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alloying and spark plasma sintering
Wear 376-377 (2017) 958–967
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Tribological behaviour of Cu based materials produced by mechanical milling/alloying and spark plasma sintering Massimo Pellizzari, Giulia Cipolloni n Department of Industrial Engineering of the University of Trento, Via Sommarive 9, 38123 Trento, Italy
art ic l e i nf o
a b s t r a c t
Article history: Received 14 September 2016 Received in revised form 4 November 2016 Accepted 25 November 2016
Aim of the work is to investigate the tribological behaviour of Cu based materials produced by mechanical milling (MM), mechanical alloying (MA) and finally sintered by Spark Plasma Sintering (SPS). MM causes a significant strain hardening of Cu, while MA by 0.5 wt%TiB2 adds a further dispersion hardening contribution. During dry sliding contact against high speed steel a close relation has been found between friction coefficient and contact temperature highlighting two regimes. The first one, occurring at an early stage of the process, is characterized by a high friction coefficient ( 1), typical of the strong adhesion of Cu–Cu contact. The second one is characterized by a lower steady-state friction ( 0.7) related to the change of wear mechanism from adhesive metallic into triboxidative. By increasing the load the friction coefficient decreases, the transition time from the adhesive to triboxidative friction regime decreases and the contact temperature increases. Under abrasive wear conditions the penetration depth of the indenter is progressively reduced by increasing hardness, highlighting the benefits of MM and, even more, of MA. The Cu þ 0.5 wt.%TiB2 composite alloyed for 240 min shows the best wear resistance under both wear conditions. The good wear resistance combined to high thermal conductivity represent very attractive properties for many applications. & 2017 Elsevier B.V. All rights reserved.
Keywords: Copper Mechanical milling Dry sliding Adhesion Tribo-oxidation Abrasion
1. Introduction In the last years the demand of materials combining high wear resistance and high thermal conductivity is continuously increasing, especially for components as plastic injection moulds, actively cooled component, continuous casting moulds, electrical contacts and welding electrodes [1–4]. All these applications demanding rapid heat removal from the surface region, require a high heat transfer capability to achieve faster processing routes and, consequently, lower costs. In this way, Cu seems to be a suitable candidate due to its high intrinsic thermal conductivity. By the way the low hardness, strength and wear resistance of Cu restrict its performance and, definitely, its applicability. Therefore, different solutions have been evaluated to improve strength without significantly affecting conductivity. In the last years, severe plastic deformation (SPD) by mechanical milling (MM) has been proposed as an efficient method to improve the strength of Cu by grain refinement and strain hardening [5–8]. An even more extended strengthening can be obtained by uniform dispersion of a harder second phase in the copper matrix by mechanical alloying (MA) [6–8]. This leads to the production of metal matrix n
Corresponding author. E-mail addresses: [email protected] (M. Pellizzari), [email protected] (G. Cipolloni). http://dx.doi.org/10.1016/j.wear.2016.11.050 0043-1648/& 2017 Elsevier B.V. All rights reserved.
composites (MMC), which have attracted growing attention [6]. The success of these materials is related to the possibility to tailor their properties using the synergistic effect of its constituents. The tribological behaviour of MMC's is associated to the different role of the metallic matrix and the reinforcement. Nevertheless it is strictly correlated to the specific testing conditions [9– 11]. Under abrasive conditions, the metal matrix must incorporate and provide suited support to the hard phase so to avoid its premature pull out or fracture. In order to guarantee a suited protection, the hardness of hard particles (HP) has to be higher than that of the abrasive medium (HA) and should also have a high fracture toughness to prevent fracture [12–15]. Usually the wear resistance of materials in abrasion test is improved by changing the abrasive particle movement from gliding to rolling. This is achieved by increasing either the hardness ratio between matrix and HA, or by increasing HP/HA [13]. Also the nature of interface between the constituents is a critical and essential feature for the performance of a MMC because it affects the ability of the matrix to hold the reinforcing particle and prevents its pull out under abrasive wear conditions. The extension of interface is strictly related to the particle size of the reinforcement because it determines the fraction of bonding area between the constituents and consequently the load bearing capability of interfaces. Usually the HP must be bigger than the abrasive particles, this reduces the
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abrasion depth if they are well bonded to the matrix [9]. Moreover the mean free path between the HP must be smaller than or at least as large as the width of the abrasive particle, if possible to act as an obstacle to the scratching abrasive [9,16]. If the HP are too small they will be easily pulled out unless they are present in large quantities obstructing the penetration of the abrasive grain into the softer matrix. By the way it must be considered that over a certain volume fraction percolation of HP may occurs forming brittle network leading to a premature failure of the component [9,17]. Under incremental load scratch testing conditions, four microscopic mechanisms and processes of scratch deformation have been observed, e.g. elastoplastic grooving, plastic plowing, microcracking, and the inter and intra-fractures and chipping of the surface scratched [18]. By increasing the amount of second phase the magnitude of damage is reduced [18]. Under dry sliding conditions the materials performance of MMC is not only related to the hardness of the material but also to the capability to activate self-protective mechanisms depending on the variable service conditions. For copper based materials the wear behaviour against steel is initially characterized by the surface plastic deformation at the contacting asperities favoured by a shear instability effect [19–23]. This effect gives rise to the occurrence of intense transfer phenomena with the formation of a mechanically mixed layer on the surface of both wearing materials [23–25]. If the interaction with the environment and the service conditions promote oxidation, a tribo-oxidative wear mechanism governs the wear behaviour. The formation of a protective oxide tribo-film improves the wear behaviour reducing the friction coefficient [26–30]. Under these conditions, a harder material provides higher load bearing capability to the oxide layer improving wear resistance [29]. In general for higher hardness the plastic deformation is resisted and material removal is more difficult, consequently the sliding marks on the surface appear shallower. Moreover the higher hardness can also help to restrain effectively nucleation or propagation of the crack generated during testing [30]. By the way the formation of groove on the worn surface cannot be completely avoided during sliding and the sub surface cracks along the sliding direction usually promote the detachment of wear particles in the form of sheet and flake leading to a delamination process [25]. In general by increasing the content of reinforcement the wear resistance increases independently from the type of reinforcement [17,31,32]. In Cu–TiB2 composite has been demonstrated that increasing the content of hard phase, wear resistance increased considerably if the distribution of reinforcement in the matrix is uniform [33,34]. Occasional fracture at the reinforcement-matrix interface was shown to occur under high load conditions, especially for large reinforcement particle size (445 mm) [34]. In this paper copper based materials have been produced by MM, MA and Spark Plasma Sintering (SPS) to increase the wear resistance by strain hardening and dispersion hardening keeping acceptable value of thermal conductivity. SPS allows sintering at lower temperature and in a shorter time comparing to more conventional processes. Indeed the high heating rate peculiar of SPS preserves the fine microstructure produced by MM and MA, and reduces any possible interaction between the metal matrix and the reinforcing particles [35,36]. Aim of the investigation is to give a contribution to the understanding of the dry sliding and abrasion wear behaviour of copper based materials and, in particular, to highlight the influence of the load and the role of the material hardness and microstructure obtained by the MM and MA. To this end, a first series of dry sliding tests against high speed steel (AISI M3:2) were carried out at different loads. Special emphasis was given to the role of microstructure in tribolayer formation. Moreover, the tribological behaviour of mechanical milled and mechanical alloyed
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materials has been also studied, under pure abrasive conditions by means of the scratch test. The results are expected to provide important criteria in designing novel Cu based tribomaterials.
2. Materials and experimental methods 2.1. Materials A water atomized copper powder (AT-Cu) with a particle size lower than 75 mm was considered for MM. For the production of MMCs 0.5 wt% of TiB2 with particle size finer than 3 mm was added to AT-Cu by MA. Both processes were carried out using a Fritsch Pulverisette 6 planetary mono mill at 400 rpm. Vial and sphere with 10 mm diameter of 100Cr6 (63 HRC) were used. The system was evacuated down to 130 Pa and 0.5 wt% of stearic acid was added as process control agent. MM was carried up to 6000 min in order to obtain different final hardness. MA up to maximum 240 min, time necessary to achieve a good dispersion of HP in the s Cu matrix. The milled powders were sintered in a DR.SINTER SPS1050 (Sumitomo Coal & Mining, now SPS Syntex, Inc.) apparatus with graphite punches and dies. SPS was performed at a nominal temperature of 950 °C, with a uniaxial pressure of 60 MPa applied at 700 °C. The heating rate was 100 °C/min up to 900 °C and 50 °C/min up to the sintering temperature. The maximum temperature and pressure were held for 5 min before allowing the furnace to cool to room temperature. The properties of MM and MA samples were compared to those of AT-Cu and a commercial Cu–Be alloy (Moldmax HH). 2.2. Experimental methods The density of the SPS-processed pellets was measured according to Archimedes' principle (ASTM B962-08). For mechanical milled samples the relative density was calculated considering the theoretical densities of Cu as 8.96 g/cm3. In the case of MMC the theoretical absolute density was calculate according to the linear rule of mixture and considering the theoretical density of TiB2 as 7.76 g/cm3. Sintered samples were cut into disk-shaped specimens of diameter 10 2.7 mm and the thermal conductivity (k) of the samples was derived by the equation: k ¼ α ρ Cp. Where α is thermal diffusivity, ρ is the bulk density, and Cp is specific heat capacity. The thermal diffusivity and the specific heat were measured using the NETZSCH laser flash apparatus LFA 467 HyperFlash. All samples were tested with the LFA between 400 and 500 °C with 20 °C temperature steps. Prior to the measurement the front and the back of the samples were coated with graphite to enhance the emission/absorption properties of the sample. The specific heat was determined by the reference method given by ASTM-E 1461– 2011. The LFA was calibrated with a Cp-standard of pure-copper. The sintered pellets were machined into block samples with the dimension of 9x13 5 mm. The 9 5 mm cross section faces were cold molded in epoxy resin for metallography. Standard metallographic preparation, including grinding with SiC papers up to 4000grit, final polishing with 3 μm and 1 μm diamond pastes, was performed. A chemical etching with 120 mL of distilled water, 30 mL of hydrochloric acid and 10 g of iron chloride was used [37]. Vickers micro-hardness was measured using an applied load of 1 N, holding time of 10 s and a loading rate of 0.1 N/s according to ASTM standard E9-08. The 9 13 mm faces of the machined block were employed as the wear surface both for sliding and abrasion wear tests (Fig. 1a). The wear surface have been polished up to 1 micron diamond paste. The dry sliding tests were carried using a block-on-disc configuration with a rotating speed of 300 rpm, corresponding to a
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Fig. 1. Sample block (a) and sliding wear configuration (b).
sliding speed of 0,63 m/s (Fig. 1b). As counterpart a disc of AISI M3:2 (830HV30) has been used. The tests have been carried out at 50 N for 15, 30 and 240 min, and for 100 and 200 N for 30 and 240 min. The worn abrasive disc and the copper based samples were substituted by a new one for every sliding time and load. During the experiments, the friction coefficient was recorded and the contact temperature was continuously measured by a thermocouple placed in a hole at 2 mm depth from the contact surface (Fig. 1b). The wear curve was recorded by periodically weighing the specimens using a precision balance (0.0001 g). The wear rate (Ka) has been calculated as ratio between volume loss and sliding distance. The wear mechanism has been analysed by light optical and scanning electron microscopy of the wear tracks and debris. Chemical analysis was carried out by energy-dispersive X-ray spectroscopy (EDXS). The abrasive wear resistance has been evaluated by the scratch test at constant load, using a Rockwell diamond indenter with 200 μm of diameter. Different loads (1, 2 and 3 N) were applied for an abrasion distance of 5 mm and with a sliding speed of 10 mm/ min. The output of the test is the penetration depth with the accuracy of 71 μm. The press surface, i.e., the sintered plane perpendicular to the pressing direction during SPS was indented/ scratched. Each test surface was polished down to 1 micron diamond paste.
3. Results and discussion 3.1. Materials characterization In Fig. 2 the OM micrographs of all the sintered samples are reported. The sample sintered from AT-Cu powder shows a uniform microstructure (Fig. 2a). Some pores are still evident and give a pinning effect for the grain growth [38]. For MM and MA the microstructure of sintered materials keeps memory of the original particles morphology, a severe anisotropy being observed when flaking prevails for MM-240’ and MMC-80’ (Fig. 2-b and -e). Instead for prolonged milling time, when welding events predominate, the microstructure becomes more uniform e.g. MM720’ and MMC-240’ (Fig. 2-c and -f). The large milled particles are clearly recognizable and they are surrounded by a network of recrystallized copper film (white region indicated by arrows) due to an intense overheating at the contact points between the powders during SPS [39]. In MMC by increasing milling time the dispersion of the hard phase becomes more homogeneous, TiB2 particles are more and more refined and the adhesion between the two
constituents is improved. MM-6000’ shows a uniform microstructure (Fig. 2-d) because the milled powder consists in small agglomerations of very thin flake like particles that are easily compacted during SPS [40]. Table 1 summarizes the properties of the six sintered sample. For MM the best combination of hardness has been obtained after 240 min, with an increase of hardness from 103HV0.1 for ATCu to 150HV0.1 for MM-240’. Despite the more extensive strain hardening induced by longer MM a decrease of hardness is observed for MM-720’ and MM-6000’. This is related to an increase of the residual porosity associated to the incomplete decomposition of the process control agent during SPS which releases gaseous species that cause the formation of pores and reduce the final density (Table 1) [41]. A different trend is shown by MMC's, where the hardness increases with milling time due to the progressive refinement of TiB2 particles and their improved distribution. The hardness of MMC-240’ almost doubles that of AT-Cu sample, increasing from 103 HV0.1 to 207 HV0.1. MMC-80’ and MMC-240’ exhibit density values of 98 and 97% respectively: the increase of hardness increasing milling time confirms that strain hardening and dispersion hardening during MA abundantly compensates for the negative effect of porosity. The thermal conductivity of AT-Cu is (300 W/mK) is slightly lower than that of pure Cu (400 W/mK) due to the deleterious presence of oxide particles and pores. The effects of MM and MA are quite different. On one side, a detrimental effect of MM has been highlighted especially for long milling time. MM-240’ exhibits an unexpected increase of thermal conductivity, meanwhile MM-720’ and MM-6000’ show a drastic decrease of thermal conductivity due to the higher porosity and the presence of crystal defects [42]. On the other side, MA exhibits some interesting and promising results in contrast with what expected. The thermal conductivity of MMC-80’ and MMC-240’ is very close to that of AT -Cu, in spite of the mechanical strain hardening and the addition of TiB2. The small amount and the good dispersion of reinforcement, especially for long milling time, provide an excellent combination of hardness and thermal conductivity. All the materials produced by MM and MA show a higher thermal conductivity than the commercial Cu–Be alloy. 3.2. Dry sliding wear 3.2.1. Friction A typical diagram of the friction coefficient and the contact temperature recorded during the dry sliding test is reported in Fig. 3, where a clear correlation between these two parameters is evident. Two distinct friction regimes can be distinguished: an
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Fig. 2. Optical micrographs of the microstructure of (a) AT-Cu, (b) MM-240’, (c) MM-720’, (d) MM-6000’, (e) MMC-80’ and (f) MMC-240’. Table 1 Properties of the sintered materials. Sample
Milling time (min)
HV 0.1
Rel. Density (%)
Thermal Conductivity (W/mK)
AT-Cu MM-240’ MM-720’ MM-6000’ MMC-80’ MMC-240’ Cu-Be
0 240 720 6000 80 240 –
103 72.1 150 71.2 1107 2.3 457 1.8 130 71.9 207 71.2 400 71.2
99 94 90 81 98 97 100
300 320 270 200 298 297 106
early stage, characterized by a high friction coefficient ( 1) and higher temperature followed by a steady-state regime showing lower friction coefficient ( 0.7) and temperature. The high value
of friction coefficient in the early stage is typical of an adhesive metal-metal contact. The observed fluctuations remark the occurrence of stick-slip phenomena accompanied by material transfer [19–21]. In this stage both temperature and friction increase to show a drop to lower steady state values after a certain transition time (tss). It is well recognized that the frictional heating facilitates surface oxidation, promoting the formation of protective oxide layers. This results in milder wear conditions characterized by metal-oxide, oxide-oxide contacts and in a decrease of the friction coefficient (Fig. 3) [29]. Definitely, after a certain time (tss) a transition from an adhesive to a triboxidative wear mechanism is evidenced [12,17,19,26]. At this stage the oxide layer acts as a solid lubricant reducing the resistance to sliding. The temperature consequently decreases keeping a constant value for the remaining duration of the test. Although all the materials show a similar macroscopic friction
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protective oxide layers. One the other one, harder materials show a more difficult transition from not conformal to conformal contact, so that wear debris are less easily entrapped in the contact region to form a protective oxide layer. This is well evident for MMC-240’ characterized by a hardness two times higher than AT-Cu (Table 1). Fig. 4 also shows that the steady state friction of MM-240’ and MA-240’ are higher than AT-Cu. The plausible reason is the improved third body abrasive contribution of the harder wear particles.
Fig. 3. Friction coefficient and temperature evolutions for 240 min of dry sliding test of MMC-240’.
Fig. 4. Comparison of the friction coefficient evolution for 240 min of dry sliding test of AT-Cu, MM-240’ and MMC-240’.
coefficient evolution under dry sliding condition, some differences can be highlighted between MM and MA (MMC) samples. In Fig. 4 the friction evolutions for 240 min sliding test of AT-Cu, MM-240’ and MMC-240’ are reported. Comparing AT-Cu with MM-240’ and MMC-240’ it can be noticed that increasing the hardness the time required to achieve the triboxidative wear regime (tss) becomes longer. On one side, harder materials produce less debris, the particles required to form
3.2.2. Wear The above statements regarding the friction behaviour and the wear mechanism can be confirmed looking at the changes in structure and chemical composition of sliding surfaces, after different sliding times, before and after tss. For this reason the sliding test were carried out for 15, 30 and 240 min to monitor and analyse the evolution of tribological layer during the transition of the wear behaviour from the adhesive mechanism to the tribo-oxidation mechanism. The characterization of the wear tracks and of the discs have been carried out, and optical micrographs after 15 and 240 min of sliding test are reported respectively in Figs. 5 and 6. After 15 min (see Fig. 3), an accumulation of material is highlighted in the wear track, (Fig. 5-a). The wear surface exhibits a quite rough aspect, with an inlet area characterized by deep abrasion scratches and an outlet area showing an intense accumulation of material. A strong materials transfer from the Cu block to the disc (high speed steel) surface is highlighted (Fig. 5-b) as confirmed by the EDXS analysis, (Fig. 7-a). Large patches of copper are transferred on the surface of the disc and tent to shear in the sliding direction, favoured by the high ductility of the Cu matrix, (Fig. 5-b). The evidence of a Cu–Cu adhesive contact is clearly proved. The high frictional heat during the running stage enhances the generation of oxide and consequently the change of the wear mechanism from adhesion mechanism to tribo oxidation. Moreover the formation of the wear track results in the change of the contact nature from not conformal to conformal. This process leads to the production of wear debris which can be transferred on the disc, i.e. adhesion, either can be packed inside the track or pulled out of the wear system. The creation of the wear track aids firstly the permanence of the wear debris in the tribological system and then their oxidation. The optical micrographs after 240 min (Fig. 6-a) of sliding confirm the change of the wear mechanism in accordance with Fig. 3. In Fig. 6-a the wear tracks is larger and cleaner than in Fig. 5-a, in addition any massive accumulation of material is detected in the outlet area and scratches due to the abrasion action
Fig. 5. Top view of wear track (a) and disc surface (b) after 15 min of dry sliding test of MMC-240’.
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Fig. 6. Top view of wear track (a) and disc surface (b) after 240 min of dry sliding test of MMC-240’.
are vanished in the inlet region. The oxide layer in the track is much thinner and homogeneous, except for some darker stripes on the left side of the track (Fig. 6-a) Also the surface of the disc drastically changes, the transfer of material is drastically decreased and a sever oxidation is evident (Fig. 7) as also confirmed by the EDXS analysis (Fig. 8). Since copper oxide is more brittle than Cu, it breaks more easily and the fragments create more quickly debris and with that the protective oxide. The brittleness of the copper oxide is also confirmed by the finer dimension of the compacted and transferred particles on the disc after 240 min in comparison to the stretched and spread patches after only 15 min, (Fig. 5-b and Fig. 6-b). The continuous formation and fragmentation of oxide layer during prolonged sliding wear consists in a dynamic phenomenon which contributes to the wear damage and to the production of wear debris. Good evidenced has been provided to show that the oxide formed by frictional heat grows until, at a critical oxide film thickness, about 10micron, it spalls off as wear debris. Once fragmentation of the oxide layer occurs part of the loose debris particle may become entrapped, fragmented and compacted in the contacting surfaces contributing to the protective action of the oxide layer (Fig. 9) [29]. Some other debris, instead, are pulled out of the wear system. The typical morphology of debris of MMC-240’ collected after 240 min is reported in Fig. 10. Large plate-like oxide fragments (Ox in Fig. 10) are present, confirming the occurrence of delamination due to the large shear strain during sliding. The dimension of the plate-like debris (4200 μm) are in accordance with the size of the oxide patches on the wear track (Fig. 9). Some debris show a shiny surface due to the intrinsic
metallic nature (M) and probably they correspond to the side attached to the wear track surface before spalling. It is also plausible that these metal debris are those produced by delamination during the early stage, as the adhesive mechanism predominates. On the other hand, the finer fraction of debris corresponds to fragmented oxide particles. 3.2.3. Wear resistance In Fig. 11 the wear rate for all the samples tested for 240 min at 50 N are reported. The dependence of wear rate on hardness is very clear, increasing hardness the wear coefficient decreases. By the way the wear rates of material produced by MM and MA are higher than Cu–Be and similar to AT-Cu, except for MM-6000’ and MM-720’. For hardness value lower than 100 Hv the soft substrate limits the load bearing capability of the material (MM-6000’) to sustain the oxide layer leading to a higher wear rate. Also the lower thermal conductivity of MM-6000’ ( 200 W/mK) (Table 1) may negatively affect the wear behaviour because, if the heat is not properly removed, the material undergoes softening and its load bearing capability is further decreased. When hardness exceed 100HV0.1 the permanence of the oxide layer is promoted and lower value of wear rate are guaranteed. Only if the oxide is firmly attached to the sliding surface it can get compacted protecting the underlying base alloy. The presence of a mixed layer in the wear track will influence its hardness and other mechanical properties, which in turn will influence friction and wear (Figs. 4 and 11). It must be considered that a higher hardness leads to the formation of harder debris, consequently the protective oxide layer will be harder providing a lower wear rate. For these materials the thermal
Fig. 7. SEM micrographs of the copper transferred particle on the disc surface after 15 min (a) and 240 min (b) of dry sliding test.
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Fig. 8. EDXS spectra of the copper transferred particle on the disc surface after 15 min (a) and 240 min (b) of dry sliding test.
conductivity higher ( 300 W/mK) than MM-6000’ aids the dissipation of the heat avoiding softening. In spite of the relatively low hardness MM-720’ shows an exceptionally low wear rate (Fig. 11), even lower than Cu–Be. The investigations carried out until now could not clarify the reason for this results. Detailed microstructural analysis are ongoing. With exception of MM-6000’ and MMC-80, the wear coefficients decrease as the load is increased to 100 N (Fig. 11). It is evident that increasing the load the positive influence of hardness becomes less important, so that all materials perform similarly to Cu–Be alloy. This is very clear for soft AT-Cu which is supposed to have a low wear resistance due to its low hardness. By the OM observation of the worn surfaces of AT-Cu an intense accumulation of wear debris inside the wear track is evident in Fig. 12-a. The wear products are stuck and compacted inside the track, thus altered the wear loss measurement and consequently the calculation of the wear rate. This is important for the performance of the material because such material accumulations inside, for example, an injection mold can lead to dimensional changes not tolerated by the design. All the others materials show a similar morphology of the wear surface, with the presence of an oxide protective layer but thinner than that in the case of AT-Cu, Fig. 12-b. In Fig. 12-b MMC-240’ exhibits a wear track smaller than for 50 N test (Fig. 6-a), but this is perfectly in accordance with the lower wear rate depicted in Fig. 11. 3.2.4. Effect of the sliding load The effect of the increasing load results in an increased contact temperature and a decreasing friction coefficient (Fig. 13). According to the mechanism explained above, a higher temperature facilitates the generation of oxide debris and assists their
Fig. 10. The morphology of the wear debris of MMC-240’ after 240 min of sliding: the metallic nature (M) and the fine oxide powder (Ox).
compaction for the formation of the wear protective oxide layers [10,11,25,27,30]. It must be considered that during the steady-state regime two competitive processes occur: the break-down of the protective layer, resulting in wear, and the compaction/sintering of wear debris, to reform the oxide layer [29]. The latter process, as the oxidation of the particle, occur more rapidly as the temperature is increased and friction decreases (Fig. 13-b) [12]. The transition from adhesion to tribo-oxidation shifts towards shorter time so that the running in almost disappears at 200 N (Fig. 13-b). Moreover, the friction becomes more stable, stick and slip phenomena are less frequent and adhesion is drastically limited. All these phenomena positively affects the wear rate which decreases
Fig. 9. Detailed view of the worn surface of MMC-240’after 240 min of sliding (a) showing the formation of protective Cu oxide patches (EDXS is referring to the darker region in Fig.a highlighted by the white cross) (b).
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Fig. 11. Dependence of wear coefficient on hardness for long sliding test.
by increasing the load (Fig. 11). All the materials show the same behaviour by increasing the load, as reported in Fig. 14. The lowest and the highest values of friction coefficient at 50 N corresponds to AT-Cu and MM-6000’ respectively, in all the other cases the values are included in a similar range. In Fig. 11 it can be noticed that MM-6000’ and MMC-80’ show the highest wear rate both at 50 and 100 N. In the case of MM6000’ this can be ascribed to the low hardness, i.e. to the insufficient support by the base material to the oxide layer, which is thus easily spalled off (Fig. 15-a). As a result a triboxidative wear regime cannot be established, as confirmed by the high value of friction from the beginning to the end of the test (Fig. 15-b). Beside MM-6000’ also MMC-80’ shows an apparently anomalous wear behaviour, namely a high wear rate (Fig. 11) in spite its high hardness (Table 1). This is associated to its anisotropic microstructure (Fig. 2-e). The flake like particle are aligned in the direction of the tangential force during sliding favouring an intense delamination wear (Fig. 16).
Fig. 13. Temperature (a) and friction coefficient (b) evolutions of MMC-240’ as function of the applied load for 240 min of sliding test.
Fig. 14. Friction coefficient and temperature ranges as function of load.
3.3. Abrasive wear Fig. 17 the penetration depth as function of applied load is shown for MM-240’ and MMC-240 as representative of the different processing routes, i.e. MM and MA. As reference also AT-Cu and Cu–Be abrasion resistance is reported. It is evident how increasing the load the penetration depth increases for all the materials [18]. For all the loads AT-Cu exhibits the lower abrasion resistance followed by the mechanical milled sample and finally by the mechanical alloyed materials which have
a penetration depth similar to the Cu–Be. Since the test is very short, frictional heating and tribo-oxidation are avoided so that the abrasive wear resistance is basically related to hardness (Fig. 18). It is evident how increasing the hardness the penetration depth decreases, of about 25% in MM samples and to more than 50% in MA ones. This confirms the benefits of strain hardening and, even more, of the combined effect of strain hardening and dispersion hardening, respectively. MMCs show a penetration depth similar to Cu–Be, and this is very satisfying and promising.
Fig. 12. Worn surfaces of AT-Cu (a) and MMC-240’ (b) after 240 min of sliding time at 100 N.
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Fig. 15. Top view of the track surface (a) and friction coefficient evolution (b) of MM-6000’ for 240 min of dry sliding test.
Fig. 16. Delamination phenomenon of MMC-80’.
Fig. 18. Penetration depth of scratch test at 3 N of all the materials as function of hardness.
Fig. 17. Penetration depth of scratch test as function of applied load.
4. Conclusions The aim of this study was to analyse the wear behaviour of copper based material produced by mechanical milling (MM) and mechanical alloying (MA) with TiB2 under dry sliding and abrasive wear conditions.
The dry sliding behaviour against high speed steel (65HRC) is characterized by the transition from an initial adhesive metallic
wear regime into a steady state tribo-oxidative one for prolonged sliding time. At low load (50 N) the wear rate decreases by increasing hardness, in view of the increasing support to the formation of protective oxide layers. At high load (4100 N), the stronger frictional heating causes higher contact temperature and definitely easier tribo-oxidative wear for all materials: accordingly, the influence of higher hardness by MM and MA becomes negligible and all materials perform similar to hard commercial Cu–Be alloy. Under pure abrasive conditions, the positive influence of strain hardening (MM samples) and strainþdispersion hardening (MA samples) becomes evident. The wear resistance of MMC is competitive with Cu–Be alloy. A promising combination of wear resistance and thermal conductivity has been obtained with MMC keeping the thermal conductivity comparable to atomized copper sample (300 W/mK) and much higher than the commercial Cu–Be alloy (106 W/mK).
Acknowledgment The authors gratefully acknowledge Netzsch group for the supply of thermal conductivity analysis, especially Dr. Baldini and Dr. Beckstein. This research was also supported by Pometon Powder Company for the supply of copper powder.
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[21] D.A. Rigney, L.H. Chen, M.G.S. Naylor, A.R. Rosenfield, Wear processes in sliding systems, Wear 100 (1984) 195–219. [22] J. Shell, P. Heilmann, D.A. Rigney, Friction and wear of Cu–Ni alloys, Wear 75 (1982) 205–220. [23] R.S. Montgomery, The sliding behaviors of copper alloys, Wear 87 (1983) 339–349. [24] P. Heilmann, J. Don, T.C. Sun, D.A. Rigney, W.A. Glaeser, Sliding wear and transfer, Wear 91 (1983) 171–190. [25] A. Onat, Mechanical and dry sliding wear properties of silicon carbide particulate reinforced aluminum-copper alloy matrix composites produced by direct squeeze casting method, J. Alloy. Compd. 489 (2010) 119–124. [26] G. Straffelini, L. Maines, M. Pellizzari, P. Scardi, Dry sliding wear of Cu-Be alloys, Wear 259 (2005) 506–511. [27] Y.S. Zhang, Z. Han, K. Lu, Fretting wear behavior of nanocrystalline surface layer of copper under dry condition, Wear 265 (2008) 396–401. [28] M.D. Sexton, A study of wear in Cu–Fe systems, Wear 94 (1984) 275–294. [29] F.H. Stott, The role of oxidation in the wear of alloys, Tribol. Int. 31 (1998) 61–71. [30] Y.S. Zhang, K. Wang, Z. Han, G. Liu, Dry sliding wear behavior of copper with nano-scaled twins, Wear 262 (2007) 1463–1470. [31] J. Lu, S. Shu, F. Qiu, Y. Wang, Q. Jiang, Compression properties and abrasive wear behavior of high volume fraction TiCx-TiB2/Cu composites fabricated by combustion synthesis and hot press consolidation, Mater. Des. 40 (2012) 157–162. [32] J.P. Tu, W. Rong, S.Y. Guo, Y.Z. Yang, Dry sliding behavior of in situ Cu-TiB2 nanocomposites against medium carbon steel, Wear 255 (2003) 832–835. [33] D.H. Kwon, T.D. Nguyen, K.X. Huynh, et al., Mechanical, electrical and wear properties of Cu-TiB2 nanocomposites fabricated by MA-SHS and SPS, J. Ceram. Process. Res. 7 (2006) 275–279. [34] S.C. Tjong, K.C. Lau, Abrasive wear behavior of TiB2 particle-reinforced copper matrix composites, Mater. Sci. Eng. A. 282 (2000) 183–186. [35] M. Tokita, Mechanism of spark plasma sintering, J. Soc. Powder Technol. 30 (1993) 790–804. [36] A. Elsayed, W. Li, O.A. El Kady, W.M. Daoush, E.A. Olevsky, Experimental investigations on the synthesis of W-Cu nanocomposite through spark plasma sintering, J. Alloy. Compd. 639 (2015) 373–380. [37] G.Petzow, MetallographicEtching, Editor by ASM International, 1999. [38] S. Diouf, C. Menapace, A. Molinari, Study of the effect of particle size on densification of copper during spark plasma sintering, Powder Metall. 55 (2012) 3–7. [39] V. Rajkovic, D. Bozic, M.T. Jovanovic, Properties of copper matrix reinforced with various size and amount of Al2O3 particles, J. Mater. Process. Technol. 200 (2008) 106–114. [40] C. Menapace, G. Cipolloni, M. Hebda, G. Ischia, Spark plasma sintering behaviour of copper powders having different particle sizes and oxygen contents, Powder Technol. 291 (2016) 170–177. [41] G. Cipolloni, M. Pellizzari, A. Molinari, M. Hebda, M. Zadra, Contamination during high energy milling of atomized copper powder and its effects on Spark Plasma Sintering, Powder Technol. 275 (2015) 51–59. [42] G. Ondracek, B. Shulz, The porosity dependence of the thermal conductivity for nuclear fuels, J. Nucl. Mater. 46 (1973) 253–258.
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Wear journal homepage: www.elsevier.com/locate/wear
Tribological behaviour of Cu based materials produced by mechanical milling/alloying and spark plasma sintering Massimo Pellizzari, Giulia Cipolloni n Department of Industrial Engineering of the University of Trento, Via Sommarive 9, 38123 Trento, Italy
art ic l e i nf o
a b s t r a c t
Article history: Received 14 September 2016 Received in revised form 4 November 2016 Accepted 25 November 2016
Aim of the work is to investigate the tribological behaviour of Cu based materials produced by mechanical milling (MM), mechanical alloying (MA) and finally sintered by Spark Plasma Sintering (SPS). MM causes a significant strain hardening of Cu, while MA by 0.5 wt%TiB2 adds a further dispersion hardening contribution. During dry sliding contact against high speed steel a close relation has been found between friction coefficient and contact temperature highlighting two regimes. The first one, occurring at an early stage of the process, is characterized by a high friction coefficient ( 1), typical of the strong adhesion of Cu–Cu contact. The second one is characterized by a lower steady-state friction ( 0.7) related to the change of wear mechanism from adhesive metallic into triboxidative. By increasing the load the friction coefficient decreases, the transition time from the adhesive to triboxidative friction regime decreases and the contact temperature increases. Under abrasive wear conditions the penetration depth of the indenter is progressively reduced by increasing hardness, highlighting the benefits of MM and, even more, of MA. The Cu þ 0.5 wt.%TiB2 composite alloyed for 240 min shows the best wear resistance under both wear conditions. The good wear resistance combined to high thermal conductivity represent very attractive properties for many applications. & 2017 Elsevier B.V. All rights reserved.
Keywords: Copper Mechanical milling Dry sliding Adhesion Tribo-oxidation Abrasion
1. Introduction In the last years the demand of materials combining high wear resistance and high thermal conductivity is continuously increasing, especially for components as plastic injection moulds, actively cooled component, continuous casting moulds, electrical contacts and welding electrodes [1–4]. All these applications demanding rapid heat removal from the surface region, require a high heat transfer capability to achieve faster processing routes and, consequently, lower costs. In this way, Cu seems to be a suitable candidate due to its high intrinsic thermal conductivity. By the way the low hardness, strength and wear resistance of Cu restrict its performance and, definitely, its applicability. Therefore, different solutions have been evaluated to improve strength without significantly affecting conductivity. In the last years, severe plastic deformation (SPD) by mechanical milling (MM) has been proposed as an efficient method to improve the strength of Cu by grain refinement and strain hardening [5–8]. An even more extended strengthening can be obtained by uniform dispersion of a harder second phase in the copper matrix by mechanical alloying (MA) [6–8]. This leads to the production of metal matrix n
Corresponding author. E-mail addresses: [email protected] (M. Pellizzari), [email protected] (G. Cipolloni). http://dx.doi.org/10.1016/j.wear.2016.11.050 0043-1648/& 2017 Elsevier B.V. All rights reserved.
composites (MMC), which have attracted growing attention [6]. The success of these materials is related to the possibility to tailor their properties using the synergistic effect of its constituents. The tribological behaviour of MMC's is associated to the different role of the metallic matrix and the reinforcement. Nevertheless it is strictly correlated to the specific testing conditions [9– 11]. Under abrasive conditions, the metal matrix must incorporate and provide suited support to the hard phase so to avoid its premature pull out or fracture. In order to guarantee a suited protection, the hardness of hard particles (HP) has to be higher than that of the abrasive medium (HA) and should also have a high fracture toughness to prevent fracture [12–15]. Usually the wear resistance of materials in abrasion test is improved by changing the abrasive particle movement from gliding to rolling. This is achieved by increasing either the hardness ratio between matrix and HA, or by increasing HP/HA [13]. Also the nature of interface between the constituents is a critical and essential feature for the performance of a MMC because it affects the ability of the matrix to hold the reinforcing particle and prevents its pull out under abrasive wear conditions. The extension of interface is strictly related to the particle size of the reinforcement because it determines the fraction of bonding area between the constituents and consequently the load bearing capability of interfaces. Usually the HP must be bigger than the abrasive particles, this reduces the
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abrasion depth if they are well bonded to the matrix [9]. Moreover the mean free path between the HP must be smaller than or at least as large as the width of the abrasive particle, if possible to act as an obstacle to the scratching abrasive [9,16]. If the HP are too small they will be easily pulled out unless they are present in large quantities obstructing the penetration of the abrasive grain into the softer matrix. By the way it must be considered that over a certain volume fraction percolation of HP may occurs forming brittle network leading to a premature failure of the component [9,17]. Under incremental load scratch testing conditions, four microscopic mechanisms and processes of scratch deformation have been observed, e.g. elastoplastic grooving, plastic plowing, microcracking, and the inter and intra-fractures and chipping of the surface scratched [18]. By increasing the amount of second phase the magnitude of damage is reduced [18]. Under dry sliding conditions the materials performance of MMC is not only related to the hardness of the material but also to the capability to activate self-protective mechanisms depending on the variable service conditions. For copper based materials the wear behaviour against steel is initially characterized by the surface plastic deformation at the contacting asperities favoured by a shear instability effect [19–23]. This effect gives rise to the occurrence of intense transfer phenomena with the formation of a mechanically mixed layer on the surface of both wearing materials [23–25]. If the interaction with the environment and the service conditions promote oxidation, a tribo-oxidative wear mechanism governs the wear behaviour. The formation of a protective oxide tribo-film improves the wear behaviour reducing the friction coefficient [26–30]. Under these conditions, a harder material provides higher load bearing capability to the oxide layer improving wear resistance [29]. In general for higher hardness the plastic deformation is resisted and material removal is more difficult, consequently the sliding marks on the surface appear shallower. Moreover the higher hardness can also help to restrain effectively nucleation or propagation of the crack generated during testing [30]. By the way the formation of groove on the worn surface cannot be completely avoided during sliding and the sub surface cracks along the sliding direction usually promote the detachment of wear particles in the form of sheet and flake leading to a delamination process [25]. In general by increasing the content of reinforcement the wear resistance increases independently from the type of reinforcement [17,31,32]. In Cu–TiB2 composite has been demonstrated that increasing the content of hard phase, wear resistance increased considerably if the distribution of reinforcement in the matrix is uniform [33,34]. Occasional fracture at the reinforcement-matrix interface was shown to occur under high load conditions, especially for large reinforcement particle size (445 mm) [34]. In this paper copper based materials have been produced by MM, MA and Spark Plasma Sintering (SPS) to increase the wear resistance by strain hardening and dispersion hardening keeping acceptable value of thermal conductivity. SPS allows sintering at lower temperature and in a shorter time comparing to more conventional processes. Indeed the high heating rate peculiar of SPS preserves the fine microstructure produced by MM and MA, and reduces any possible interaction between the metal matrix and the reinforcing particles [35,36]. Aim of the investigation is to give a contribution to the understanding of the dry sliding and abrasion wear behaviour of copper based materials and, in particular, to highlight the influence of the load and the role of the material hardness and microstructure obtained by the MM and MA. To this end, a first series of dry sliding tests against high speed steel (AISI M3:2) were carried out at different loads. Special emphasis was given to the role of microstructure in tribolayer formation. Moreover, the tribological behaviour of mechanical milled and mechanical alloyed
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materials has been also studied, under pure abrasive conditions by means of the scratch test. The results are expected to provide important criteria in designing novel Cu based tribomaterials.
2. Materials and experimental methods 2.1. Materials A water atomized copper powder (AT-Cu) with a particle size lower than 75 mm was considered for MM. For the production of MMCs 0.5 wt% of TiB2 with particle size finer than 3 mm was added to AT-Cu by MA. Both processes were carried out using a Fritsch Pulverisette 6 planetary mono mill at 400 rpm. Vial and sphere with 10 mm diameter of 100Cr6 (63 HRC) were used. The system was evacuated down to 130 Pa and 0.5 wt% of stearic acid was added as process control agent. MM was carried up to 6000 min in order to obtain different final hardness. MA up to maximum 240 min, time necessary to achieve a good dispersion of HP in the s Cu matrix. The milled powders were sintered in a DR.SINTER SPS1050 (Sumitomo Coal & Mining, now SPS Syntex, Inc.) apparatus with graphite punches and dies. SPS was performed at a nominal temperature of 950 °C, with a uniaxial pressure of 60 MPa applied at 700 °C. The heating rate was 100 °C/min up to 900 °C and 50 °C/min up to the sintering temperature. The maximum temperature and pressure were held for 5 min before allowing the furnace to cool to room temperature. The properties of MM and MA samples were compared to those of AT-Cu and a commercial Cu–Be alloy (Moldmax HH). 2.2. Experimental methods The density of the SPS-processed pellets was measured according to Archimedes' principle (ASTM B962-08). For mechanical milled samples the relative density was calculated considering the theoretical densities of Cu as 8.96 g/cm3. In the case of MMC the theoretical absolute density was calculate according to the linear rule of mixture and considering the theoretical density of TiB2 as 7.76 g/cm3. Sintered samples were cut into disk-shaped specimens of diameter 10 2.7 mm and the thermal conductivity (k) of the samples was derived by the equation: k ¼ α ρ Cp. Where α is thermal diffusivity, ρ is the bulk density, and Cp is specific heat capacity. The thermal diffusivity and the specific heat were measured using the NETZSCH laser flash apparatus LFA 467 HyperFlash. All samples were tested with the LFA between 400 and 500 °C with 20 °C temperature steps. Prior to the measurement the front and the back of the samples were coated with graphite to enhance the emission/absorption properties of the sample. The specific heat was determined by the reference method given by ASTM-E 1461– 2011. The LFA was calibrated with a Cp-standard of pure-copper. The sintered pellets were machined into block samples with the dimension of 9x13 5 mm. The 9 5 mm cross section faces were cold molded in epoxy resin for metallography. Standard metallographic preparation, including grinding with SiC papers up to 4000grit, final polishing with 3 μm and 1 μm diamond pastes, was performed. A chemical etching with 120 mL of distilled water, 30 mL of hydrochloric acid and 10 g of iron chloride was used [37]. Vickers micro-hardness was measured using an applied load of 1 N, holding time of 10 s and a loading rate of 0.1 N/s according to ASTM standard E9-08. The 9 13 mm faces of the machined block were employed as the wear surface both for sliding and abrasion wear tests (Fig. 1a). The wear surface have been polished up to 1 micron diamond paste. The dry sliding tests were carried using a block-on-disc configuration with a rotating speed of 300 rpm, corresponding to a
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Fig. 1. Sample block (a) and sliding wear configuration (b).
sliding speed of 0,63 m/s (Fig. 1b). As counterpart a disc of AISI M3:2 (830HV30) has been used. The tests have been carried out at 50 N for 15, 30 and 240 min, and for 100 and 200 N for 30 and 240 min. The worn abrasive disc and the copper based samples were substituted by a new one for every sliding time and load. During the experiments, the friction coefficient was recorded and the contact temperature was continuously measured by a thermocouple placed in a hole at 2 mm depth from the contact surface (Fig. 1b). The wear curve was recorded by periodically weighing the specimens using a precision balance (0.0001 g). The wear rate (Ka) has been calculated as ratio between volume loss and sliding distance. The wear mechanism has been analysed by light optical and scanning electron microscopy of the wear tracks and debris. Chemical analysis was carried out by energy-dispersive X-ray spectroscopy (EDXS). The abrasive wear resistance has been evaluated by the scratch test at constant load, using a Rockwell diamond indenter with 200 μm of diameter. Different loads (1, 2 and 3 N) were applied for an abrasion distance of 5 mm and with a sliding speed of 10 mm/ min. The output of the test is the penetration depth with the accuracy of 71 μm. The press surface, i.e., the sintered plane perpendicular to the pressing direction during SPS was indented/ scratched. Each test surface was polished down to 1 micron diamond paste.
3. Results and discussion 3.1. Materials characterization In Fig. 2 the OM micrographs of all the sintered samples are reported. The sample sintered from AT-Cu powder shows a uniform microstructure (Fig. 2a). Some pores are still evident and give a pinning effect for the grain growth [38]. For MM and MA the microstructure of sintered materials keeps memory of the original particles morphology, a severe anisotropy being observed when flaking prevails for MM-240’ and MMC-80’ (Fig. 2-b and -e). Instead for prolonged milling time, when welding events predominate, the microstructure becomes more uniform e.g. MM720’ and MMC-240’ (Fig. 2-c and -f). The large milled particles are clearly recognizable and they are surrounded by a network of recrystallized copper film (white region indicated by arrows) due to an intense overheating at the contact points between the powders during SPS [39]. In MMC by increasing milling time the dispersion of the hard phase becomes more homogeneous, TiB2 particles are more and more refined and the adhesion between the two
constituents is improved. MM-6000’ shows a uniform microstructure (Fig. 2-d) because the milled powder consists in small agglomerations of very thin flake like particles that are easily compacted during SPS [40]. Table 1 summarizes the properties of the six sintered sample. For MM the best combination of hardness has been obtained after 240 min, with an increase of hardness from 103HV0.1 for ATCu to 150HV0.1 for MM-240’. Despite the more extensive strain hardening induced by longer MM a decrease of hardness is observed for MM-720’ and MM-6000’. This is related to an increase of the residual porosity associated to the incomplete decomposition of the process control agent during SPS which releases gaseous species that cause the formation of pores and reduce the final density (Table 1) [41]. A different trend is shown by MMC's, where the hardness increases with milling time due to the progressive refinement of TiB2 particles and their improved distribution. The hardness of MMC-240’ almost doubles that of AT-Cu sample, increasing from 103 HV0.1 to 207 HV0.1. MMC-80’ and MMC-240’ exhibit density values of 98 and 97% respectively: the increase of hardness increasing milling time confirms that strain hardening and dispersion hardening during MA abundantly compensates for the negative effect of porosity. The thermal conductivity of AT-Cu is (300 W/mK) is slightly lower than that of pure Cu (400 W/mK) due to the deleterious presence of oxide particles and pores. The effects of MM and MA are quite different. On one side, a detrimental effect of MM has been highlighted especially for long milling time. MM-240’ exhibits an unexpected increase of thermal conductivity, meanwhile MM-720’ and MM-6000’ show a drastic decrease of thermal conductivity due to the higher porosity and the presence of crystal defects [42]. On the other side, MA exhibits some interesting and promising results in contrast with what expected. The thermal conductivity of MMC-80’ and MMC-240’ is very close to that of AT -Cu, in spite of the mechanical strain hardening and the addition of TiB2. The small amount and the good dispersion of reinforcement, especially for long milling time, provide an excellent combination of hardness and thermal conductivity. All the materials produced by MM and MA show a higher thermal conductivity than the commercial Cu–Be alloy. 3.2. Dry sliding wear 3.2.1. Friction A typical diagram of the friction coefficient and the contact temperature recorded during the dry sliding test is reported in Fig. 3, where a clear correlation between these two parameters is evident. Two distinct friction regimes can be distinguished: an
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Fig. 2. Optical micrographs of the microstructure of (a) AT-Cu, (b) MM-240’, (c) MM-720’, (d) MM-6000’, (e) MMC-80’ and (f) MMC-240’. Table 1 Properties of the sintered materials. Sample
Milling time (min)
HV 0.1
Rel. Density (%)
Thermal Conductivity (W/mK)
AT-Cu MM-240’ MM-720’ MM-6000’ MMC-80’ MMC-240’ Cu-Be
0 240 720 6000 80 240 –
103 72.1 150 71.2 1107 2.3 457 1.8 130 71.9 207 71.2 400 71.2
99 94 90 81 98 97 100
300 320 270 200 298 297 106
early stage, characterized by a high friction coefficient ( 1) and higher temperature followed by a steady-state regime showing lower friction coefficient ( 0.7) and temperature. The high value
of friction coefficient in the early stage is typical of an adhesive metal-metal contact. The observed fluctuations remark the occurrence of stick-slip phenomena accompanied by material transfer [19–21]. In this stage both temperature and friction increase to show a drop to lower steady state values after a certain transition time (tss). It is well recognized that the frictional heating facilitates surface oxidation, promoting the formation of protective oxide layers. This results in milder wear conditions characterized by metal-oxide, oxide-oxide contacts and in a decrease of the friction coefficient (Fig. 3) [29]. Definitely, after a certain time (tss) a transition from an adhesive to a triboxidative wear mechanism is evidenced [12,17,19,26]. At this stage the oxide layer acts as a solid lubricant reducing the resistance to sliding. The temperature consequently decreases keeping a constant value for the remaining duration of the test. Although all the materials show a similar macroscopic friction
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protective oxide layers. One the other one, harder materials show a more difficult transition from not conformal to conformal contact, so that wear debris are less easily entrapped in the contact region to form a protective oxide layer. This is well evident for MMC-240’ characterized by a hardness two times higher than AT-Cu (Table 1). Fig. 4 also shows that the steady state friction of MM-240’ and MA-240’ are higher than AT-Cu. The plausible reason is the improved third body abrasive contribution of the harder wear particles.
Fig. 3. Friction coefficient and temperature evolutions for 240 min of dry sliding test of MMC-240’.
Fig. 4. Comparison of the friction coefficient evolution for 240 min of dry sliding test of AT-Cu, MM-240’ and MMC-240’.
coefficient evolution under dry sliding condition, some differences can be highlighted between MM and MA (MMC) samples. In Fig. 4 the friction evolutions for 240 min sliding test of AT-Cu, MM-240’ and MMC-240’ are reported. Comparing AT-Cu with MM-240’ and MMC-240’ it can be noticed that increasing the hardness the time required to achieve the triboxidative wear regime (tss) becomes longer. On one side, harder materials produce less debris, the particles required to form
3.2.2. Wear The above statements regarding the friction behaviour and the wear mechanism can be confirmed looking at the changes in structure and chemical composition of sliding surfaces, after different sliding times, before and after tss. For this reason the sliding test were carried out for 15, 30 and 240 min to monitor and analyse the evolution of tribological layer during the transition of the wear behaviour from the adhesive mechanism to the tribo-oxidation mechanism. The characterization of the wear tracks and of the discs have been carried out, and optical micrographs after 15 and 240 min of sliding test are reported respectively in Figs. 5 and 6. After 15 min (see Fig. 3), an accumulation of material is highlighted in the wear track, (Fig. 5-a). The wear surface exhibits a quite rough aspect, with an inlet area characterized by deep abrasion scratches and an outlet area showing an intense accumulation of material. A strong materials transfer from the Cu block to the disc (high speed steel) surface is highlighted (Fig. 5-b) as confirmed by the EDXS analysis, (Fig. 7-a). Large patches of copper are transferred on the surface of the disc and tent to shear in the sliding direction, favoured by the high ductility of the Cu matrix, (Fig. 5-b). The evidence of a Cu–Cu adhesive contact is clearly proved. The high frictional heat during the running stage enhances the generation of oxide and consequently the change of the wear mechanism from adhesion mechanism to tribo oxidation. Moreover the formation of the wear track results in the change of the contact nature from not conformal to conformal. This process leads to the production of wear debris which can be transferred on the disc, i.e. adhesion, either can be packed inside the track or pulled out of the wear system. The creation of the wear track aids firstly the permanence of the wear debris in the tribological system and then their oxidation. The optical micrographs after 240 min (Fig. 6-a) of sliding confirm the change of the wear mechanism in accordance with Fig. 3. In Fig. 6-a the wear tracks is larger and cleaner than in Fig. 5-a, in addition any massive accumulation of material is detected in the outlet area and scratches due to the abrasion action
Fig. 5. Top view of wear track (a) and disc surface (b) after 15 min of dry sliding test of MMC-240’.
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Fig. 6. Top view of wear track (a) and disc surface (b) after 240 min of dry sliding test of MMC-240’.
are vanished in the inlet region. The oxide layer in the track is much thinner and homogeneous, except for some darker stripes on the left side of the track (Fig. 6-a) Also the surface of the disc drastically changes, the transfer of material is drastically decreased and a sever oxidation is evident (Fig. 7) as also confirmed by the EDXS analysis (Fig. 8). Since copper oxide is more brittle than Cu, it breaks more easily and the fragments create more quickly debris and with that the protective oxide. The brittleness of the copper oxide is also confirmed by the finer dimension of the compacted and transferred particles on the disc after 240 min in comparison to the stretched and spread patches after only 15 min, (Fig. 5-b and Fig. 6-b). The continuous formation and fragmentation of oxide layer during prolonged sliding wear consists in a dynamic phenomenon which contributes to the wear damage and to the production of wear debris. Good evidenced has been provided to show that the oxide formed by frictional heat grows until, at a critical oxide film thickness, about 10micron, it spalls off as wear debris. Once fragmentation of the oxide layer occurs part of the loose debris particle may become entrapped, fragmented and compacted in the contacting surfaces contributing to the protective action of the oxide layer (Fig. 9) [29]. Some other debris, instead, are pulled out of the wear system. The typical morphology of debris of MMC-240’ collected after 240 min is reported in Fig. 10. Large plate-like oxide fragments (Ox in Fig. 10) are present, confirming the occurrence of delamination due to the large shear strain during sliding. The dimension of the plate-like debris (4200 μm) are in accordance with the size of the oxide patches on the wear track (Fig. 9). Some debris show a shiny surface due to the intrinsic
metallic nature (M) and probably they correspond to the side attached to the wear track surface before spalling. It is also plausible that these metal debris are those produced by delamination during the early stage, as the adhesive mechanism predominates. On the other hand, the finer fraction of debris corresponds to fragmented oxide particles. 3.2.3. Wear resistance In Fig. 11 the wear rate for all the samples tested for 240 min at 50 N are reported. The dependence of wear rate on hardness is very clear, increasing hardness the wear coefficient decreases. By the way the wear rates of material produced by MM and MA are higher than Cu–Be and similar to AT-Cu, except for MM-6000’ and MM-720’. For hardness value lower than 100 Hv the soft substrate limits the load bearing capability of the material (MM-6000’) to sustain the oxide layer leading to a higher wear rate. Also the lower thermal conductivity of MM-6000’ ( 200 W/mK) (Table 1) may negatively affect the wear behaviour because, if the heat is not properly removed, the material undergoes softening and its load bearing capability is further decreased. When hardness exceed 100HV0.1 the permanence of the oxide layer is promoted and lower value of wear rate are guaranteed. Only if the oxide is firmly attached to the sliding surface it can get compacted protecting the underlying base alloy. The presence of a mixed layer in the wear track will influence its hardness and other mechanical properties, which in turn will influence friction and wear (Figs. 4 and 11). It must be considered that a higher hardness leads to the formation of harder debris, consequently the protective oxide layer will be harder providing a lower wear rate. For these materials the thermal
Fig. 7. SEM micrographs of the copper transferred particle on the disc surface after 15 min (a) and 240 min (b) of dry sliding test.
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Fig. 8. EDXS spectra of the copper transferred particle on the disc surface after 15 min (a) and 240 min (b) of dry sliding test.
conductivity higher ( 300 W/mK) than MM-6000’ aids the dissipation of the heat avoiding softening. In spite of the relatively low hardness MM-720’ shows an exceptionally low wear rate (Fig. 11), even lower than Cu–Be. The investigations carried out until now could not clarify the reason for this results. Detailed microstructural analysis are ongoing. With exception of MM-6000’ and MMC-80, the wear coefficients decrease as the load is increased to 100 N (Fig. 11). It is evident that increasing the load the positive influence of hardness becomes less important, so that all materials perform similarly to Cu–Be alloy. This is very clear for soft AT-Cu which is supposed to have a low wear resistance due to its low hardness. By the OM observation of the worn surfaces of AT-Cu an intense accumulation of wear debris inside the wear track is evident in Fig. 12-a. The wear products are stuck and compacted inside the track, thus altered the wear loss measurement and consequently the calculation of the wear rate. This is important for the performance of the material because such material accumulations inside, for example, an injection mold can lead to dimensional changes not tolerated by the design. All the others materials show a similar morphology of the wear surface, with the presence of an oxide protective layer but thinner than that in the case of AT-Cu, Fig. 12-b. In Fig. 12-b MMC-240’ exhibits a wear track smaller than for 50 N test (Fig. 6-a), but this is perfectly in accordance with the lower wear rate depicted in Fig. 11. 3.2.4. Effect of the sliding load The effect of the increasing load results in an increased contact temperature and a decreasing friction coefficient (Fig. 13). According to the mechanism explained above, a higher temperature facilitates the generation of oxide debris and assists their
Fig. 10. The morphology of the wear debris of MMC-240’ after 240 min of sliding: the metallic nature (M) and the fine oxide powder (Ox).
compaction for the formation of the wear protective oxide layers [10,11,25,27,30]. It must be considered that during the steady-state regime two competitive processes occur: the break-down of the protective layer, resulting in wear, and the compaction/sintering of wear debris, to reform the oxide layer [29]. The latter process, as the oxidation of the particle, occur more rapidly as the temperature is increased and friction decreases (Fig. 13-b) [12]. The transition from adhesion to tribo-oxidation shifts towards shorter time so that the running in almost disappears at 200 N (Fig. 13-b). Moreover, the friction becomes more stable, stick and slip phenomena are less frequent and adhesion is drastically limited. All these phenomena positively affects the wear rate which decreases
Fig. 9. Detailed view of the worn surface of MMC-240’after 240 min of sliding (a) showing the formation of protective Cu oxide patches (EDXS is referring to the darker region in Fig.a highlighted by the white cross) (b).
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Fig. 11. Dependence of wear coefficient on hardness for long sliding test.
by increasing the load (Fig. 11). All the materials show the same behaviour by increasing the load, as reported in Fig. 14. The lowest and the highest values of friction coefficient at 50 N corresponds to AT-Cu and MM-6000’ respectively, in all the other cases the values are included in a similar range. In Fig. 11 it can be noticed that MM-6000’ and MMC-80’ show the highest wear rate both at 50 and 100 N. In the case of MM6000’ this can be ascribed to the low hardness, i.e. to the insufficient support by the base material to the oxide layer, which is thus easily spalled off (Fig. 15-a). As a result a triboxidative wear regime cannot be established, as confirmed by the high value of friction from the beginning to the end of the test (Fig. 15-b). Beside MM-6000’ also MMC-80’ shows an apparently anomalous wear behaviour, namely a high wear rate (Fig. 11) in spite its high hardness (Table 1). This is associated to its anisotropic microstructure (Fig. 2-e). The flake like particle are aligned in the direction of the tangential force during sliding favouring an intense delamination wear (Fig. 16).
Fig. 13. Temperature (a) and friction coefficient (b) evolutions of MMC-240’ as function of the applied load for 240 min of sliding test.
Fig. 14. Friction coefficient and temperature ranges as function of load.
3.3. Abrasive wear Fig. 17 the penetration depth as function of applied load is shown for MM-240’ and MMC-240 as representative of the different processing routes, i.e. MM and MA. As reference also AT-Cu and Cu–Be abrasion resistance is reported. It is evident how increasing the load the penetration depth increases for all the materials [18]. For all the loads AT-Cu exhibits the lower abrasion resistance followed by the mechanical milled sample and finally by the mechanical alloyed materials which have
a penetration depth similar to the Cu–Be. Since the test is very short, frictional heating and tribo-oxidation are avoided so that the abrasive wear resistance is basically related to hardness (Fig. 18). It is evident how increasing the hardness the penetration depth decreases, of about 25% in MM samples and to more than 50% in MA ones. This confirms the benefits of strain hardening and, even more, of the combined effect of strain hardening and dispersion hardening, respectively. MMCs show a penetration depth similar to Cu–Be, and this is very satisfying and promising.
Fig. 12. Worn surfaces of AT-Cu (a) and MMC-240’ (b) after 240 min of sliding time at 100 N.
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Fig. 15. Top view of the track surface (a) and friction coefficient evolution (b) of MM-6000’ for 240 min of dry sliding test.
Fig. 16. Delamination phenomenon of MMC-80’.
Fig. 18. Penetration depth of scratch test at 3 N of all the materials as function of hardness.
Fig. 17. Penetration depth of scratch test as function of applied load.
4. Conclusions The aim of this study was to analyse the wear behaviour of copper based material produced by mechanical milling (MM) and mechanical alloying (MA) with TiB2 under dry sliding and abrasive wear conditions.
The dry sliding behaviour against high speed steel (65HRC) is characterized by the transition from an initial adhesive metallic
wear regime into a steady state tribo-oxidative one for prolonged sliding time. At low load (50 N) the wear rate decreases by increasing hardness, in view of the increasing support to the formation of protective oxide layers. At high load (4100 N), the stronger frictional heating causes higher contact temperature and definitely easier tribo-oxidative wear for all materials: accordingly, the influence of higher hardness by MM and MA becomes negligible and all materials perform similar to hard commercial Cu–Be alloy. Under pure abrasive conditions, the positive influence of strain hardening (MM samples) and strainþdispersion hardening (MA samples) becomes evident. The wear resistance of MMC is competitive with Cu–Be alloy. A promising combination of wear resistance and thermal conductivity has been obtained with MMC keeping the thermal conductivity comparable to atomized copper sample (300 W/mK) and much higher than the commercial Cu–Be alloy (106 W/mK).
Acknowledgment The authors gratefully acknowledge Netzsch group for the supply of thermal conductivity analysis, especially Dr. Baldini and Dr. Beckstein. This research was also supported by Pometon Powder Company for the supply of copper powder.
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