic properties of the gear teeth and the applied load have relatively small influences on the lubricant film thickness. Increasing the load will only increase the elastic flattening (the width of the Hertzian contact band) and the contact area, without changing the geometry of the inlet region. Apart from the benefits of a thicker film, improving the performance of SP surfaces might be due to lower surface roughness. This can be achieved by either using shots having a smaller diameter, or by grinding/polishing the surface

the surface as a result of a presumably high coefficient of friction.

(η) fatigue lives for as-austempered and shot-peened ADI specimens [19, 45].

However, the surfaces of the SP specimens have a higher surface roughness (*Ra* = 3.1 μm) than the counterpart as-austempered specimens (*Ra* = 0.4 μm). This leads to an extremely low

**Figure 12.** Micro-graph showing the influence of graphite nodules on crack initiation and/or propagation [43].

starved lubrication conditions was also mentioned by Adamović et al. [28]. In fact, the hard-

It was also noted that the graphite nodules play a part in determining the scuffing resistance. In fact, as seen in **Figure 12** cracks emanate or stop at the nodules. This observation is similar to that mentioned earlier during bending fatigue [18] and dry sliding wear of ADI [16].

As-austempered and SP Cu-Ni ADI specimens were tested up to pitting failure using a conethree ball tribosystem at an applied stress of 2.56 GPa [19]. **Figure 13** presents the Weibull probability plot for the tests, showing the percentage of specimens that will fail up to a specific number of cycles. The cumulative distribution function (CDF) is shown on the y-axis, while the number of rolling cycles to contact fatigue is shown on the x-axis. In this plot, the

**Figure 13(b)** shows the average contact fatigue lives for the as-austempered DI and SP specimens. Results show that the average contact fatigue life of SP specimens decreased by 72% when compared to the performance of the as-austempered specimens. Similarly, Sharma et al. [31] reported that SP lowered the contact fatigue life of ADI by 60%. Also, Vrbka et al. [46] report an 82% decrease in life of steel specimens following rolling contact fatigue tests. In the current study, austempering at 360°C resulted in a hardness of 370 HV, while SP increased the surface hardness to approximately 530 HV. Based only on hardness, one would expect an improvement in the contact fatigue resistance. Also, the residual compressive stresses present in SP layers (**Figure 4(b)**) should effectively reduce the maximum shear stress inside the

However, the surfaces of the SP specimens have a higher surface roughness (*Ra* = 3.1 μm) than the counterpart as-austempered specimens (*Ra* = 0.4 μm). This leads to an extremely low

**Figure 12.** Micro-graph showing the influence of graphite nodules on crack initiation and/or propagation [43].

ness of mating surfaces is a crucial factor in avoiding scuffing.

38 Advanced Surface Engineering Research

*3.3.4. Rolling contact fatigue resistance of shot-peened Cu-Ni ADI*

data points represent the number of cycles to failure of the ADI specimens.

Hertzian contact field, delaying crack nucleation and propagation.

**Figure 13.** (a) Weibull probability plot for data obtained from rolling contact fatigue tests and (b) L10, L50, characteristic (η) fatigue lives for as-austempered and shot-peened ADI specimens [19, 45].

value of the specific film thickness *λ* of 0.05, which contrast with a corresponding value of 2.81 for the as-austempered specimens. A value of *λ* smaller than 0.4 denotes that a boundary lubricated condition exists, implying that peaks of the asperities of the SP specimens are penetrating the lubricant film. Contact, therefore, occurs between the asperities of the SP ADI cones and balls. Metal-to-metal contact is unavoidable and nearly all the load is supported by the asperities, defying the principal purpose of lubrication. Contact fatigue failure starts from the surface as a result of a presumably high coefficient of friction.

This might seem anomalous after having shown that under starved lubricated sliding wear, the dimples forming the rough surface of the SP specimens act as lubricant reservoirs, aiding in keeping the lubricant between the surfaces and serving to delay failure. One may have expected this positive attribute to apply also for the rolling tests. However, with the application of larger quantity of lubricant, smoother surfaces of the as-austempered samples proved to be more beneficial as it resulted in the formation of a full lubricating film. One should note that in contrast, all the specimens in the starved lubricated sliding wear tests were tested under boundary lubricated conditions. These results are in agreement with Zhai et al. [44] who report that the effect of surface dents is favourable under poorly lubricated conditions, but adverse under a well-lubricated environment. This means that the influence of dimpled surfaces on the tribological characteristics of the material depends on the lubrication regime.

For rough SP surfaces to operate in a fully lubricated condition, the minimum film thickness should be thicker than the combined surface roughness of the two interacting components. A thicker film is obtained by increasing the operating speed, the oil viscosity or the relative radius of curvature of the surfaces in contact. The elastic properties of the gear teeth and the applied load have relatively small influences on the lubricant film thickness. Increasing the load will only increase the elastic flattening (the width of the Hertzian contact band) and the contact area, without changing the geometry of the inlet region. Apart from the benefits of a thicker film, improving the performance of SP surfaces might be due to lower surface roughness. This can be achieved by either using shots having a smaller diameter, or by grinding/polishing the surface after SP. In fact, rolling contact fatigue tests carried out by Ohba et al. [30] showed nearly equal fatigue lives for ground as-austempered DI and SP ADI using shots with a diameter of 0.1 mm. On the other hand, Vrbka et al. [46] report a deterioration in the rolling contact fatigue resistance of steel despite the use of shots having smaller diameters of 0.07 and 0.11 mm. One notes that shots used in the present study had diameters in the range of 0.85–1.2 mm. Results in this study [46] were improved when testing specimens, which were polished after SP, thus creating relatively smoother surfaces, in which the asperities did not protrude the lubricant film. However, grinding or polishing of SP surfaces might be challenging since extra care must be taken so as not to remove the SP layer and hence eliminate the beneficial effects, which result from SP (high hardness and compressive stresses at the surface).

**5.** Lubricated rolling contact fatigue tests revealed that SP resulted in a 72% decrease in the average contact fatigue life when compared to the resulting fatigue life obtained by the asaustempered specimens. This was attributed to the rough surfaces of SP specimens, which in turn caused a low specific film thickness, leading to rolling in the boundary lubrication regime. In contrast, rolling of the polished as-austempered specimens was conducted in the presence of a full lubricant film, which is the ideal lubrication regime of components

Shot Peening of Austempered Ductile Iron http://dx.doi.org/10.5772/intechopen.79316 41

The SP process is constantly maturing, and many questions still remain open as the industry is continuously on the search for process improvements that improve and extent the service lifetime of components. For example, the improvement in surface roughness has improved the tribological characteristics, the ability to create textured nanostructured surface layers and

The authors would like to acknowledge the positive impact of ERDF funding and the purchase of the testing equipment through the project: Developing an Interdisciplinary Material

Department of Metallurgy and Materials Engineering, University of Malta, Msida, Malta

[1] Zammit A. Tribological and mechanical characteristics of surface modified austempered

[2] Borui Casting International Ltd. Austempering Process for Ductile Iron [Internet]. Available from: http://www.metals-china.com/austempering-process-for-ductile-iron-

[3] Harding RA. The production, properties and automotive applications of austempered

[4] Shanmugam P, Rao PP, Udupa KR, Venkataraman N. Effect of microstructure on the fatigue strength of an austempered ductile iron. Journal of Materials Science. 1994;

also new equipment and techniques to characterise the treated surfaces [11].

Testing and Rapid Prototyping R&D Facility (Ref. no. 012).

Address all correspondence to: ann.zammit@um.edu.mt

ductile iron [PhD thesis]. University of Malta; 2014

austempered-ductile-iron.html [Accessed: 17-05-2018]

ductile iron. Kovove Materialy. 2007;**45**:1-16

**29**:4933-4940. DOI: 10.1007/BF00356546

under rolling contact.

**Acknowledgements**

**Author details**

Ann Zammit

**References**
