Dry sliding of plasma-sintered iron—The influence of nitriding on wear resistance

Dry sliding of plasma-sintered iron—The influence of nitriding on wear resistance

Available online at www.sciencedirect.com Wear 265 (2008) 301–310 Dry sliding of plasma-sintered iron—The influence of nitriding on wear resistance ...

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Available online at www.sciencedirect.com

Wear 265 (2008) 301–310

Dry sliding of plasma-sintered iron—The influence of nitriding on wear resistance H.C. Pavanati a,∗ , G. Straffelini b , A.M. Maliska a , A.N. Klein a a

LABMAT, Departamento de Engenharia Mecˆanica, Universidade Federal de Santa Catarina - 88040-970, Florian´opolis, SC, Brazil b Dipartimento di Ingegneria dei Materiali, Universit` a di Trento, Via Mesiano 77, 38050 Trento, Italy Received 14 March 2007; received in revised form 28 August 2007; accepted 22 October 2007 Available online 20 February 2008

Abstract Samples of unalloyed iron were sintered in the presence of an abnormal glow discharge using the confined anode–cathode configuration in order to allow surface enrichment of the sample with atoms from the cathode. AISI 430 and AISI 1020 cathodes were used to sinter the iron with and without Cr enrichment, respectively. The chromium favors the formation of a hard compound layer after nitriding (with a microhardness up to 1300 HV) due to the presence of chromium nitrides. In order to evaluate the wear resistance and surface damage to the samples, dry sliding wear tests were carried out using a block-on-ring system with applied normal loads of 25–150 N. A large scatter in the experimental results was observed for the unnitrided samples. This effect was attributed to the capacity of the materials to undergo a severe-to-mild wear transition during sliding. Nitriding was found to reduce the wear rate by at least one order of magnitude. In particular, the nitrided samples enriched with chromium (540 ◦ C) did not display any damage at 25 N, the lowest normal load used in the present investigation. © 2007 Elsevier B.V. All rights reserved. Keywords: Powder metallurgy; Plasma nitriding; Confined anode–cathode configuration; Chromium enrichment; Dry sliding wear resistance

1. Introduction Powder metallurgy is increasingly used in the production of mechanical parts. Sintered parts are now mainly used in applications where static or alternating loads are encountered. In order to favor their use in applications where contact loads are present, a complete understanding of the tribological behavior of these materials is needed. The role of the surface treatments needed to improve their wear resistance also has to be understood, since the residual porosity will noticeably decrease their mechanical strength [1–3]. In the present investigation, the potentiality of the abnormal glow discharge technology as a means to sinter metallic components and, at the same time, to enrich the surface of the parts by Cr for a successive nitriding treatment is investigated. Using this technique the specimens could work as either the cathode or the anode of the discharge [4–6]. In the case of the anode–cathode configuration, simultaneously to the sintering process, it is possi-



Corresponding author. Tel.: +55 48 37219268; fax: +55 48 32340059. E-mail address: [email protected] (H.C. Pavanati).

0043-1648/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2007.10.014

ble to add atoms from the cathode onto the sample surface due to the cathodic sputtering, as has been shown by Pavanati et al. [7]. As Cr has a relatively high affinity with nitrogen, the presence of such an alloy element in solid solution in steels could be of interest for nitriding treatments. An improvement in the mechanical properties after nitriding, using Cr as the alloy element, has been demonstrated [8–10]. A nitriding treatment of Cr-enriched specimens, as well as mechanical and microstructure characterizations had previously been carried out [11]. It was observed that the hardness values and the depth of the hardened layer were increased when enriched with Cr before nitriding. In the present work, dry sliding tests were carried out in order to study the wear behavior of the iron samples sintered in plasma, with and without Cr enrichment, after plasma nitriding treatment. The influence of the nitriding treatment on the wear resistance of the sintered samples was evaluated. 2. Experimental techniques and materials The plasma reactor used to sinter the materials has been described elsewhere [7]. The samples were 12.7 mm × 32.0 mm ×6.0 mm and made from atomized iron powder DC177 (sup-

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Fig. 1. Schematic of the reactor with confined anode–cathode configuration.

plied by H¨ogan¨as do Brasil Ltda.) mixed with 0.6 wt% of zinc stearate. The powder particles were in the range 30–200 ␮m and the mean particle size distribution was around 100 ␮m. After an hour of mixing the powder was compacted to a pressure of 600 MPa using a double action press with a moving die body. After compaction, the green density of the samples was approximately 7.05 × 10−6 g m−3 . The arrangement of the electrodes in the plasma reactor was the confined anode–cathode configuration shown in Fig. 1. The sample holder was 50 mm long and 10 mm in width and the cylindrical outer cathode was 60 mm long, 35 mm in diameter, and 1 mm thick, as shown in Fig. 1. The anode was grounded and the cathode was negatively biased by using a DC pulsed power source with output voltages of 400, 500, 600, and 700 V, a maximum current of 5 A and a pulse period of 10–250 ␮s. The temperature of the sample was a function of the number and the energy of the ions impinging on the cathode in the plasma reactor. Using a pulsed power source, it is possible to adjust the

tension applied to the cathode (which is directly proportional to the energy of the ions) and the pulse length “time switched on” (corresponding to the effective time of bombardment). Raising these parameters would also raise the temperature of the cathode. The pressure of the gas mixture in the reactor chamber is proportional to the number of ions impinging on the cathode. Therefore, the temperature would be raised when the gas pressure is increased. A type K thermocouple protected by an Inconel cover, 1.5 mm in diameter, electrically insulated with Al2 O3 was inserted into the sample holder at a previously known T point between the sample and the holder (Fig. 1) in order to take the temperatures. Delubrication was carried out in the presence of an abnormal glow discharge as described by Santos et al. [12]. A hydrogen discharge was generated at a pressure of 400 Pa using the anode–cathode configuration (Fig. 1) and adjusting the gas flow to 3.33 × 10−6 standard m3 s−1 (200 sccm). The power supply voltage was fixed at 500 V and the temperature was maintained at 350 ◦ C for 30 minutes, controlled by the “time switched on” of the power supply (ton ). A plasma sintering cycle occurred after the delubrication using a gas mixture of 80%Ar + 20%H2 , a gas flow of 4.00 × 10−6 standard m3 s−1 and an output voltage of 500 V. The ton was fixed at 150 ␮s and the sintering temperature was reached by increasing the pressure. The sintering was carried out at 1150 ◦ C for 60 min and the heating rate was about 20 ◦ C min−1 . The pressure of the gas inside the chamber required to reach that temperature, using 500 V was about 1300 Pa. The choice of cathode material depended on whether or not surface enrichment was desired. Therefore, two sets of samples were produced using an AISI 430 and AISI 1020 steel cathode for sintering with and without chromium enrichment, respectively. For comparison, another set of samples were sintered in a resistance furnace, following the same thermal cycle as was used in the plasma reactor. After the sintering cycle, a set of samples was nitrided in an abnormal glow discharge of a gas mixture of 25%H2 + 75%N2 at a pressure of 400 Pa for 120 min at different temperatures. Table 1 shows the variable parameters of the treatments. The dry sliding tests were carried out using an Amsler A135 tribometer test machine with block-on-ring configuration. In this case, the sample was the block and the ring was a disc, 40 mm in diameter and 10 mm in height, produced from AISI 52100 steel, which had been treated to reach a hardness of around 750 HV. Before the test the surface of the block was grinded with a SiC

Table 1 Sintering and nitriding treatment conditions for unalloyed iron Symbol

F-Fe F-FeN540 P-Fe P-FeN540 P-Cr P-CrN540 P-CrN450

Description

Furnace sintered Furnace sintered and plasma nitrided Plasma sintered Plasma sintered and plasma nitrided Plasma sintered with Cr enrichment Plasma sintered with Cr enrichment and plasma nitrided Plasma sintered with Cr enrichment and plasma nitrided

Sintering cycle

Nitriding cycle

Atmosphere

Cathode

100%H2 100%H2 20%H2 80%Ar 20%H2 80%Ar 20%H2 80%Ar 20%H2 80%Ar 20%H2 80%Ar

AISI 1020 AISI 1020 AISI 430 AISI 430 AISI 430

Temperature (◦ C) 540 540 540 450

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Table 2 Sliding wear test conditions Test type Block material Ring material Ring dimensions (mm) Normal force (as-sintered samples), W (N) Normal force (nitrided samples), W (N) Tangential velocity, v (m s−1 ) Total distance, D (m) Medium Lubricant

Block-on-ring Sample AISI 52100 steel—750 HV30/10 Diameter 40, height 10 25, 35, 50, 70, 100, 150 25, 50, 70 0.314 565 Air None

paper (600 meshes). The specimen remained fixed while the disc rotated against it with an applied normal force. The tests were carried out by varying only the normal force, and keeping the others parameters fixed. Table 2 shows the test conditions. After the tribological tests the wear tracks were characterized by scanning electron microscopy (SEM) (Philips XL-30) of the surface and the cross-section. The semi-quantitative elementary composition analysis was carried out with energy dispersive X-ray (EDX) apparatus coupled to the scanning electron microscope. The microhardness profile was built using a MHT-4 Carl Zeiss microhardness tester. The X-ray diffraction (XRD) of the debris was carried out by an X’Pert Philips difractometer with a Cu anode. The roughness of the sample surface and the volume of the track was measured by a profilometer (Hommelwerke T8000), with a scanning velocity of 0.05 mm s−1 .

Fig. 2. Micrograph of the cross-section of the Cr-enriched sample.

3. Results and discussion 3.1. As-sintered specimens The hardness and microhardness of the materials are listed in Table 3. In the case of the P-Cr samples, a Cr-layer is present on the surface as a result of Cr enrichment during the sintering. This process has been explained in detail elsewhere [7]. The samples present about 10 wt% of Cr in solid solution only on the surface up to a depth of about 20–30 ␮m (Fig. 2). As has been shown [7], the porosity in the Cr-enriched layer is lower than that in a layer without Cr, as a consequence of the ␣-Fe stabilization at the sintering temperature (1150 ◦ C). The ferrite stabilization enhances the sintering, leading to a reduction in the surface porosity. The roughness of the samples as well as of the counter-body is shown in Fig. 3. The porosity effect of the sintered samples was eliminated in the analysis. Fig. 4 shows the experimental worn volume as a function of the applied load for all the assintered materials, i.e., F-Fe, P-Fe and P-Cr samples, indicated by squares, circles and triangles, respectively. It should be noted

Fig. 3. Roughness values (Ra ) for all samples and counter-body sliding surface.

that the wear increases approximately with the applied load, but that there is a large experimental scattering. In addition, no relevant difference in wear can be detected among the three materials. This agrees with the fact that they are characterized

Table 3 Summary of characteristics of as-sintered specimens Symbol

Hardness (HV 10)

Ferrite microhardness (HV 0.02)

F-Fe P-Fe F-Cr

55.7 ± 4.0 52.1 ± 2.1 50.7 ± 2.4

94.8 ± 4.0 85.1 ± 3.9 Cr layer, 170 ± 36; core, 92 ± 3.2

Fig. 4. Worn volume vs. applied normal force for as-sintered samples (F-Fe, unalloyed iron sintered in furnace; P-Fe, unalloyed iron sintered in DC-plasma; P-Cr, surface Cr-enriched iron sintered in DC-plasma).

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Fig. 6. Surface SEM image of the counter-face with adhered wear debris.

Fig. 5. Wear surface of material F-Fe tested at 70 N. (a) L70, lower measured trace volume. (b) H70, higher measured trace volume.

by very similar microhardness values (Table 3). The counterbody has lower Ra than the samples values (Fig. 3). Such result suggests a less important effect of the abrasive wear in the early stage of the sliding of the counter-body asperities. A microstructural characterization of the wear traces after the tests, analyses of the wear debris and of the subsurface damage were carried out in order to identify the wear mechanisms present during the tribologic tests. In order to elucidate the mechanisms responsible for the large scatter in the experimental results, the samples with the highest and lowest wear rate measured for a normal force of 70 N, and

designated by H70 and L70, respectively, were analyzed in detail. The sample wear track volumes for H70 and L70 were, respectively, 4.3 and 1.5 mm3 . Examining the wear track of these specimens (Fig. 5), different morphologies of the worn surface can be noted. In the L70 sample (Fig. 5a) there is a dark gray area identified with an (A) and a light gray region identified with a (B). An EDX analysis carried out in regions (A) and (B) in Fig. 5a showed that the (A) areas have high oxygen content whereas no oxygen was detected in the (B) areas. These results suggest that the (A) areas are formed by metallic and oxidized sintered debris. It is assumed that the sample track was subjected to sliding with a protective layer during the wear test. The SEM image of the wear track of the H70 sample presents microploughs in the sliding direction which are most probably correlated with the abrasive wear. Such microploughs were probably produced by the agglomerated debris adhering to the counter-body face sliding over the body surface. Fig. 6 shows a SEM image of the counter-body with agglomerate adhering to the sliding face. However, such evidence only illustrates the state of the sample immediately after the test. It is not possible to confirm the predominant wear mechanism. Abrasive wear would lead to more metallic debris, as a result of the formation of more debris. In the case of the fragments which stay in the friction region, a higher proportion of oxidized debris would be expected as a result of the high level of plastic strain, as has already been

Fig. 7. XRD spectra of the wear debris. (a) L70, lower measured trace volume. (b) H70, higher measured trace volume.

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mentioned. The debris produced by the H70 and L70 samples was collected and XRD analysis was carried out. The XRD patterns of the debris collected at the end of the tests are shown in Fig. 7. In both cases the presence of ␣-Fe and iron oxide peaks was observed. However, in the case of the L70 fragments (aspectrum) the metallic iron peak (␣-Fe) is proportionally lower than in the case of the H70 fragments (b-spectrum). As the porosity plays an important role in the formation and trapping of the debris [14] much of the debris produced by the dry sliding contact is assumed to remain in the friction region during several revolutions of the counter-body. Such fragments probably adhere to the face and the counter-face and are subjected to a high local pressure. They are mechanically comminuted on the revolutions and despite the relatively low temperature they undergo oxidation due to the high level of plastic strain [3,13]. The oxidized fragments could either leave the friction region or remain in the contact area [15,16]. In the latter case, the oxidized particles mix with metallic debris and as a result of high local pressure and temperature, the metallic debris could sinter forming a resistant agglomeration of metallic and oxidized wear debris which adheres to either the body or the counter-body faces. Considering the measured wear values (Fig. 4), the images of the worn traces (Fig. 5a and b) and the debris XRD analyses (Fig. 7), it can reasonably be supposed that in some cases a transition from adhered wear debris on the face and the counterface occurred during sliding, implying a wide variation in the total measured wear. It could be also argued that the porosity contributes to the large scatter in the experimental results, i.e., to determining the adhesion of the wear debris on either the face or the counter-face of the worn parts. The cross-sections of the wear tracks of samples H70 and L70 are shown in Fig. 8. In both cases, the grain boundaries and pores are deformed along the sliding direction. This is more obvious in the case of sample L70, i.e., when a low worn volume was measured. Fig. 9 shows the microhardness profile close to the wear track for both samples. It can be observed that the surface strain-hardening was particularly intense on the sample with a severe-to-mild transition. This result can be explained by considering that: when the debris is attached in the counter-face, the wear fragments are removed more rapidly and thus the plastic deformation in the sub-surface layers is reduced. In the material with debris adhered to the face, the internal material remains under stress for longer, thus inducing deeper strain hardening and closing the surface pores.

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Fig. 8. Cross-section micrographs of wear traces of F-Fe sample. (a) L70, lower measured trace volume. (b) H70, higher measured trace volume.

the furnace without Cr enrichment, the characteristics of the nitrided layer are similar. The presence of extensions of the compound layer in the grain boundaries near the surface of the F-FeN540 sample (Fig. 10a) should be noted. This characteristic was rarely observed in the P-FeN540 sample, which only showed a uniform white layer (Fig. 10b). The Cr-enriched samples sintered in plasma show that Cr considerably enhances the hardness of the compound layer and the depth of the hardened region (Fig. 12). Spalvins [17] and many others [18–20] have observed that the relatively high concentration of Cr (>5 wt%)

3.2. Nitrided specimens Fig. 10 shows the microstructure of the nitrided samples (FFeN540, P-FeN540, P-CrN450 and P-CrN540). Fig. 11 shows the XRD spectra of the nitrided samples. In the case of the unalloyed iron nitrided samples, only the ␧-phase (Fe2–3 N) and the ␥ -phase (Fe4 N) are observed. Cr nitrides were detected in the Cr-enriched specimens. The ␧-phase, traces of the ␥ -phase, Cr2 N, and CrN were detected in the P-CrN450 specimen. In the case of the P-CrN540 sample only the ␧-phase and Cr2 N were noticed. The relevant microhardness profiles are presented in Fig. 12. In the case of the specimens sintered in plasma and in

Fig. 9. Microhardness profiles of the samples tested at 70 N with lower (L70) and higher (H70) measured trace volume.

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Fig. 10. Micrographs of nitrided samples. (a) F-FeN540 sintered in furnace and nitrided in plasma at 540 ◦ C. (b) P-FeN540 sintered in plasma and nitrided in plasma at 540 ◦ C. (c) PCrN450 sintered in plasma with Cr enrichment and nitrided in plasma at 450 ◦ C. (d) PCrN540 sintered in plasma with Cr enrichment and nitrided in plasma at 540 ◦ C.

in solid solution promotes the precipitation of Cr nitrides. If the concentration is higher than 10%, an abrupt steep decrease in the hardness profile was observed [18–20]. In the present work, the depth enriched with chromium in the samples is about 20–30 ␮m and the concentration in the surface is about 10% in weight [7] forming, as shown in the literature, high values of hardness and an abrupt, steep decrease in the microhardness profile. The results of the wear tests of the nitrided samples are shown in Fig. 13a. Under some conditions (mainly with an applied load of 25 N) the measured volume has a negative value. In such

cases, the track was not effectively formed, however, an area with wear debris transferred from the counter-face had adhered in the friction region. Thus, the method used here to measure the volume has indicated a negative worn volume. Fig. 3 is observed that the Ra values for nitrided samples were higher than the unnitrided ones. Then, higher abrasion of the counter-face is expected as a result of the sliding against higher and harder asperities of the nitrided samples. The counter-body underwent more wear due to the hardness of the white layer of the sample which exceeds that of the counter-body itself. Such an effect could be considered observing the mass variation of the samples

Fig. 11. XRD spectra of F-FeN540, P-CrN540 and P-CrN450 samples.

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Fig. 12. Microhardness profiles of nitrided samples.

Fig. 13. Dry sliding results for nitrided samples. (a) Average worn volume vs. applied normal force of nitrided samples. (b) Mass variation of samples tested with applied load of 25 N.

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(Fig. 13b). An increase in weight was noted for all the samples tested with an applied load of 25 N. Analyzing Fig. 13a one can suppose that the F-FeN540 and P-FeN540 show better dry sliding behavior than the Crenriched samples. However, when analyzing the cross-section micrographs of the samples tested with an applied load of 25 N (Fig. 14) a distinctive behavior among the unalloyed iron and Cr-enriched samples can be noted. No wear was observed in the samples nitrided without Cr enrichment (F-Fe540 and P-Fe540) when subjected to dry sliding with an applied load of 25 N. However, in the P-Fe540 sample (Fig. 14a) there were several transversal cracks in the compound layer. Moreover, in some regions the compound layer had been removed. Therefore, the sample shows surface damage as a result of the wear test, which is different from that suggested by the track volume results (Fig. 13a). The F-Fe540 sample showed identical behavior. The compound layer was not removed (Fig. 14b) in the Cr-enriched sample nitrided at 450 ◦ C (P-CrN450). Nevertheless, fewer transversal fissures were present in the compound layer, as indicated in Fig. 14b. It is believed that such fissures were produced by the combined effect of the applied load (perpendicular to the surface) and the adhesive forces (longitudinal to the surface). It was assumed that the hardened region was not sufficient to support the applied loads, so it is supposed to the substrate underwent plastic deformation as a consequence, tensile stresses arose in the compound layer leading to transversal fissures in the compound layer. The P-CrN540 sample has a hardened depth (>500 HV) higher than that of the other samples (Fig. 12). In this case, the substrate can support the applied load despite the brittle characteristic of the treated surface and therefore there was no damage in the compound layer (Fig. 14c). Applying 50 N of normal load, the F-FeN540 and P-FeN540 have behaved similarly when tested with 25 N. Fissures were observed in the compound layer. Moreover, due to the adhesive forces in the surface the fractured compound layer rotates following the sliding direction (Fig. 15a). Therefore, a relatively high level of damage in the surface of the F-FeN540 and P-FeN540 should be noted (Fig. 15a). A surface with such a characteristic favors abrasive wear and the accumulation of debris in the surface inducing the agglomeration of a protective layer of oxidized and metallic sintered debris. These agglomerated particles protect the compound layer from subsequent wear and surface damage [15,16]. So, for these cases, it can be supposed that the surface was damaged in the early stages of the tests and, as a result of the formation of the agglomerated particles layer, the surface damage was considerably reduced. A total removal of the compound layer (Fig. 15b) was observed in the Cr-enriched samples nitrided at 540 and 450 ◦ C. Due to the high hardness of the compound layer and the abrupt, steep decrease in the microhardness profile, compound-layer fragmentation occurs more easily than in samples without Cr-enrichment. The chromium nitride debris, present in the friction region, probably acts as abrasive particles, significantly increasing the wear rate of the material while it is present in the friction area. However, the compound layer of the Cr-enriched samples may resist more than the samples without Cr-enrichment. The wear process probably occurs progressively, but throughout the test, and as a conse-

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Fig. 14. Wear traces micrographs of cross-sectioned nitrided samples tested with applied load of 25 N. (a) P-FeN540 sintered in plasma and nitrided at 540 ◦ C. (b) P-CrN450 sintered in plasma with Cr enrichment nitrided at 450 ◦ C. (c) P-CrN540 sintered in plasma with Cr enrichment nitrided at 540 ◦ C.

quence, the total wear volume is higher than that observed for the samples without Cr enrichment. Therefore, the Cr-enriched samples which were subsequently nitrided have a worn volume higher than that measured for the samples treated without Cr enrichment for applied loads of 50 N.

Fig. 15. Wear traces micrographs of cross-sectioned nitrided samples tested with applied load of 50 N. (a) P-FeN540 sintered in plasma and nitrided at 540 ◦ C. (b) P-CrN540 sintered in plasma with Cr enrichment and nitrided in plasma at 540 ◦ C.

The samples tested with an applied load of 70 N show a similar wear behavior to those tested with an applied load of 50 N. The exception was the P-FeN540 sample. As discussed earlier, due to the large grains the compound layer extensions are less frequently observed. As a result, the compound layer in this sample does not last as long. Such layer extensions may play an important role in the integrity, and rapid deterioration of the compound layers, regardless of the presence of fissures in the layers. As a consequence, the sample experienced a higher wear rate than the F-FeN540 sample when subjected to tests with 70 N as a result of the total elimination of the compound layer. Observing the wear behavior of the as-sintered sample and the nitrided samples, it can be said that, in general, the nitriding treatment increased the wear resistance of the sintered parts. The as-sintered samples had a wear rate one order of magnitude higher than the nitrided samples. If the selection criterion is the mass loss, then the nitrided unalloyed samples sintered in the furnace behaved better. However, if the criterion is related to the surface integrity, the Cr-enriched sample nitrided at 540 ◦ C behaved better. When subjected to low applied loads (25 N) the integrity of the compound layer of this sample remained intact. Table 4 summarizes the results of all the wear tests. The values of the specific wear rate (W) were calculated using the maximum and minimum measured wear track volume. For the as-sintered samples, the specific wear rate (W) was inside the range observed by Straffelini and Molinari [3]. Moreover, these values were about one order of magnitude higher than those of the nitrided samples. The friction coefficient measured in the test is also given in Table 4. Moreover, friction coefficient diagram was shown in Fig. 16 for samples tested under 50 N of applied load. Unfortunately, no wear mechanism transition could be detected in the friction coefficient during the test. Considering the scattering, it should be noted that the friction coefficient of the nitrided sample is slightly higher than that of the as-sintered samples. Despite the results of some authors whose friction coefficient was lower for nitrided samples when compared with unnitrided ones [9,21], others have presented an increase in the friction coefficient after nitriding [10] similar to that presented

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Table 4 Summarized results obtained with the wear tests Sample

Specific wear rate, W (mm3 mm−1 N−1 )

Friction coefficient, μ (50 N)

Wear mechanism

F-Fe

1.9 × 10−8 –1.1 × 10−7

0.60–0.75

P-Fe

2.6 × 10−8 –1.6 × 10−7

0.60–0.75

P-Cr

1.3 × 10−8 –9.3 × 10−8

0.60–0.70

F-FeN540



0.70–0.85

P-FeN540

0–2.7 × 10−9

0.70–0.85

P-CrN540

0–4.5 × 10−9

0.70–0.80

P-CrN450

0–6.26 × 10−9

0.70–0.80

Adhesive wear, debris oxidation and wear by microploughing with eventual localized oxidation Adhesive wear, debris oxidation and wear by microploughing with eventual localized oxidation Adhesive wear, debris oxidation and wear by microploughing with eventual localized oxidation Plastic deformation in the substrate, fissures in compound layer and subsequent compound layer damage. Formation of protective layer of agglomerated particles Plastic deformation in the substrate, fissures in compound layer and subsequent compound layer removal Plastic deformation in the substrate, fissures and eventual fragmentation of the compound layer producing abrasive wear Plastic deformation in the substrate, fissures and eventual fragmentation of the compound layer producing abrasive wear

Fig. 16. Friction coefficients diagram for nitrided and unnitrided samples (FN = 50 N).

of the sample to support the applied normal load. For soft materials, it is assumed that a deeper, hardened layer is preferable to a harder and thinner layer. For the samples nitrided without Crenrichment, it was found that the compound layer was damaged in the early stages of the wear tests and a tribolayer protected the sample from subsequent wear. The Cr-enriched nitrided samples are supposed to show a lower wear rate in the compound layer, but wear takes place progressively throughout the entire wear test, resulting in a higher worn volume. Moreover, the Crenriched sample nitrided at 540 ◦ C did not show measurable wear or compound layer deterioration when tested with 25 N of normal load. Under the test conditions used in this work, it was verified that the nitriding treatment improved the wear resistance of the sintered samples by one order of magnitude. Acknowledgements

in Table 4. Discussion of these results is particularly complex. Detailed experiments are needed to elucidate such behavior.

This work was supported by grants from CAPES, Brazil, Finep/MCT (PRONEX), Brazil, and CNPq, Brazil (PADCT).

4. Conclusions References Dry sliding tests were performed using a block-on-ring configuration on powder metallurgy iron specimens sintered in a furnace and in an abnormal glow discharge with and without superficial Cr enrichment. Some of the specimens were thermochemically treated by plasma nitriding. In the as-sintered samples, the wear mechanism was probably controlled by the formation of an agglomerated of wear debris particle on either the face or the counter-face. If these agglomerated particles adhered to the counter-face, the sample underwent abrasive wear, and if it adhered to the sample surface, the wear debris protected the sample from subsequent wear. It is assumed that the alternation in these wear mechanisms was caused by the surface porosity and the test geometry leading to high scattering in the wear results for all the as-sintered samples. The nitrided samples without Cr enrichment did not show measurable wear, however, the compound layer lost its integrity. It is assumed that the integrity of the compound layer is a function of the capacity

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