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Vol. 8. Issue 1.
Pages 436-446 (January - March 2019)
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Vol. 8. Issue 1.
Pages 436-446 (January - March 2019)
Original Article
DOI: 10.1016/j.jmrt.2018.04.004
Open Access
Mechanical and tribological characterization of AlCrN coated spark plasma sintered W–25%Re–Hfc composite material for FSW tool application
Akeem Yusuf Adesinaa,
Corresponding author

Corresponding author.
, Zafar Iqbalb, Fadi A. Al-Badoura, Zuhair M. Gasema
a Department of Mechanical Engineering, King Fahd University of Petroleum and Minerals, Dhahran 31261, Saudi Arabia
b Centre of Research Excellence in Corrosion, King Fahd University of Petroleum and Minerals, Dhahran 31261, Saudi Arabia
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Figures (12)
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Tables (4)
Table 1. Deposition parameter for all coatings.
Table 2. Experimental parameters for the wear test.
Table 3. Composition and surface properties of the AlCrN coating.
Table 4. Mechanical properties of AlCrN coated sample.
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In this study, the mechanical and tribological properties of cathodic arc physical vapor deposited AlCrN coating on spark plasma sintered W–25%Re–HfC composite tool material were investigated. AlCrN coated and uncoated W–25%Re–HfC samples were tested using pin-on-disk wear test configuration to evaluate the coating performance. Scratch test result shows that the adhesion strength of the coating is about 25N, which indicates that the coating exhibited good adhesion to the composite material. Specific wear rate of the coated sample is 10 times lower than that of the uncoated sample under identical conditions. The high wear rate of the uncoated W–25%Re–HfC sample is due to extensive abrasive wear. However, the coated sample is dominated by oxidation wear mechanism leading to the formation of dense Al2O3 and Cr2O3 oxides with good wear resistance properties. The improved wear resistance of the coating is attributable to the combined excellent mechanical properties, high adhesion to the substrate, low coefficient of friction and the formation of protective oxides. This study demonstrates by way of tribological analysis the feasibility of improving the life and performance of expensive FSW tools by the application of cathodic arc AlCrN PVD coating.

Friction stir welding
PVD AlCrN coating
FSW tool wear
W–Re alloy
Full Text

Friction stir welding (FSW) has developed as a novel technology for obtaining improved quality welding of high strength and high softening temperature (HS-HST) materials such as steels, titanium alloys, etc. However, the severe plastic deformation coupled with high temperature and pressure encountered during the process subject the tool to extreme wear leading to weld contamination and tool failure. Thus, the tool is critical to FSW process, and as a result several tool materials such as pcBN, WC-base alloy, molybdenum-base alloy, W–Re base alloy/composite, etc., are being developed to abate the problem of tool wear during FSW of HS-HST materials [1–9]. Nevertheless, tool wear still remains a problem and a major setback [4,10,11]. Besides other well-known wear mechanisms such as abrasion and adhesion, tool wear may also induce chemical reaction and subsequent formation of second phases which may lead to the depletion of an essential element from the weld joint as a result cause weld quality deterioration. Park et al. [12] showed that the wear of pcBN tool caused the depletion of Cr and the formation of secondary phases consisting of Cr2Br and Cr5Br3 in the stir zone during FSW of stainless steel. Wang et al. [13] also reported severe deformation and adhesive wear of tool when W-1.1%La2O3 and WC-Co tools were used in the FSW of Ti–6Al–4V alloy. In general, the expensive cost and processing required for preparing these materials are tremendous.

In recent times, the use of metallurgically hard physical vapor coatings (PVD) has been identified as a promising and more economical solution to the wear of FSW tools as experienced in the cutting tool industries. In a preliminary study, AlCrN coated FSW tool was employed in welding 6mm thick 6061-T6 aluminum alloy plate, it was found that the coated FSW tool produces defect-free and full penetration weld without failure, delamination of coating or diffusion of coating material into the weld [14]. It was also shown in the study that the AlCrN coated FSW tool possesses enhanced resistance to deformation and wear due to the improved surface properties and up to 55% reduction in wear was achieved when compared to uncoated FSW tool [14]. Also, Ohashi et al. [15] illustrated the effect of Si3N4 tool wear in contaminating the steel weld with Si and N during Friction Stir Spot Welding (FSSW). This contamination was avoided with the use of TiC/TiN coated tool. Also, TiAlN coated high-speed steel was used in welding A2124/25%SiC Al MMC plates so as to prevent the tool from wear [16].

Due to the significant wear of WC base tools in FSW of HS-HST, some studies have been carried out to abate its wear using PVD coatings. WC-6%Co FSW tool was coated with a 5μm thick AlCrN PVD coating for FSW of Ti6Al4V alloy by Batalha et al. [17]. However, no significant inference can be drawn regarding the contribution of the coating because the tool pin fractured during the plunging stage. This was due to the severe process parameters used leading to excessive plunging force that exceeded the tool compressive strength. In addition, they reported significant adhesion of the coating material to the workpiece and thus, the coating was completely delaminated from the substrate.

Composites such as, W–Re based reinforced with HfC have also been developed [18–20] for FSW tool materials, which possessed better machinability and high fracture toughness as compared to common tools such as W–C, W–Re, and pcBN. However, considering the high material cost and processing involved in developing W–Re–HfC composites, tool wear cannot be tolerated. Thus, the use of PVD coating can be employed to extend the life and improve the performance of such expensive tool material.

In the present study, the mechanical and tribological properties of AlCrN physical vapor deposited (PVD) coating on spark plasma sintered W–25Re–HfC composite material is investigated for possible application to FSW tools for welding HS-HST materials such as stainless steels and titanium alloys. AlCrN coated and uncoated W–25%Re–HfC samples were comparatively tested using pin-on-disk wear test configuration under a high-load bearing condition to evaluate the performance of the coating. Results from the scratch test, structural and morphological analysis of the coating were also presented and discussed.

2Experimental details2.1Tool material development

The tool material was developed from a semi-alloyed tungsten-rhenium alloy (W–25wt.% Re) and hafnium carbide (HfC) powders (received in the semi-alloyed form from Rhenium Alloys Inc., USA) via mechanical alloying (MA) technique to obtain nanostructured particle sized powders with a uniform distribution of the reinforcement. During the mechanical alloying process, 10vol.% of HfC powder was dispersed as reinforcement in the nanocrystalline W–25% Re alloy. Spark plasma sintering (SPS) was utilized for the consolidation of the composite powder material. Details of the synthesis and consolidation procedures of the W–25%Re–10vol.% HfC FSW composite tool material is similar to the procedure in an earlier published work [21]. The main MA parameters used for the tool material in the study are milling rotational speed of 150rpm for 15h milling time with a ball-to-powder ratio of 5:1. The SPS consolidation parameters are 50MPa compaction pressure at a temperature of 1800°C for a duration of 10min. The mechanical, structural and morphological properties of the developed composite material have been presented elsewhere [21].

2.2Coating deposition and characterization

The cathodic arc aluminum chromium nitride (AlCrN) PVD coating was deposited using an industrial size cathodic arc PVD machine (Metaplas.Domino Mini, Oerlikon, Germany). The machine is equipped with two evaporators carrying the AlCr targets for depositing AlCrN coating. The target has an elemental composition of 68at.% Al, 29.5at.% Cr, 1.5at.% Mg and 1at.% Si. Magnesium (Mg) and silicon (Si) are to enhance the target fluidity. The W–25%Re–10vol.% HfC composite substrate was ground, polished and cleaned with acetone before it was mounted on the hexagonal cylindrical holder with two axes of rotation and then positioned on the planetary system of the PVD machine. The W–25%Re–10vol.% HfC substrate sample was fixed on the cylindrical holder with the help of a screw-adaptable sample holder. The substrate roughness before the coating was deposited, expressed in the arithmetic mean (Ra) and root-mean-square (RMS) values, was 0.25±0.04μm and 0.35±0.07μm, respectively.

The coating deposition procedure consisted of five stages: (i) Vacuum evacuation to low pressure and testing to ensure chamber is not leaking. (ii) Arc enhanced glow discharge (AEGD) etching, and sample cleaning. (iii) Heating the chamber to the deposition temperature and evacuating to the desired pressure for the deposition process. (iv) Coating deposition process. Finally, (v) cooling stage to avoid oxidation of the coating. Table 1 summarizes the details of the coating deposition parameters used in the current work.

Table 1.

Deposition parameter for all coatings.

Coating parameters  AlCrN coating on W–25%Re–HfC 
AEGD etching duration  30min 
Nitrogen gas flow rate  500sccm (99.996% purity) 
Planetary rotating speed  2rpm 
Cathode charge  400Ah 
Deposition temperature  450±20°C 
Deposition pressure  8.5×10−2–6.5×10−2mbar 
Bias voltage  Between −80 and −100
Deposition rate  ∼3.5μm/h 
Coating thickness  4.5±0.5μm 

Morphological and compositional characterization of coated samples were conducted by scanning electron microscope (SEM) (JOEL, Japan) and energy dispersive spectroscopy (EDS) which is attached to the SEM, respectively. X-ray diffractometer (D8 advance XRD, Bruker, USA) with CuKα radiation (λ=0.154186nm) and optical profilometer (ContourGT-K, Bruker, Germany) were used for the structural analysis and for measuring the surface roughness respectively.

2.3Mechanical and tribological properties

Microindentation machine (MicroCombi Tester, CSM Instruments, Switzerland) was used to measure the mechanical properties and scratch resistance or adhesion of the coating to the tool material. The indentation hardness (H) and elastic modulus (E) were measured according to Oliver and Pharr model [22]. H is a measure of the resistance to permanent deformation or damage and it is determined according to Eq. (1)[22] while E was calculated from the plane strain modulus E* using an estimated sample poison ratio according to Eqs. (2)–(4)[22].

where Fmax is the maximum test load, Ap is the projected contact area, Er is the reduced modulus, S is the contact stiffness at Fmax, β is indenter geometry constant (β=0.0134 for triangular indenter) [23], and vi is poison ratio of the indenter (0.07), Ei is the elastic modulus of the indenter (1141GPa) and vs is the Poison ratio of the coated specimen (Poison's ratio of 0.3 is typical for ceramic coating) [24,25]. The indentation test was conducted for the coated sample with a normal load of 20mN applied for 10s dwell time at loading and unloading rates of 40mN/min. The load was appropriate since the maximum penetration depth is below 10% of the coating thickness. A normal load of 2N was utilized for the indentation test of the uncoated W–25%Re–HfC sample while other parameters remain the same.

The adhesion of the coating to the composite tool material was investigated through the scratch test with a Rockwell C indenter of 100μm tip radius. The indenter was pressed against the coatings with an initial applied load of 30mN. Then, it was pulled across the coating surface with progressively increasing load at a rate of 0.1N/s until the maximum applied load of 30N. The scratch test was conducted over a scanning length of 10mm with a scratch speed of 2mm/min. During the test, the normal load, penetration depth, acoustic emission (AE), the frictional force () and the coefficient of friction (COF) were continuously measured and recorded. By observing changes in and/or AE signals, the cohesive failure (Lc1) and adhesive failure (Lc2) were determined.

The wear test was conducted on a multipurpose tribometer (UMT-3, Bruker, Germany) with a ball-on-disk configuration. Hardened steel counterpart of 6.3mm diameter and HRC 62 (∼750HV) hardness was used for the test. The Young's modulus and Poisson's ratio of 207GPa and 0.28 respectively are typical for hardened steel ball, thus these values were used in calculating the Hertzian contact pressure. The tests were performed in ambient atmospheric conditions of 23±2°C temperature and 40±5% relative humidity. The experimental parameters for the wear test are shown in Table 2. The wear behavior of the AlCrN coated and uncoated W–25%Re–HfC composite tool material was comparatively evaluated. The specific wear rate was obtained by measuring the net weight using a high-resolution analytical weight balance (AUW220D, Shimadzu Analytical Balance, Japan). The resolution and repeatability of the weight balance are 0.01mg and ≤0.1mg, respectively. The specific wear rate is expressed as mass loss per unit sliding distance per applied load (mg/kNm) accordingly.

Table 2.

Experimental parameters for the wear test.

Load (N)  Rotational speed (rpm)  Track diameter (mm)  Sliding distance (m)  Sliding speed (m/s) 
30  800  12  500  0.50 
3Results and discussion3.1Characterization of coated sample

The surface morphology of the uncoated W–25%Re–HfC sample, as-deposited coated sample and the cross section of the coating showing the thickness are shown in Fig. 1. Microparticles and pore defects which are typical of cathodic arc PVD coatings can be observed from the surface morphology. These defects are one of the main setbacks of PVD coatings and are unwanted, as they affect the surface and corrosion properties of the coating such as causing sticking to work-piece, localized coating failure and pitting corrosion [26]. However, modern development in the PVD machine has significantly enabled the reduction in the droplets population and size. As also shown in the authors earlier study [14], the droplets of the coating are mostly less than 1μm, which is very small as compared to up to 40μm droplet size, which is considered typical [27,28].

Fig. 1.

Surface morphology of the (a) W–Re–HfC sample, (b) as-deposited AlCrN coating and (c) cross-section image showing the coating thickness.


The roughness parameters of the coating as obtained from optical profilometer are presented in Table 3. The coating roughness depends on the target elemental composition and the deposition parameters [29]. The more the composition of low-melting-point elements, such as aluminum, in the target, the higher is the probability of rougher surface due to increase formation of macro and microparticles. Selection of optimum deposition parameters as well as careful surface preparation can help minimized surface defects for a given target composition [30]. The values of the arithmetic mean (RA), root-mean-square (RMS) and the peak-to-valley (PV) presented in Table 3, show that the coated sample possessed a low surface roughness as compared to the polished uncoated sample. The thickness of the coating as shown in Fig. 1(c) is between 4 and 5μm. The elemental compositions of the coating are also shown in Table 3, the composition corresponds to a non-stoichiometric Al0.68Cr0.32N coating with 1.77 Al/Cr ratio, similar to reported value in an earlier study [14].

Table 3.

Composition and surface properties of the AlCrN coating.

Coating chemical composition
Elements  Al  Cr 
Composition (at.%)|(wt.%)  33.3|(35.4)  18.6|(38.1)  48.1|(26.5) 
Surface roughness
Roughness parameters  RA (μm)  RMS (μm)  PV (μm) 
Uncoated W–25%Re–HfC  0.35±0.01  0.48±0.02  18.61±6.66 
AlCrN coated sample  0.15±0.02  0.23±0.01  8.85±0.13 

The XRD patterns of the substrate and coated sample are presented in Fig. 2. The peaks of the W–Re alloy and HfC reinforcement were detected in the composite substrate material and they matched with JCPDS PDF No. 03-065-8387 and 00-039-1491, respectively. The XRD pattern of AlCrN coated sample indicates the presence of the substrate material although some of the peaks have been depressed. The AlCrN coating crystalizes into the B1-NaCl cubic structure with typical diffraction peaks. The two main diffraction peaks of AlCrN coating are identified corresponding to the (111) and (200) planes. These peaks correspond to the peaks of CrN with JCPDS PDF No. 00-011-0065 and of AlN with JCPDS PDF No. 00-046-1200, which was also reported in earlier studies [31,32]. The estimated average grain size of the coating calculated based on the (111) and (200) peaks using the Scherrer equation is about 9.30nm. Considering that the instrumental broadening and possible microstrain were not decomposed from the observed peaks, the calculated average grain size is the minimum grain size of the coating obtainable.

Fig. 2.

X-ray diffraction pattern of W–25%Re–HfC and AlCrN coated sample.

3.2Micromechanical and scratch properties

Table 4 shows the micromechanical properties of the coated and uncoated samples from a minimum of six measurements. The coating possessed higher hardness value as compared to the hardness values of the substrate material. The Vickers hardness value of the coating was 2563HV against 717HV of the composite substrate. Wear resistance behavior of the coating is better understood by the relationship between the hardness and elastic modulus rather than just the hardness. Thus, the parameters such as the elastic strain to failure (H/E) and plasticity index (H3/E2) are of paramount importance in describing the wear resistance of coatings [33,34]. The coated sample exhibited relatively high H/E and H3/E2 values of 0.1 and 0.3, respectively, as compared to the uncoated sample with H/E and H3/E2 values of 0.034 and 0.0087, respectively; as shown in Table 4. These values for the coating were similar to that obtained from a previous study [14]. This substantiate the fact that the measured mechanical properties of the coating were not influenced by the substrate.

Table 4.

Mechanical properties of AlCrN coated sample.

Samples  Hardness (GPa)  Vickers (HV)  Elastic modulus (GPa)  H/E  H3/E2 (GPa) 
Uncoated W–25%Re–HfC  7.6±0.2  717±15  224.1±4.3  0.034  0.0087 
AlCrN coated sample  27.2±8.1  2563±763  264.2±54.6  0.101  0.306 

Scratch test was used to determine the cohesion and adhesion strength of the coating. Fig. 3 shows the AE signal and the COF as a function of the applied normal load. The AE signal is used to determine the cohesion force of the coating, denoted as Lc1, while the change in the COF is used to determine the adhesion force, denoted as Lc2, in Fig. 3. The coating showed high cohesion and adhesion strength of 9.5 and 25N, respectively. This is an indication of the coating good adhesion to the composite tool material. The AE signal is suitable for detecting the initiation of crack within the coating. This is by detecting the acoustic (elastic) waves radiated due to irreversible changes in the internal structure of the coating caused by cracks, which can be observed from the sudden increase in the signal, as shown in Fig. 3. The crack consequently increases as the applied load increases until a sudden increase in the COF occurs (illustrated in Fig. 3) indicating a major failure in the coating. The COF is observed to be constant after the coating failure, this indicates a change in the counterface-substrate interaction at the interface.

Fig. 3.

Scratch acoustic emission and COF as a function of the applied load.


Fig. 4 shows the optical images of the scratch track and the excerpts X and Y show the micrograph around the point of cohesive and adhesive failures respectively. The failure mechanism shows recovery spallation behind the indenter on both sides of the track with very small coating chipping. It is observed that no delamination of the coating occurs inside the scratch track even up to the maximum applied load, rather, tensile cracks failure was evident. Fig. 5(a) shows SEM images of the track and the compositional analysis within the track. It is interesting that the coating was still present after the adhesive failure. This suggests that the mode of failure of the coating inside the track is neither by spallation nor delamination of the coating, the failure was by tensile cracking and thinning. Thus, the coating combines between recovery spallation at the edges of the track and tensile cracking inside the track [35] as shown in Fig. 5(b) and (c), respectively. These failure modes are typical of PVD nitride coatings deposited on tough substrates and thus necessitate a high adhesion of the coating as illustrated by Bull [36]. W–25%Re–HfC is a tough and easily machinable composite material and this can be sensed from the large groove in the SEM image of the scratch track in Fig. 5.

Fig. 4.

Optical micrograph of the scratch track and excerpts X and Y showing micrograph around cohesive and adhesive failures, respectively.

Fig. 5.

SEM images of the scratch track and the composition analysis of the track after adhesion failure.

3.3Wear test analysis

Ball-on-disk wear test configuration was used for the wear resistance evaluation. The initial Hertzian contact pressure was about 1GPa. High contact pressure is an advantage of ball-on-disk test configuration, however, after run-in, the contact area is enlarged by wear thus reducing the contact pressure as the test progress further. Fig. 6 shows the COF over a total sliding distance of 500m. It is found that AlCrN coated sample exhibited lower and steadier COF with minimal fluctuation as compared to the uncoated sample. The COF of the coated sample is 30% lower compared to uncoated sample during the steady-state regime. The COF of uncoated and coated samples in the steady-state regime is about 0.68 and 0.47, respectively. The COF of the AlCrN coating in this study is also much lower than the COF reported for similar coating (0.65–0.75) by Mo et al. [37,38]. High COF is undesirable as they often lead to significant wear of the sliding interfaces. This is because high COF indicates a high tangential force, which implies that higher energy is required to overcome these forces at the interface between the sample disk and the ball counterface. Hence, for the continuous relative motion between the contacting parts, this high energy needs to be dissipated, and consequently, this results in large plastic deformation which often leads to high wear.

Fig. 6.

COF as a function of the sliding distance.


Figs. 7(a), (b) and 8 respectively show the 3D profilometer images, and the depth profile across the wear track. Fig. 7(c) and (d) show the surface profile from which the roughness inside the wear track was measured for the uncoated and coated samples, respectively. Roughness measurement shows that the wear track of the uncoated sample is rougher than that of the coated sample. The roughness parameters RA and RMS inside the wear track of the AlCrN coated sample are 0.19±0.03μm and 0.24±0.05μm, respectively. However, for the uncoated sample, RA and RMS inside the wear track are 0.45±0.07μm and 0.60±0.09μm respectively, which are about 60% higher than the coated sample. This is an indication of the difference in the wear mechanism for these samples. The high roughness inside the wear track is obviously due to the deep and wide grooves from the plowing action of the asperities or a three-body abrasive wear mechanism. Also, it can be observed that the wear volume differs significantly.

Fig. 7.

Optical 3D micrograph showing; wear track of (a) uncoated and (b) AlCrN-coated samples, and surface profile inside the wear track of (c) uncoated and (d) AlCrN-coated samples.

Fig. 8.

Wear track depth for the uncoated and coated samples.


The uncoated sample experienced considerably high wear compared to the coated sample. The maximum wear track depth of the coated sample was around 1.35μm, about 90% lower than the depth of the uncoated sample with a wear track depth of about 12μm. The depth profile (Fig. 8) indicates that the coating experience very low wear and the substrate was fully protected all through the sliding distance. The corresponding volume of the wear track was 4.8 and 177.8μm3 for coated and the uncoated samples, respectively. Interestingly, the volume of material loss by wear for the uncoated sample was found 97% higher than that of the coated sample. This observed increase in the wear resistance of the coated sample can also be attributed to the high elastic strain to failure and plasticity index as compared to the uncoated sample. The specific wear rate of the samples and the weight loss of the respective counterface steel ball are shown in Fig. 9. AlCrN coated sample possessed lower wear rate of 42.83mg/kNm against the uncoated sample with 443.33mg/kNm specific wear rate. The use of the AlCrN coating resulted in reducing the wear rate of the composite material by about 90%. It is also worth mentioning that the AlCrN coating exhibited higher wear resistance properties despite the high load and sliding speed used in this study as compared to similar coating reported in [37–39]. For example, Vettivel et al. [39] reported the wear rate of AlCrN to be 50, 79, and 120103mg/kNm at 10, 20, and 30N, respectively. These values are much higher than the value reported for the AlCrN coating in this study at 30N applied normal load.

Fig. 9.

Specific weight loss for the samples and the ball counterface.


Similarly, the steel ball counterface against the coated sample was not severely worn as compared to the steel ball counterface against the uncoated composite sample. Fig. 10 shows the worn ball surfaces and it is found the ball scar diameter is about 1 and 1.58mm when used against coated and uncoated samples respectively. The ball scar diameter against the uncoated sample is about 1.5 times larger than ball scar against the coated composite sample. This indicates that the interface between the ball and the uncoated composite sample experienced high plastic deformation required to overcome the high interfacial forces. A close observation of the steel ball surface after wear test as seen in Fig. 10(a) shows significant grooves along the sliding direction due to plastic deformation. This substantiate the fact that the wear mechanism leading to the high wear rate of the uncoated sample is preponderantly abrasive wear by plastic deformation. In this type of wear mechanism, the asperities undergo severe plastic deformation leading to the subsurface strain hardening. This is followed by plowing of the material through a series of grooves as shown in the SEM image of the wear track of the uncoated sample in Fig. 11(a) and (b). The SEM image (Fig. 11(b)) also indicates somewhat oxidation wear mechanism of the uncoated samples as expected due to the elevated temperature at the interface.

Fig. 10.

Optical images of the ball scar after wear test against (a) uncoated and (b) coated composite samples.

Fig. 11.

SEM micrograph of the wear track showing the oxide layer and wear debris of (a), (b) uncoated and (c), (d) AlCrN coated samples.


The wear mechanism of the AlCrN coated composite sample is depicted in Figs. 11(c), (d) and 12. The magnified SEM micrograph of the wear track is shown in Fig. 11(d) while Fig. 12 shows the EDS analysis of the different section observed in the micrograph. It can be found that there is a formation of an oxidized layer with a high composition of iron and chromium which are obviously from the counterface steel ball. This oxide layer as shown in the EDS spectrum is basically oxides of iron and chromium. Fractured debris, due to low cycle fatigue, and the formed oxidized layer can be seen in the wear track. EDS analysis of the exposed section revealed that coating is intact though oxidized. The oxidation of the coating is indicated by the reduction in the content of nitrogen in the coating from about 26wt.% to 17wt.% and by the presence of significant amount of oxygen of about 10wt.%. The content of Al and Cr in the wear track, 32wt.% and 35wt.%, respectively, is about the same amount as in the elemental composition of the coating. This confirmed that the coating is intact and that it protects the substrate from wear. Al2O3 and Cr2O3, which are typical oxides formed on AlCrN coating at elevated temperatures, are known for their excellent mechanical properties, good adhesion, and high density thus making them very protective. Mo and Zhu [40] have also reported the formation of these oxides on AlCrN coating and how they influence the tribological properties during sliding. Therefore, the wear mechanism of the coated sample is predominantly oxidation which is accompanied by low cycle fatigue thus led to the fracturing of the oxide layer to form the observed debris.

Fig. 12.

EDS analysis of the wear track of AlCrN coated sample showing the wear debris and wear mechanism.


The mechanical and tribological properties of PVD AlCrN coating on spark plasma sintered W–25%Re–HfC composite material is considered. PVD AlCrN coating was deposited on W–25%Re–HfC composite material intended for FSW tool application. This study demonstrates by way of tribological analysis the feasibility of improving the life and performance of expensive FSW tools by applying cathodic arc AlCrN PVD coating. The mechanical, scratch and wear resistance properties were characterized, and following are some important conclusions from the study:

  • 1.

    AlCrN coated W–25%Re–HfC sample exhibits improved mechanical properties with hardness and modulus of elasticity of about 27 and 264GPa, respectively.

    The coating exhibited excellent adhesion to the substrate and high cohesive strength.

  • 2.

    The COF of AlCrN coated sample is 30% lower than the COF of the uncoated sample. Similarly, the measured wear volume of the coated sample is significantly lower as compared to the uncoated sample. Quantitatively, the specific wear rate of the coated sample is 90% lower than the wear rate of the uncoated sample.

  • 3.

    The improved wear resistance of the coating is attributable to the combined excellent mechanical properties, high adhesion to the substrate, the relatively low COF and the formation of oxides which is well-bonded, highly dense and protective tribolayer.

Conflicts of interest

The authors declare no conflicts of interest.


The authors acknowledge the support received from the Deanship of Research, King Fahd University of Petroleum and Minerals for funding this work under the internal research grant (Project IN161018). Also, the Ph.D. Scholarship provided by King Fahd University of Petroleum and Minerals (KFUPM) for A.Y. Adesina is gratefully acknowledged.

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Journal of Materials Research and Technology

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