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Influence of temperature of sub-zero treatments on the wear behaviour of die steel
Wear 267 (2009) 1361–1370
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Influence of temperature of sub-zero treatments on the wear behaviour of die steel D. Das a , K.K. Ray b , A.K. Dutta c,∗ a b c
Department of Metallurgy and Materials Engineering, Bengal Engineering and Science University, Shibpur, Howrah 711103, India Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India Department of Mechanical Engineering, Bengal Engineering and Science University, Shibpur, Howrah 711103, India
a r t i c l e
i n f o
Article history: Received 1 September 2008 Received in revised form 25 November 2008 Accepted 25 November 2008 Keywords: Sub-zero treatment Cryogenic treatment Die steel Wear behaviour Mode of wear Mechanism of wear
a b s t r a c t This study examines the influence of temperature of sub-zero treatment on the wear behaviour of AISI D2 steel. A series of dry sliding wear studies have been made under constant normal load at varying sliding velocities. Emphasis has been laid to understand the operative modes and mechanisms of wear by the estimation of specific wear rates and detailed characterizations of the worn surfaces, wear debris and subsurfaces with the help of scanning electron microscope (SEM) examinations coupled with energy dispersive X-ray (EDX) microanalyses. The obtained results unambiguously infer that lower the temperature of sub-zero treatment higher is the improvement in wear resistance. Wear resistance can increase by 1.5–125 times depending on sliding velocity while hardness increases only by 4.2% at the lowest temperature of sub-zero treatment (77 K) compared to the conventionally treated specimens. These results corroborate well with the reduction in retained austenite content associated with simultaneous increase in the amount of secondary carbide particles with lowering of sub-zero treatment temperature. The operative modes and mechanisms of wear are identified as either mild oxidative or severe delaminative, which depends on the temperature of sub-zero treatment and the sliding velocity of the wear test. © 2009 Elsevier B.V. All rights reserved.
1. Introduction It is well established that sub-zero treatment improves several mechanical properties of tool steels like wear resistance [1–10], hardness [5–12], fatigue resistance [12], toughness [13], etc., apart from improving dimensional stability of engineering components [14]. The merit of sub-zero treatment is popularly attributed to the transformation of retained austenite (␥R ) to martensite. As a consequence, sub-zero treatment is being commercially used for the improvement of wear resistance of tool/die steels. In a series of recent investigations [7–10], the present authors have examined the nature of phase transformation at deep cryogenic temperature and have studied their role on the wear behaviour of AISI D2 steel. The outcomes of these investigations have indicated that apart from the transformation of ␥R to martensite, nucleation and growth of secondary carbides play an important role. In addition, these studies have also shown that the degree of improvement in wear resistance of tool steels is dependent on the selection of wear test parameters.
∗ Corresponding author. Tel.: +91 33 2668 4561/62/63; fax: +91 33 2668 4564/2916. E-mail address: [email protected] (A.K. Dutta). 0043-1648/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2008.11.029
The process of sub-zero treatment consists of controlled cooling of conventionally hardened specimens to a selected sub-zero temperature followed by controlled heating of the specimens back to the ambient temperature for subsequent tempering treatment. The sub-zero treatment can be classified into three different temperature regimes: cold treatment (CT, ≥193 K), shallow cryogenic treatment (SCT, 193–113 K) and deep cryogenic treatment (DCT, 113–77 K). The influence of these temperature regimes on the wear characteristics of sub-zero treated steels has not been systematically examined so far. The major aim of this report is to bring forth the influence of temperature of sub-zero treatment on the wear behaviour of steels under different wear test conditions. Until recently, the accepted temperature for sub-zero treatment has been 193 K where dry ice can be used for cooling. However, the results of few recent studies [1–4] suggest that the temperature of sub-zero treatment should be <193 K in order to obtain the maximum improvement in mechanical properties of tool steels; the lowest temperature may be 77 K, the boiling temperature of liquid nitrogen at normal atmospheric pressure. Mohan Lal et al. [2] have shown that lower is the temperature higher is the improvement in wear resistance for D3 and M2 steels. In contrast, Moore and Collins [15] have suggested that there is an optimal temperature for sub-zero treatment of tool steel. Schiradelly and Diekman [16] have applied Taguchi Design of Experiment (DOE) to identify and to optimize the critical parameters of cryogenic treatment for a
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Table 1 Nominal composition of the investigated AISI D2 steel. Element
C
Mn
Si
S
P
Cr
Mo
V
Fe
Weight (%)
1.49
0.29
0.42
0.028
0.029
11.38
0.80
0.68
Balance
martensitic stainless steel. Using the analysis of variance (ANOVA) associated with the results of wear tests, these authors have shown that the most significant parameter is the soaking temperature (72% in contribution). The reported improvements in wear resistance of die/tool steels due to cryogenic treatment varies widely, for example from a few percent to a few hundred percent for the same material [9]. Such inconsistent reports are hindering the exploitation of the inherent potential of cryogenic treatment for its commercial exploitations. The lack of systematic studies of wear behaviour of cryogenically treated tool steels with proper identification of the operative modes and mechanisms of wear is mainly responsible for the reported wide scatter in the degree of improvement of wear resistance of tool steels by cryogenic treatment [8–10]. Therefore, it is necessary to compare the wear resistance of tool/die steels subjected to different sub-zero temperature treatments with proper identification of the operative modes and mechanisms of wear which is the focus of this study. 2. Experimental procedure 2.1. Material and treatments A commercial forged AISI D2 steel bar has been chosen for this investigation. The nominal chemical composition of the steel is presented in Table 1. Specimens of this steel have been subjected to conventional treatment (CONT), cold treatment (CT), shallow cryogenic treatment (SCT) and deep cryogenic treatment (DCT) in separate batches. The CONT consists of hardening and tempering; while in CT, SCT and DCT, an additional step of controlled sub-zero treatment with lowest quenching temperature (TLQ ) as 198, 148 and 77 K, respectively, has been incorporated in-between hardening and tempering treatments. In sub-zero treatment, conventionally hardened specimens have been cooled from ambient temperature to the selected TLQ , held there for ∼5 min for homogenization of temperature before heating back to the ambient temperature for subsequent tempering treatment. Uniform cooling and heating rates of ∼0.75 K min−1 were maintained in each of the sub-zero treatments. The hardening and single tempering treatments following ASM standard [17] have been carried out at 1297 K for 30 min and at 483 K for 120 min, respectively. In order to distinguish the specimens subjected to different TLQ , these have been ascribed with codes as shown in Table 2. These designations are followed in the subsequent discussions. One can note here that TLQ for CONT is 303 K, i.e., with out any sub-zero treatment. 2.2. Microstructural characterizations Microstructural examinations have been carried out on polished and picral-etched specimens using both optical and scanning elecTable 2 Different heat treatments and the sample codes. Sample code
CONT CT SCT DCT
1293 K, 30 min
Lowest quenching temperature, TLQ (K) 303 198 148 77
2.3. Hardness measurement The bulk hardness of the differently treated specimens has been measured by Vickers hardness tester using 60 kgf load. At least fifteen readings are considered for estimating the average value of hardness, and the estimated standard errors associated with the measured Vickers hardness numbers (VHN) were ±5. 2.4. Sliding wear tests Dry sliding wear tests have been performed following ASTM standard G99-05 by using a computerized pin-on-disc wear testing machine. Cylindrical specimens of 4 mm diameter and 30 mm length have been used as static pins, whereas, tungsten carbide coated En 35 steel disc (surface hardness HV ≈ 1750) of diameter 160 mm is selected as the rotating counter surface. The wear tests have been carried out using a fixed normal load (FN ) of 98.1 N (10 kgf) at different sliding velocities (SV ) of 1.00, 1.25 and 1.50 m s−1 and under ambient conditions of 298 K and 60% relative humidity. The rpm of the disc was adjusted for a selected track diameter (100–140 mm) for achieving a predetermined SV . The wear tests were continued either for sliding distance of ∼2000 m or for the duration that resulted into cumulative height loss of the pin specimen of ∼2 mm (equivalent to ∼25 mm3 of cumulative wear volume loss), whichever occurred earlier. Average wear rates have been estimated from the recorded cumulative height loss of the specimens with respect to sliding distance in the steady-state wear regime considering at least three test results under identical conditions. 2.5. Characterizations of the worn surfaces, wear debris and subsurfaces The worn surfaces at the end of the wear tests and generated debris collected in the steady-state wear regime have been examined under SEM along with the energy dispersive X-ray (EDX) microanalyses for all specimens at all wear test conditions. Some selected worn samples were sectioned perpendicular to the worn surface and parallel to the sliding direction using wire electrodischarge machining. The sectioned faces of the specimens were mounted, polished to 0.25 m finish, etched using picral solution and were examined using SEM in back-scatter mode to record the subsurface features of the worn specimens. 3. Results and discussion 3.1. Microstructures and hardness
Description of heat treatment cycles Hardening
tron microscopes (SEMs). The carbide particles have been classified as primary carbides (PCs) and secondary carbides (SCs) considering equivalent spherical diameter for PCs as >5 m, while that for SCs ≤ 5 m. The volume percents of PCs and SCs have been estimated by image analyses of digitally acquired suitable optical and SEM micrographs using Leica QMetals software in Leica QWin V3 environment. Identification of the phases, and the measurement of retained austenite (␥R ) following ASTM standard E975-00 have been done by X-ray diffraction (XRD) analyses with Mo K␣ radiation at 0.01◦ s−1 scan rate on bulk specimens with the help of PHILIPS X’Pert software.
Tempering
483 K, 120 min
Fig. 1 depicts representative SEM micrographs of conventionally treated and a typical sub-zero treated specimens. These exhibit large elongated dendritic type of primary carbides (PCs) and nearly spherical secondary carbides (SCs) on tempered martensite matrix. The occasional patches of retained austenite (␥R ) have been detected only in SEM micrographs of CONT [7] and CT specimens. The volume fraction of ␥R has been estimated by XRD technique,
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Fig. 1. Typical representative SEM micrographs of (a) CONT and (b) SCT specimens exhibiting primary carbides (PCs) and secondary carbides (SCs) of different sizes. The smaller SCs are shown inside circles.
detailed procedure of which has been reported earlier [9]. The amounts of ␥R in the different types of specimens are shown in Fig. 2. The amount of ␥R in CT specimen is 4.6 ± 0.5 vol.% in comparison to 9.8 ± 0.7 vol.% in CONT specimen. The amount of ␥R in both SCT and DCT specimens is considered to be below the detection limit (<2 vol.%) of XRD technique. This observation is in agreement with the reported results for AISI D2 steel in DCT [6–10] and SCT [14] conditions. The obtained results suggest that CT reduces the ␥R substantially as compared to CONT; while both SCT and DCT in between conventional hardening and tempering treatments almost completely remove it. Since the martensite finish (Mf ) temperature for AISI D2 steel is ≈148 K, SCT and DCT with TLQ ≤ 148 K almost completely transforms ␥R to martensite, and hence the observed results are expected. The amount of PCs is found to be invariant (≈7.0 vol.%) while the amount of SCs varied in CONT, CT, SCT and DCT specimens. The PCs in the investigated specimens have been identified mainly as M7 C3 (M = Fe, Cr, Mo, V) with small amount of Cr7 C3 by XRD analyses. The variations in the amounts of SCs with TLQ as determined by image analyses are illustrated in Fig. 2. The SCs in all the specimens have been identified as M23 C6 (M = Fe, Cr, Mo, V) type and this is in agreement with earlier reports [7–10,13]. It is interesting to note that the application of sub-zero treatments after conventional hardening does not alter the nature of PCs and SCs. The amount of SCs increases with decreasing temperature of sub-zero treatment. As the temperature of sub-zero treatment is lowered, the amount of ␥R
decreases resulting in higher amount of martensite; the increased amount of martensite naturally leads to higher amount of carbide precipitation on subsequent tempering treatment as observed in Fig. 2. The variation of hardness with TLQ is shown in Fig. 3. The results indicate that sub-zero treatment increases hardness of D2 steel though by a smaller margin of up to 4.2%. The increase in hardness is attributed to reduction in ␥R content with associated increase in the amount of SCs (Fig. 2). It is observed from the results in Fig. 3 that the rate of increase in hardness with TLQ diminishes with lowering of TLQ . In addition, the nature of distribution of SCs (Fig. 1) can also influence the increment in hardness with TLQ due to alteration in the inter-particle spacing leading to the variation in micro-residual stresses and sub-structural changes like dislocation density.
Fig. 2. Variation of amount of retained austenite (␥R ) and secondary carbides (SCs) as a function of lowest quenching temperature.
Fig. 3. Influence of lowest quenching temperature on the bulk hardness of differently treatment specimens.
3.2. Wear rates Typical plots of cumulative wear volume loss versus sliding distance for all the specimens tested at SV of 1.25 m s−1 under FN of 98.1 N are shown in Fig. 4. The results in Fig. 4 exhibit distinguished regimes of ‘running-in’ and ‘steady-state’ wear [18–20] for each specimen and the difference in the nature of variation in the wear behaviour of these specimens. The results exhibit that the volume losses for CONT and CT specimens are significantly higher than those of the SCT and DCT specimens. Furthermore, volume
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Fig. 4. Typical plot of cumulative wear volume loss (CWVL) versus sliding distance (SD ) for differently treated specimens tested at sliding velocity of 1.25 m s−1 . Inset shows enlarged view of same for CONT and CT specimens. The arrow indicates the transition of initial running-in wear to the final steady-state wear.
loss of CONT specimen is appreciably higher than that of CT specimens, whereas the volume loss for SCT specimens is marginally higher than that of DCT specimen. The results obtained for tests conducted at SV of 1.00 and 1.50 m s−1 is similar in nature. The extent of running-in wear depends on the experimental conditions like (i) surface roughness of both the pin and the disc, (ii) perpendicularity of the pin with respect to the disc and (iii) the rise in temperature at the contact area, etc. [20]. The results in Fig. 4 indicate that (i) the wear rates of all specimens are considerably higher in the runningin regime than that in the steady-state regime and (ii) the wear rates in the running-in regime decreases with lowering of TLQ . Steady-state wear, which is known to depend on the material’s characteristics [18–20], has been analyzed to reveal the effect of TLQ on the wear behaviour of the selected steel. For this purpose, the specific wear rate (WS ) has been calculated by the equation: WS = VL /(FN × SD ); where VL is the wear volume loss of the pin specimens in the steady-state regime in mm3 , FN is the applied normal load in N and SD is the sliding distance corresponding to the steady-state wear regime in mm. The magnitudes of WS estimated at different SV for specimens with different TLQ are compiled in Fig. 5. The results in Fig. 5 assist to infer that: (i) at SV = 1.00 m s−1 , WS sharply decreases for TLQ = 303–198 K and then its magnitude decreases marginally with further lowering of TLQ , (ii) at SV = 1.25 m s−1 , WS marginally decreases between TLQ = 303 and 198 K and between TLQ = 148 and 77 K, whereas WS drastically decreases between TLQ = 198 and 148 K, and (iii) at SV = 1.50 m s−1 , the magnitude of WS marginally decreases for the entire range of investigated TLQ . The nature of variation of WS for different specimens at the selected SV or for the same specimen at different SV can be explained only with the knowledge of the operative mechanisms of wear and this is the content of the next section. 3.3. Mechanisms of wear In order to identify the operative wear mechanisms, the topography of the worn surfaces as well as the nature and morphology of the generated wear debris have been characterized for all the tested specimens. Some typical representative examples in Figs. 6–10
Fig. 5. Variation of estimated specific wear rate (WS ) with lowest quenching temperature (TLQ ) of the specimens tested at different sliding velocities (SV ). Shading has been used to demarcate WS into two different regimes based on the operative modes and mechanisms of wear (see text for details).
illustrate the operative wear mechanisms of differently treated specimens tested at varying SV . The micrographs in Figs. 6 and 7 bring to light the difference in the operative wear mechanisms for CONT specimens compared to that for different sub-zero treated specimens at SV = 1.00 m s−1 . The worn surface of CONT specimen is rough but metallic in nature, and also exhibits fracture ridges and deformation lips stretched parallel to the sliding direction (Fig. 6a). The presence of deformation lip infers that the CONT specimen has undergone heavy plastic deformation during wear test. In contrast, the worn surfaces of the CT (Fig. 6b), SCT (Fig. 6c) and DCT (Fig. 6d) specimens are smoother and exhibit presence of patches of oxide. The micrograph of oxide depleted region of worn surfaces of these specimens shows cracking and pull-out of PCs apart from the presence of surface grooves as illustrated in Fig. 6e. Surface grooves on the worn surface are formed due to the trapping of loose hard PC particles in-between the pin and disc surfaces. The micrograph of the wear debris of CONT specimen (Fig. 7a) exhibits shiny metallic appearance and a plate-like morphology. In contrast, the debris of the CT (Fig. 7b), SCT (Fig. 7c) and DCT (Fig. 7d) specimens appear as black coloured powder in naked eye; and these are much finer in size and oxide in nature (Fig. 7e). These observations reveal that the wear mechanism at SV = 1.00 m s−1 and FN = 98.1 N for CONT specimen are plastic deformation induced delaminative wear [21], whereas that for the sub-zero treated specimens is predominantly oxidative [22] coupled with cracking and pull-out of PCs. Representative features on the worn surfaces of CONT, CT, SCT and DCT specimens tested at SV = 1.25 m s−1 are shown in Fig. 8; the corresponding morphology of the wear debris and their typical representative EDX profiles are presented in Fig. 9. The observed features on the worn surfaces for both CONT and CT specimens tested at SV = 1.25 m s−1 (Fig. 8a and b) closely resemble with that observed for CONT specimen tested at SV = 1.00 m s−1 (Fig. 6a). Worn surface features for SCT and DCT specimens tested at SV = 1.25 m s−1 (Fig. 8c and d) are similar to that of CT, SCT and DCT specimens tested at SV = 1.00 m s−1 (Fig. 6b–d). However, the worn surfaces of SCT and DCT specimens show regions of thick compacted oxide layer at SV = 1.25 m s−1 (Fig. 8c and d) unlike that in the worn surface of SCT and DCT specimens tested at SV = 1.00 m s−1 (Fig. 6c and d). Furthermore, examinations of worn surfaces of SCT and DCT specimens tested at SV = 1.25 m s−1 does not reveal any obvious signature of cracking and/or pull-out of PCs even at higher magnifications.
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Fig. 6. Representative secondary electron SEM micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.00 m s−1 . (e) A typical high magnification backscatter electron SEM micrograph of oxide depleted region as marked by area-1 in the worn surfaces of CT, SCT and DCT specimens.
The wear debris of CONT and CT specimens show large metallic platelets unlike fine oxide debris shown by SCT and DCT specimens (Fig. 9). Typical representative EDX profile of the wear debris of CONT and CT specimens is shown in Fig. 9e and that for SCT and DCT specimens is presented in Fig. 9f. The recorded EDX profiles exhibit oxygen peak of higher intensity for SCT and DCT specimens compared to that from CONT and CT specimens (Fig. 9f vis-à-vis Fig. 9e). This observation suggests that the generated wear debris for CONT and CT specimens is predominantly metallic in nature and that for SCT and DCT specimens is predominantly oxide in nature. The EDX profiles of both types of wear debris indicate the presence of tungsten that has been transferred from the counter body to the specimen surface by adhesion. Evidences in Figs. 8 and 9 assist to reveal that the operative mechanisms of wear for both CONT and CT specimens is delaminative at SV = 1.25 m s−1 , whereas that for SCT and DCT specimens is oxidative wear under same experimental condition. The surface topography of all specimens tested at SV = 1.50 m s−1 indicates delaminative wear. Representative evidences are shown
in Fig. 10 for SCT and DCT specimens. The wear debris generated during wear tests of these specimens is metallic platelets (inserts in Fig. 10) corroborating with the observed features of the worn surfaces. In summary, the observed wear mechanism for CONT specimens for all tested SV is delaminative wear; whereas the operative wear mechanisms change from predominantly oxidative to delaminative at SV = 1.25 m s−1 for CT specimens and at SV = 1.50 m s−1 for SCT and DCT specimens. 3.4. Transition of modes and mechanisms of wear The modes of wear may be categorized using wear mechanism map as suggested by Lim and Ashby [23] or by magnitude of WS as suggested by Wang et al. [24] or through magnitude of wear coefficient as suggested by Rabinowicz [20]. Identification of the operative wear mechanisms at different test conditions is known to assist in revealing the operative mode of wear [23]. In the present investigation, the operative mechanisms of wear have been identified by the detailed characterizations (for all specimens tested
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Fig. 7. Representative secondary electron SEM micrographs of generated wear debris corresponds to the steady-state wear regime of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.00 m s−1 . All these micrographs are taken at the same magnification of 250X, whereas (e) is a typical high magnification view of wear debris of CT, SCT and DCT specimens.
at different SV ) of worn surfaces and generated debris. These are shown in Fig. 5, which reveals that the data related to the WS clusters into two different regimes depending on the mechanisms of wear, and the magnitude of WS between these regimes varies about an order of magnitude. Considering the magnitude of WS and the wear mechanisms, these regimes are termed as ‘mild-oxidative’ and ‘severe-delaminative’, in which ‘mild’ and ‘severe’ refer the mode of wear. The results in Fig. 5 indicate that the transition SV for mild-tosevere wear for CT is 1.00 m s−1 and that for SCT and DCT specimens are 1.25 m s−1 . Within the tested range of SV (1.00–1.50 m s−1 ), the WS for CONT specimens lies in the severe wear regime (Fig. 5) and corresponding wear mechanism has been identified as delaminative wear (Figs. 6–9). These results suggest that transition SV for mild-to-severe wear for CONT specimens occurs at SV lower than 1.00 m s−1 . Additional experiments were carried out to find out the transition SV for CONT specimens and this has been found to be 0.75 m s−1 corresponding to predominantly oxidative wear. The WS of the CONT specimen at the transition SV is 8.10 × 10−8 mm3 N−1 mm−1 . The transition SV for mild-to-severe wear is a function of microstructure and mechanical property of tested pin specimens, when the other wear test parameters remain constant [23,24]. The observed variation in the transition SV for CONT, CT and SCT specimens is expected, since these specimens exhibit significant differences in their microstructures (Fig. 2) and hardness values
(Fig. 3). The obtained transition SV for both SCT and DCT specimens are same (1.25 m s−1 ), because the microstructures of SCT and DCT specimens are nearly similar, both being free from ␥R and having almost identical amounts of SCs (Fig. 2). In addition, the difference in hardness values of these specimens is <0.5%. Hence, it is natural that both SCT and DCT specimens exhibit same transition SV . It is also interesting to note that the WS of both SCT and DCT specimens decreases about an order of magnitude with increase of SV from 1.00 to 1.25 m s−1 even when the operative mode and mechanism of wear (mild-oxidative) remain unchanged. Similar observation has been reported earlier for dry sliding wear of steel specimens; for example, Sullivan and Hodgson [25] have shown that the wear rate of 52100 steel decreases by an order of magnitude with increase in SV , when SV lies below transition SV for mildto-severe wear. This behaviour is related to the rate of formation and retention of oxide at the contact surface of the pin specimen. Increase in SV increases the flash temperature that, on one hand, accelerates the formation of tribochemical reactions at the contact junctions [22,23]. On the other hand, it increases the degree of work-hardening of the material beneath the contact junctions leading to the development of increased depth of work-hardened layer with enhanced hardness [25]; this, in turn, assists in retaining oxides at the contact junctions resulting into development of thicker oxide layer [22,25]. Therefore, increased SV increases the thickness of oxide layer at the contact surfaces which is evident
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Fig. 8. Representative secondary electron SEM micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.25 m s−1 .
from the SEM micrographs of wear debris of both SCT and DCT specimens (Fig. 7 vis-à-vis Fig. 9). Development of thicker oxide layer is thus responsible for the reduction in the WS values with increasing SV from 1.00 to 1.25 m s−1 for both SCT and DCT specimens [22,25]. The results in Fig. 5 assist to infer that the maximum value of WS for mild-oxidative wear and the minimum value of WS for severe-delaminative wear are 9 × 10−8 and 4 × 10−7 mm3 N−1 mm−1 , respectively. Wang et al. [24] have earlier suggested for steel specimens that the upper limit of WS for mild-oxidative wear is 1 × 10−8 mm3 N−1 mm−1 and the lower limit of WS for severedelaminative wear is 2 × 10−8 mm3 N−1 mm−1 . These authors have also indicated that one can use the term ‘transition wear’, if the WS is in the region of 1 × 10−8 to 2 × 10−8 mm3 N−1 mm−1 . The present results, however, indicate that the demarcating lines for transition wear could be wider than that suggested by Wang et al. [24]. Kato [26] has earlier suggested that the WS of steel specimens
can vary over a wide range, from 10−5 to 10−13 mm3 N−1 mm−1 , and the exact magnitude of WS are influenced by several statistical variables involved in the wear tests. The major variables that determine the magnitude of WS are macro-material structures, surface roughness, flash temperature, local contamination, adhesive transfers, free wear particles and tribochemical reactions on the micro-scale on the contact surfaces [26]. Thus, the differences in the magnitudes of WS as obtained in the present investigation and that reported by Wang et al. [24] for demarcating the mild and the severe wear regimes may be attributed to the differences in the microstructures and properties of the selected materials, employed wear test methodologies and the chosen test conditions. Wang et al. [24] have studied the wear behaviour of 52100 and 1080 steels employing pin-on-ring testing method, whereas the material and test methodology of the present study are AISI D2 steel and pin-ondisc.
Table 3 Ratio of specific wear rates (WS ) and the operative modes and mechanisms of wear. Sliding velocity (m s−1 )
Specimen codes
Wear rate ratio, a
Mode and mechanism of wear Mode
Mechanism
1.00
CONT CT SCT DCT
1 6.5 7.3 7.9
Severe Mild Mild Mild
Delaminative Oxidative Oxidative Oxidative
1.25
CONT CT SCT DCT
1 1.8 103 125
Severe Severe Mild Mild
Delaminative Delaminative Oxidative Oxidative
1.50
CONT CT SCT DCT
1 1.2 1.3 1.5
Severe Severe Severe Severe
Delaminative Delaminative Delaminative Delaminative
a
is ratio of WS of CONT specimen to that of CT, SCT and DCT specimens.
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Fig. 9. Representative SEM micrographs of generated wear debris corresponds to the steady-state wear regime of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.25 m s−1 . All micrographs taken at the same magnification of 250×, whereas the insets in (c) and (d) at 1500× are detailing the features of the same micrographs. (e) and (f) are representative EDX profiles taken from the rubbed surfaces of the wear debris as marked by area-1 in (a) and (b), and by area-2 in the inset of (c) and (d), respectively.
Fig. 10. Representative SEM micrographs of worn surfaces at the end of wear tests of (a) SCT and (b) DCT specimens tested at SV = 1.50 m s−1 . Insets represent the generated wear debris corresponds to the steady-state wear regime for the same samples.
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Fig. 11. Representative back scattered SEM subsurface micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (b) DCT specimens tested at SV = 1.25 m s−1 . WL: while layer; PDL: plastically deformed layer.
3.5. Influence of temperature of sub-zero treatments on wear behaviour In order to assess the influence of temperature of sub-zero treatments on the degree of improvement in wear resistance, a new parameter has been incorporated here which is considered as the ratio of WS of CONT specimen to that of the different sub-zero treated specimens at a given SV . The estimated values of for different SV are compiled in Table 3 along with the operative modes and mechanisms of wear. The results in Table 3 reveal that the improvement in wear resistance by sub-zero treatment is marginal when the operative mode and mechanism between CONT and sub-zero treated specimens are similar (severe-delaminative). Under this condition, the value of is 1.8 for CT specimen at SV = 1.25 m s−1 and it varies from 1.2 to 1.5 at SV = 1.50 m s−1 for all the sub-zero treated specimens. In contrast, the magnitude of varies over two orders of magnitude when the operative mode and mechanism of wear for CONT specimens is severe-delaminative and that for the sub-zero treated specimens is mild-oxidative. Such wide variation in the magnitudes of is due to the similar or dissimilar operative modes and mechanisms of wear amongst the specimens considered (Table 3). These findings are in excellent agreement with the recent reports by the present authors [8,10] on the deep cryogenically treated and conventionally treated tool steels. However, comparison of values for SCT and DCT specimens at SV = 1.00 and 1.25 m s−1 indicates that the increases significantly at higher SV for these specimens, though the operative mode and mechanism of wear for the CONT (severe-delaminative) and that for the SCT and DCT (mild-oxidative) remains unchanged for above-mentioned SV values (Table 3). Such variation in occurs due to increase in WS for the CONT specimens and reduction of the same for the SCT and DCT specimens when SV increases from 1.00 to 1.25 m s−1 (Fig. 5). This apparent contradiction, however, is in agreement with earlier suggestions by Welsh [18,19]. He elucidated that the WS of steels increases with increasing SV, if the SV lies between the characteristics critical SV related to the mildto-severe and the severe-to-mild wear transitions. Identification of operative wear mechanism as delaminative for CONT specimen for
all the tested SV (1.00–1.50 m s−1 ) confirms that these SV indeed lie in between the two transitions SV for this specimen. Thus, according to this explanation of Welsh [18,19] it is natural that WS would increase with the increase in SV from 1.00 to 1.25 m s−1 as observed for the CONT specimens (Fig. 5). Welsh [18,19] and Sullivan and Hodgson [25] have earlier also shown that the WS of steels tend to decrease with increasing SV , if the applied SV is lower than the transition SV at which mildto-severe transition takes place. For both SCT and DCT specimens, mode and mechanism of wear are ‘mild-oxidative’ at SV = 1.00 and 1.25 m s−1 and ‘severe-delaminative’ at SV = 1.50 ms−1 . Thus, SV = 1.00 and 1.25 ms−1 are below the transition SV for these specimens. Therefore, it is natural that WS of these specimens would decrease with the increase of SV from 1.00 to 1.25 m s−1 (Fig. 5). This behaviour is related to the rate of tribochemical reaction at the contact surfaces as the increase in SV increases the flash temperature, which accelerates the formation of oxide layer at the contact junctions. At higher SV , development of more compacted oxide layer at the contact is expected to reduce the magnitude of WS as observed in the present study for the SCT and DCT specimens when SV increased from 1.00 to 1.25 m s−1 . This is well supported from the observations that (i) the area of worn surface covered by the compacted oxide layer increases with increasing SV (Fig. 6c and d vis-à-vis Fig. 8c and d) and (ii) the generated oxide debris are much thicker at higher SV (Fig. 7e vis-à-vis insets in Fig. 9c and d). An attempt to correlate the results of wear tests in Table 3 with the results of microstructural analyses in Figs. 1 and 2 reveal that TLQ has considerable influence on generating the nature of the microstructures and hence on the resulted wear properties of the selected steel. In order to get insights on the role of microstructural constituents on the variation of wear behaviour, the subsurfaces of all types of specimens tested at SV = 1.25 m s−1 were examined under SEM. Typical micrographs of the subsurfaces are presented in Fig. 11. The features beneath the worn surface of CONT and CT specimens are markedly different than that for SCT and DCT specimens (Fig. 11). The SCT and DCT specimens show only plastically deformed layer; in contrast, the CONT and CT specimens exhibit a well-defined ‘heavily deformed layer’, termed popularly
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in the literature as white (non-etching) layer [27], and followed by a plastically deformed layer. The presence of the white layer is the signature of deformation induced delaminative wear [21] which is indeed the governing wear mechanism for CONT and CT specimens at SV = 1.25 m s−1 . The presence of only plastically deformed layer and the absence of any white layer in the subsurfaces of SCT and DCT specimens are indicative of mild-oxidative wear for these specimens. The white layer of CONT specimen (Fig. 11a) shows presence of cracks that originated at a given depth beneath the worn surface and propagated parallel to the sliding direction; whereas, the cracking of white layer is less prominent for CT specimen (Fig. 11b). This observation is in agreement with the lower estimated WS of the CT specimen than that of CONT specimen at identical test conditions (Fig. 5). Severe-delaminative wear of steels occurs by continued plastic deformation at the contact surfaces of a test specimen leading to nucleation of cracks followed by its propagation as suggested by Suh [21]. The presence of soft ␥R associated with lower amount of hard SCs in the CONT and CT specimens as compared to those in the SCT and DCT specimens (Fig. 2) suggests that the former specimens are more prone to plastic deformation during wear at a lower SV for a constant FN or at lower FN at a constant SV . Therefore, specimens having less resistance to plastic deformation are expected to undergo mild-to-severe wear transition at a lower FN and/or SV . Thus, the present observation of higher SV for mild-tosevere wear transition for SCT and DCT specimens as compared to that for CONT and CT specimens is natural. Moreover, the fact that at SV = 1.00 m s−1 , the operative wear mechanisms for CONT and CT specimens are delaminative and oxidative wear, respectively, also supports the above argument; since CT specimen possesses higher amount of SCs and lower amount of ␥R than those in CONT specimens (Fig. 2). In summary, this investigation highlights the influence of temperature of sub-zero treatments on the microstructure and wear behaviour of die steel. 4. Conclusions The experimental results related to the influence of temperature of sub-zero treatments on the wear behaviour of selected AISI D2 steel and their pertinent analyses assist to infer the following major conclusions: (i) All types of sub-zero treatments appreciably improve the wear resistance of the die steels compared to the conventional treated (CONT) ones. However, the improvement in wear resistance by shallow (SCT) and deep cryogenic treatment (DCT) is significantly higher than that achieved by CT, and the maximum improvement is obtained by DCT. This is attributed to the decrease in the retained austenite content associated with the increasing amount of secondary carbide particles with lowering of temperature of the sub-zero treatments. (ii) The obtained hardness of AISI D2 steel for CONT and DCT are 759 and 791 VHN, respectively and typical values of their specific wear rates (WS ) are 1.03 × 10−6 and 8.26 × 10−9 mm3 N−1 mm−1 , respectively, at sliding velocity (SV ) of 1.25 m s−1 . (iii) Estimation of WS and detailed characterization of worn surfaces, generated debris and subsurfaces of all types of specimens assist to reveal that the severe mode of wear is associated with the deformation induced delaminative wear mechanism, whereas the mild mode of wear is associated with the predominantly oxidative wear mechanism coupled with occasional cracking and pull-out of primary carbide particles. (iv) The modes and mechanisms of wear are synergistically affected by the temperature of sub-zero treatments and the
SV of wear tests at a constant normal load. The mild-to-severe wear transition occurs at SV = 0.75 m s−1 for CONT specimen and at SV = 1.00 m s−1 for CT specimen, and at SV = 1.25 m s−1 for both SCT and DCT specimens. (v) The obtained results unambiguously lead to infer that lower the temperature of sub-zero treatment higher is the improvement in wear resistance. However, the degree of improvement is dependent on the SV of wear tests and/or the temperature of sub-zero treatment, which govern the operative modes and mechanisms of wear. Acknowledgement The financial assistance received from the University Grants Commission, Government of India [Grant No. F. No. 31-48/ 2005(SR)] to carry out a part of this research is gratefully acknowledged. References [1] E.A. Carlson, Cold treating and cryogenic treatment of steel, in: ASM Handbook, vol. 4, Heat Treating, 10th ed., ASM International, Metals Park, OH, 1990, pp. 203–206. [2] D. Mohan Lal, S. Renganarayanan, A. Kalanidhi, Cryogenic treatment to argument wear resistance of tool and die steels, Cryogenics 41 (2001) 149–155. [3] R.F. Barron, Cryogenic treatment of metals to improve wear resistance, Cryogenics 22 (1982) 409–413. [4] F. Meng, K. Tagashira, R. Azuma, H. Sohma, Role of eta-carbide precipitations in the wear resistance improvements of Fe–12Cr–Mo–V–1.4C tool steel by cryogenic treatment, ISIJ Int. 34 (1994) 205–210. [5] A. Molinari, M. Pellizzari, S. Gialanella, G. Straffelini, K.H. Stiasny, Effect of deep cryogenic treatment on the mechanical properties of tool steels, J. Mater. Proc. Tech. 118 (2001) 350–355. [6] D.N. Collins, J. Dormer, Deep cryogenic treatment of a D2 cold-worked tool steel, Heat Treat. Met. 3 (1997) 71–74. [7] D. Das, A.K. Dutta, V. Toppo, K.K. Ray, The Effect of cryogenic treatment on the carbide precipitation and tribological behavior of D2 steel, Mater. Manuf. Process. 22 (2007) 474–480. [8] D. Das, A.K. Dutta, K.K. Ray, On the enhancement of wear resistance of tool steels by cryogenic treatment, Philos. Mag. Lett. 88 (2008) 801–811. [9] D. Das, A.K. Dutta, K.K. Ray, Influence of varied cryotreatment on the wear behavior of AISI D2 steel, Wear 266 (2009) 297–309. [10] D. Das, A.K. Dutta, K.K. Ray, Inconsistent wear behavior of cryotreated tool steels: role of Mode and mechanism, Mater. Sci. Technol. (2008), doi:10.1179/174328408X374685. [11] D. Yun, L. Xiaoping, X. Hongshen, Deep cryogenic treatment of high-speed steel and its mechanism, Heat Treat. Met. 3 (1998) 55–59. [12] S. Zhirafar, A. Rezaeian, M. Pugh, Effect of cryogenic treatment on the mechanical properties of 4340 steel, J. Mater. Proc. Tech. 186 (2007) 298–303. [13] Y.M. Rhyim, S.H. Han, Y.S. Na, J.H. Lee, Effect of deep cryogenic treatment on carbide precipitation and mechanical properties of tool steel, Solid State Phenom. 118 (2006) 9–14. [14] C.H. Surberg, P. Stratton, Klaus Lingenhöle, The effect of some heat treatment parameters on the dimensional stability of AISI D2, Cryogenics 48 (2008) 42–47. [15] K.E. Moore, D.N. Collins, Cryogenic treatment of three heat treated tool steels, Key Eng. Mater. 86–87 (1993) 47–54. [16] R. Schiradelly, F.J. Diekman, Cryogenics–the Racer’s edge, Heat Treat. Progr. Nov. (2001) 43–50. [17] P.M. Unterweiser, H.E. Boyer, J.J. Kubbs (Eds.), Hear Treater’s Guide— Standard Practices and Procedures for Steel, 4th ed., ASM, Metal Park, OH, 1987, pp. 300–312. [18] N.C. Welsh, The dry wear of steels. I. The general pattern of behaviour, Philos. Trans. R. Soc., Ser. A 257 (1965) 31–50. [19] N.C. Welsh, The dry wear of steels. II. Interpretation and special features, Philos. Trans. R. Soc., Ser. A 257 (1965) 51–70. [20] E. Rabinowicz, Friction and Wear of Materials, 2nd ed., John Wiley and Sons, New York, United States, 1995, pp. 124–136. [21] N.P. Suh, The delamination theory of wear, Wear 25 (1973) 111–124. [22] T.F.J. Quinn, D.M. Rowson, J.L. Sullivan, Application of the oxidational theory of mild wear to the sliding wear of low alloy steel, Wear 65 (1980) 1–20. [23] S.C. Lim, M.F. Ashby, Wear mechanism maps, Acta Metall. 35 (1987) 1–24. [24] Y. Wang, T. Lei, J. Liu, Tribo-metallographic behavior of high carbon steels in dry sliding. I. Wear mechanisms and their transition, Wear 231 (1999) 1–11. [25] J.L. Sullivan, S.G. Hodgson, A study of mild oxidational wear for conditions of low load and speed, Wear 121 (1988) 95–106. [26] K. Kato, Classification of wear mechanisms/models, Proc. Inst. Mech. Eng. Part J: J. Eng. Tribol. 216 (2002) 349–355. [27] Y. Wang, T. Lei, J. Liu, Tribo-metallographic behavior of high carbon steels in dry sliding. III. Dynamic microstructural changes and wear, Wear 231 (1999) 20–37.
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Wear journal homepage: www.elsevier.com/locate/wear
Influence of temperature of sub-zero treatments on the wear behaviour of die steel D. Das a , K.K. Ray b , A.K. Dutta c,∗ a b c
Department of Metallurgy and Materials Engineering, Bengal Engineering and Science University, Shibpur, Howrah 711103, India Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India Department of Mechanical Engineering, Bengal Engineering and Science University, Shibpur, Howrah 711103, India
a r t i c l e
i n f o
Article history: Received 1 September 2008 Received in revised form 25 November 2008 Accepted 25 November 2008 Keywords: Sub-zero treatment Cryogenic treatment Die steel Wear behaviour Mode of wear Mechanism of wear
a b s t r a c t This study examines the influence of temperature of sub-zero treatment on the wear behaviour of AISI D2 steel. A series of dry sliding wear studies have been made under constant normal load at varying sliding velocities. Emphasis has been laid to understand the operative modes and mechanisms of wear by the estimation of specific wear rates and detailed characterizations of the worn surfaces, wear debris and subsurfaces with the help of scanning electron microscope (SEM) examinations coupled with energy dispersive X-ray (EDX) microanalyses. The obtained results unambiguously infer that lower the temperature of sub-zero treatment higher is the improvement in wear resistance. Wear resistance can increase by 1.5–125 times depending on sliding velocity while hardness increases only by 4.2% at the lowest temperature of sub-zero treatment (77 K) compared to the conventionally treated specimens. These results corroborate well with the reduction in retained austenite content associated with simultaneous increase in the amount of secondary carbide particles with lowering of sub-zero treatment temperature. The operative modes and mechanisms of wear are identified as either mild oxidative or severe delaminative, which depends on the temperature of sub-zero treatment and the sliding velocity of the wear test. © 2009 Elsevier B.V. All rights reserved.
1. Introduction It is well established that sub-zero treatment improves several mechanical properties of tool steels like wear resistance [1–10], hardness [5–12], fatigue resistance [12], toughness [13], etc., apart from improving dimensional stability of engineering components [14]. The merit of sub-zero treatment is popularly attributed to the transformation of retained austenite (␥R ) to martensite. As a consequence, sub-zero treatment is being commercially used for the improvement of wear resistance of tool/die steels. In a series of recent investigations [7–10], the present authors have examined the nature of phase transformation at deep cryogenic temperature and have studied their role on the wear behaviour of AISI D2 steel. The outcomes of these investigations have indicated that apart from the transformation of ␥R to martensite, nucleation and growth of secondary carbides play an important role. In addition, these studies have also shown that the degree of improvement in wear resistance of tool steels is dependent on the selection of wear test parameters.
∗ Corresponding author. Tel.: +91 33 2668 4561/62/63; fax: +91 33 2668 4564/2916. E-mail address: [email protected] (A.K. Dutta). 0043-1648/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2008.11.029
The process of sub-zero treatment consists of controlled cooling of conventionally hardened specimens to a selected sub-zero temperature followed by controlled heating of the specimens back to the ambient temperature for subsequent tempering treatment. The sub-zero treatment can be classified into three different temperature regimes: cold treatment (CT, ≥193 K), shallow cryogenic treatment (SCT, 193–113 K) and deep cryogenic treatment (DCT, 113–77 K). The influence of these temperature regimes on the wear characteristics of sub-zero treated steels has not been systematically examined so far. The major aim of this report is to bring forth the influence of temperature of sub-zero treatment on the wear behaviour of steels under different wear test conditions. Until recently, the accepted temperature for sub-zero treatment has been 193 K where dry ice can be used for cooling. However, the results of few recent studies [1–4] suggest that the temperature of sub-zero treatment should be <193 K in order to obtain the maximum improvement in mechanical properties of tool steels; the lowest temperature may be 77 K, the boiling temperature of liquid nitrogen at normal atmospheric pressure. Mohan Lal et al. [2] have shown that lower is the temperature higher is the improvement in wear resistance for D3 and M2 steels. In contrast, Moore and Collins [15] have suggested that there is an optimal temperature for sub-zero treatment of tool steel. Schiradelly and Diekman [16] have applied Taguchi Design of Experiment (DOE) to identify and to optimize the critical parameters of cryogenic treatment for a
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Table 1 Nominal composition of the investigated AISI D2 steel. Element
C
Mn
Si
S
P
Cr
Mo
V
Fe
Weight (%)
1.49
0.29
0.42
0.028
0.029
11.38
0.80
0.68
Balance
martensitic stainless steel. Using the analysis of variance (ANOVA) associated with the results of wear tests, these authors have shown that the most significant parameter is the soaking temperature (72% in contribution). The reported improvements in wear resistance of die/tool steels due to cryogenic treatment varies widely, for example from a few percent to a few hundred percent for the same material [9]. Such inconsistent reports are hindering the exploitation of the inherent potential of cryogenic treatment for its commercial exploitations. The lack of systematic studies of wear behaviour of cryogenically treated tool steels with proper identification of the operative modes and mechanisms of wear is mainly responsible for the reported wide scatter in the degree of improvement of wear resistance of tool steels by cryogenic treatment [8–10]. Therefore, it is necessary to compare the wear resistance of tool/die steels subjected to different sub-zero temperature treatments with proper identification of the operative modes and mechanisms of wear which is the focus of this study. 2. Experimental procedure 2.1. Material and treatments A commercial forged AISI D2 steel bar has been chosen for this investigation. The nominal chemical composition of the steel is presented in Table 1. Specimens of this steel have been subjected to conventional treatment (CONT), cold treatment (CT), shallow cryogenic treatment (SCT) and deep cryogenic treatment (DCT) in separate batches. The CONT consists of hardening and tempering; while in CT, SCT and DCT, an additional step of controlled sub-zero treatment with lowest quenching temperature (TLQ ) as 198, 148 and 77 K, respectively, has been incorporated in-between hardening and tempering treatments. In sub-zero treatment, conventionally hardened specimens have been cooled from ambient temperature to the selected TLQ , held there for ∼5 min for homogenization of temperature before heating back to the ambient temperature for subsequent tempering treatment. Uniform cooling and heating rates of ∼0.75 K min−1 were maintained in each of the sub-zero treatments. The hardening and single tempering treatments following ASM standard [17] have been carried out at 1297 K for 30 min and at 483 K for 120 min, respectively. In order to distinguish the specimens subjected to different TLQ , these have been ascribed with codes as shown in Table 2. These designations are followed in the subsequent discussions. One can note here that TLQ for CONT is 303 K, i.e., with out any sub-zero treatment. 2.2. Microstructural characterizations Microstructural examinations have been carried out on polished and picral-etched specimens using both optical and scanning elecTable 2 Different heat treatments and the sample codes. Sample code
CONT CT SCT DCT
1293 K, 30 min
Lowest quenching temperature, TLQ (K) 303 198 148 77
2.3. Hardness measurement The bulk hardness of the differently treated specimens has been measured by Vickers hardness tester using 60 kgf load. At least fifteen readings are considered for estimating the average value of hardness, and the estimated standard errors associated with the measured Vickers hardness numbers (VHN) were ±5. 2.4. Sliding wear tests Dry sliding wear tests have been performed following ASTM standard G99-05 by using a computerized pin-on-disc wear testing machine. Cylindrical specimens of 4 mm diameter and 30 mm length have been used as static pins, whereas, tungsten carbide coated En 35 steel disc (surface hardness HV ≈ 1750) of diameter 160 mm is selected as the rotating counter surface. The wear tests have been carried out using a fixed normal load (FN ) of 98.1 N (10 kgf) at different sliding velocities (SV ) of 1.00, 1.25 and 1.50 m s−1 and under ambient conditions of 298 K and 60% relative humidity. The rpm of the disc was adjusted for a selected track diameter (100–140 mm) for achieving a predetermined SV . The wear tests were continued either for sliding distance of ∼2000 m or for the duration that resulted into cumulative height loss of the pin specimen of ∼2 mm (equivalent to ∼25 mm3 of cumulative wear volume loss), whichever occurred earlier. Average wear rates have been estimated from the recorded cumulative height loss of the specimens with respect to sliding distance in the steady-state wear regime considering at least three test results under identical conditions. 2.5. Characterizations of the worn surfaces, wear debris and subsurfaces The worn surfaces at the end of the wear tests and generated debris collected in the steady-state wear regime have been examined under SEM along with the energy dispersive X-ray (EDX) microanalyses for all specimens at all wear test conditions. Some selected worn samples were sectioned perpendicular to the worn surface and parallel to the sliding direction using wire electrodischarge machining. The sectioned faces of the specimens were mounted, polished to 0.25 m finish, etched using picral solution and were examined using SEM in back-scatter mode to record the subsurface features of the worn specimens. 3. Results and discussion 3.1. Microstructures and hardness
Description of heat treatment cycles Hardening
tron microscopes (SEMs). The carbide particles have been classified as primary carbides (PCs) and secondary carbides (SCs) considering equivalent spherical diameter for PCs as >5 m, while that for SCs ≤ 5 m. The volume percents of PCs and SCs have been estimated by image analyses of digitally acquired suitable optical and SEM micrographs using Leica QMetals software in Leica QWin V3 environment. Identification of the phases, and the measurement of retained austenite (␥R ) following ASTM standard E975-00 have been done by X-ray diffraction (XRD) analyses with Mo K␣ radiation at 0.01◦ s−1 scan rate on bulk specimens with the help of PHILIPS X’Pert software.
Tempering
483 K, 120 min
Fig. 1 depicts representative SEM micrographs of conventionally treated and a typical sub-zero treated specimens. These exhibit large elongated dendritic type of primary carbides (PCs) and nearly spherical secondary carbides (SCs) on tempered martensite matrix. The occasional patches of retained austenite (␥R ) have been detected only in SEM micrographs of CONT [7] and CT specimens. The volume fraction of ␥R has been estimated by XRD technique,
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Fig. 1. Typical representative SEM micrographs of (a) CONT and (b) SCT specimens exhibiting primary carbides (PCs) and secondary carbides (SCs) of different sizes. The smaller SCs are shown inside circles.
detailed procedure of which has been reported earlier [9]. The amounts of ␥R in the different types of specimens are shown in Fig. 2. The amount of ␥R in CT specimen is 4.6 ± 0.5 vol.% in comparison to 9.8 ± 0.7 vol.% in CONT specimen. The amount of ␥R in both SCT and DCT specimens is considered to be below the detection limit (<2 vol.%) of XRD technique. This observation is in agreement with the reported results for AISI D2 steel in DCT [6–10] and SCT [14] conditions. The obtained results suggest that CT reduces the ␥R substantially as compared to CONT; while both SCT and DCT in between conventional hardening and tempering treatments almost completely remove it. Since the martensite finish (Mf ) temperature for AISI D2 steel is ≈148 K, SCT and DCT with TLQ ≤ 148 K almost completely transforms ␥R to martensite, and hence the observed results are expected. The amount of PCs is found to be invariant (≈7.0 vol.%) while the amount of SCs varied in CONT, CT, SCT and DCT specimens. The PCs in the investigated specimens have been identified mainly as M7 C3 (M = Fe, Cr, Mo, V) with small amount of Cr7 C3 by XRD analyses. The variations in the amounts of SCs with TLQ as determined by image analyses are illustrated in Fig. 2. The SCs in all the specimens have been identified as M23 C6 (M = Fe, Cr, Mo, V) type and this is in agreement with earlier reports [7–10,13]. It is interesting to note that the application of sub-zero treatments after conventional hardening does not alter the nature of PCs and SCs. The amount of SCs increases with decreasing temperature of sub-zero treatment. As the temperature of sub-zero treatment is lowered, the amount of ␥R
decreases resulting in higher amount of martensite; the increased amount of martensite naturally leads to higher amount of carbide precipitation on subsequent tempering treatment as observed in Fig. 2. The variation of hardness with TLQ is shown in Fig. 3. The results indicate that sub-zero treatment increases hardness of D2 steel though by a smaller margin of up to 4.2%. The increase in hardness is attributed to reduction in ␥R content with associated increase in the amount of SCs (Fig. 2). It is observed from the results in Fig. 3 that the rate of increase in hardness with TLQ diminishes with lowering of TLQ . In addition, the nature of distribution of SCs (Fig. 1) can also influence the increment in hardness with TLQ due to alteration in the inter-particle spacing leading to the variation in micro-residual stresses and sub-structural changes like dislocation density.
Fig. 2. Variation of amount of retained austenite (␥R ) and secondary carbides (SCs) as a function of lowest quenching temperature.
Fig. 3. Influence of lowest quenching temperature on the bulk hardness of differently treatment specimens.
3.2. Wear rates Typical plots of cumulative wear volume loss versus sliding distance for all the specimens tested at SV of 1.25 m s−1 under FN of 98.1 N are shown in Fig. 4. The results in Fig. 4 exhibit distinguished regimes of ‘running-in’ and ‘steady-state’ wear [18–20] for each specimen and the difference in the nature of variation in the wear behaviour of these specimens. The results exhibit that the volume losses for CONT and CT specimens are significantly higher than those of the SCT and DCT specimens. Furthermore, volume
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Fig. 4. Typical plot of cumulative wear volume loss (CWVL) versus sliding distance (SD ) for differently treated specimens tested at sliding velocity of 1.25 m s−1 . Inset shows enlarged view of same for CONT and CT specimens. The arrow indicates the transition of initial running-in wear to the final steady-state wear.
loss of CONT specimen is appreciably higher than that of CT specimens, whereas the volume loss for SCT specimens is marginally higher than that of DCT specimen. The results obtained for tests conducted at SV of 1.00 and 1.50 m s−1 is similar in nature. The extent of running-in wear depends on the experimental conditions like (i) surface roughness of both the pin and the disc, (ii) perpendicularity of the pin with respect to the disc and (iii) the rise in temperature at the contact area, etc. [20]. The results in Fig. 4 indicate that (i) the wear rates of all specimens are considerably higher in the runningin regime than that in the steady-state regime and (ii) the wear rates in the running-in regime decreases with lowering of TLQ . Steady-state wear, which is known to depend on the material’s characteristics [18–20], has been analyzed to reveal the effect of TLQ on the wear behaviour of the selected steel. For this purpose, the specific wear rate (WS ) has been calculated by the equation: WS = VL /(FN × SD ); where VL is the wear volume loss of the pin specimens in the steady-state regime in mm3 , FN is the applied normal load in N and SD is the sliding distance corresponding to the steady-state wear regime in mm. The magnitudes of WS estimated at different SV for specimens with different TLQ are compiled in Fig. 5. The results in Fig. 5 assist to infer that: (i) at SV = 1.00 m s−1 , WS sharply decreases for TLQ = 303–198 K and then its magnitude decreases marginally with further lowering of TLQ , (ii) at SV = 1.25 m s−1 , WS marginally decreases between TLQ = 303 and 198 K and between TLQ = 148 and 77 K, whereas WS drastically decreases between TLQ = 198 and 148 K, and (iii) at SV = 1.50 m s−1 , the magnitude of WS marginally decreases for the entire range of investigated TLQ . The nature of variation of WS for different specimens at the selected SV or for the same specimen at different SV can be explained only with the knowledge of the operative mechanisms of wear and this is the content of the next section. 3.3. Mechanisms of wear In order to identify the operative wear mechanisms, the topography of the worn surfaces as well as the nature and morphology of the generated wear debris have been characterized for all the tested specimens. Some typical representative examples in Figs. 6–10
Fig. 5. Variation of estimated specific wear rate (WS ) with lowest quenching temperature (TLQ ) of the specimens tested at different sliding velocities (SV ). Shading has been used to demarcate WS into two different regimes based on the operative modes and mechanisms of wear (see text for details).
illustrate the operative wear mechanisms of differently treated specimens tested at varying SV . The micrographs in Figs. 6 and 7 bring to light the difference in the operative wear mechanisms for CONT specimens compared to that for different sub-zero treated specimens at SV = 1.00 m s−1 . The worn surface of CONT specimen is rough but metallic in nature, and also exhibits fracture ridges and deformation lips stretched parallel to the sliding direction (Fig. 6a). The presence of deformation lip infers that the CONT specimen has undergone heavy plastic deformation during wear test. In contrast, the worn surfaces of the CT (Fig. 6b), SCT (Fig. 6c) and DCT (Fig. 6d) specimens are smoother and exhibit presence of patches of oxide. The micrograph of oxide depleted region of worn surfaces of these specimens shows cracking and pull-out of PCs apart from the presence of surface grooves as illustrated in Fig. 6e. Surface grooves on the worn surface are formed due to the trapping of loose hard PC particles in-between the pin and disc surfaces. The micrograph of the wear debris of CONT specimen (Fig. 7a) exhibits shiny metallic appearance and a plate-like morphology. In contrast, the debris of the CT (Fig. 7b), SCT (Fig. 7c) and DCT (Fig. 7d) specimens appear as black coloured powder in naked eye; and these are much finer in size and oxide in nature (Fig. 7e). These observations reveal that the wear mechanism at SV = 1.00 m s−1 and FN = 98.1 N for CONT specimen are plastic deformation induced delaminative wear [21], whereas that for the sub-zero treated specimens is predominantly oxidative [22] coupled with cracking and pull-out of PCs. Representative features on the worn surfaces of CONT, CT, SCT and DCT specimens tested at SV = 1.25 m s−1 are shown in Fig. 8; the corresponding morphology of the wear debris and their typical representative EDX profiles are presented in Fig. 9. The observed features on the worn surfaces for both CONT and CT specimens tested at SV = 1.25 m s−1 (Fig. 8a and b) closely resemble with that observed for CONT specimen tested at SV = 1.00 m s−1 (Fig. 6a). Worn surface features for SCT and DCT specimens tested at SV = 1.25 m s−1 (Fig. 8c and d) are similar to that of CT, SCT and DCT specimens tested at SV = 1.00 m s−1 (Fig. 6b–d). However, the worn surfaces of SCT and DCT specimens show regions of thick compacted oxide layer at SV = 1.25 m s−1 (Fig. 8c and d) unlike that in the worn surface of SCT and DCT specimens tested at SV = 1.00 m s−1 (Fig. 6c and d). Furthermore, examinations of worn surfaces of SCT and DCT specimens tested at SV = 1.25 m s−1 does not reveal any obvious signature of cracking and/or pull-out of PCs even at higher magnifications.
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Fig. 6. Representative secondary electron SEM micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.00 m s−1 . (e) A typical high magnification backscatter electron SEM micrograph of oxide depleted region as marked by area-1 in the worn surfaces of CT, SCT and DCT specimens.
The wear debris of CONT and CT specimens show large metallic platelets unlike fine oxide debris shown by SCT and DCT specimens (Fig. 9). Typical representative EDX profile of the wear debris of CONT and CT specimens is shown in Fig. 9e and that for SCT and DCT specimens is presented in Fig. 9f. The recorded EDX profiles exhibit oxygen peak of higher intensity for SCT and DCT specimens compared to that from CONT and CT specimens (Fig. 9f vis-à-vis Fig. 9e). This observation suggests that the generated wear debris for CONT and CT specimens is predominantly metallic in nature and that for SCT and DCT specimens is predominantly oxide in nature. The EDX profiles of both types of wear debris indicate the presence of tungsten that has been transferred from the counter body to the specimen surface by adhesion. Evidences in Figs. 8 and 9 assist to reveal that the operative mechanisms of wear for both CONT and CT specimens is delaminative at SV = 1.25 m s−1 , whereas that for SCT and DCT specimens is oxidative wear under same experimental condition. The surface topography of all specimens tested at SV = 1.50 m s−1 indicates delaminative wear. Representative evidences are shown
in Fig. 10 for SCT and DCT specimens. The wear debris generated during wear tests of these specimens is metallic platelets (inserts in Fig. 10) corroborating with the observed features of the worn surfaces. In summary, the observed wear mechanism for CONT specimens for all tested SV is delaminative wear; whereas the operative wear mechanisms change from predominantly oxidative to delaminative at SV = 1.25 m s−1 for CT specimens and at SV = 1.50 m s−1 for SCT and DCT specimens. 3.4. Transition of modes and mechanisms of wear The modes of wear may be categorized using wear mechanism map as suggested by Lim and Ashby [23] or by magnitude of WS as suggested by Wang et al. [24] or through magnitude of wear coefficient as suggested by Rabinowicz [20]. Identification of the operative wear mechanisms at different test conditions is known to assist in revealing the operative mode of wear [23]. In the present investigation, the operative mechanisms of wear have been identified by the detailed characterizations (for all specimens tested
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Fig. 7. Representative secondary electron SEM micrographs of generated wear debris corresponds to the steady-state wear regime of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.00 m s−1 . All these micrographs are taken at the same magnification of 250X, whereas (e) is a typical high magnification view of wear debris of CT, SCT and DCT specimens.
at different SV ) of worn surfaces and generated debris. These are shown in Fig. 5, which reveals that the data related to the WS clusters into two different regimes depending on the mechanisms of wear, and the magnitude of WS between these regimes varies about an order of magnitude. Considering the magnitude of WS and the wear mechanisms, these regimes are termed as ‘mild-oxidative’ and ‘severe-delaminative’, in which ‘mild’ and ‘severe’ refer the mode of wear. The results in Fig. 5 indicate that the transition SV for mild-tosevere wear for CT is 1.00 m s−1 and that for SCT and DCT specimens are 1.25 m s−1 . Within the tested range of SV (1.00–1.50 m s−1 ), the WS for CONT specimens lies in the severe wear regime (Fig. 5) and corresponding wear mechanism has been identified as delaminative wear (Figs. 6–9). These results suggest that transition SV for mild-to-severe wear for CONT specimens occurs at SV lower than 1.00 m s−1 . Additional experiments were carried out to find out the transition SV for CONT specimens and this has been found to be 0.75 m s−1 corresponding to predominantly oxidative wear. The WS of the CONT specimen at the transition SV is 8.10 × 10−8 mm3 N−1 mm−1 . The transition SV for mild-to-severe wear is a function of microstructure and mechanical property of tested pin specimens, when the other wear test parameters remain constant [23,24]. The observed variation in the transition SV for CONT, CT and SCT specimens is expected, since these specimens exhibit significant differences in their microstructures (Fig. 2) and hardness values
(Fig. 3). The obtained transition SV for both SCT and DCT specimens are same (1.25 m s−1 ), because the microstructures of SCT and DCT specimens are nearly similar, both being free from ␥R and having almost identical amounts of SCs (Fig. 2). In addition, the difference in hardness values of these specimens is <0.5%. Hence, it is natural that both SCT and DCT specimens exhibit same transition SV . It is also interesting to note that the WS of both SCT and DCT specimens decreases about an order of magnitude with increase of SV from 1.00 to 1.25 m s−1 even when the operative mode and mechanism of wear (mild-oxidative) remain unchanged. Similar observation has been reported earlier for dry sliding wear of steel specimens; for example, Sullivan and Hodgson [25] have shown that the wear rate of 52100 steel decreases by an order of magnitude with increase in SV , when SV lies below transition SV for mildto-severe wear. This behaviour is related to the rate of formation and retention of oxide at the contact surface of the pin specimen. Increase in SV increases the flash temperature that, on one hand, accelerates the formation of tribochemical reactions at the contact junctions [22,23]. On the other hand, it increases the degree of work-hardening of the material beneath the contact junctions leading to the development of increased depth of work-hardened layer with enhanced hardness [25]; this, in turn, assists in retaining oxides at the contact junctions resulting into development of thicker oxide layer [22,25]. Therefore, increased SV increases the thickness of oxide layer at the contact surfaces which is evident
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Fig. 8. Representative secondary electron SEM micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.25 m s−1 .
from the SEM micrographs of wear debris of both SCT and DCT specimens (Fig. 7 vis-à-vis Fig. 9). Development of thicker oxide layer is thus responsible for the reduction in the WS values with increasing SV from 1.00 to 1.25 m s−1 for both SCT and DCT specimens [22,25]. The results in Fig. 5 assist to infer that the maximum value of WS for mild-oxidative wear and the minimum value of WS for severe-delaminative wear are 9 × 10−8 and 4 × 10−7 mm3 N−1 mm−1 , respectively. Wang et al. [24] have earlier suggested for steel specimens that the upper limit of WS for mild-oxidative wear is 1 × 10−8 mm3 N−1 mm−1 and the lower limit of WS for severedelaminative wear is 2 × 10−8 mm3 N−1 mm−1 . These authors have also indicated that one can use the term ‘transition wear’, if the WS is in the region of 1 × 10−8 to 2 × 10−8 mm3 N−1 mm−1 . The present results, however, indicate that the demarcating lines for transition wear could be wider than that suggested by Wang et al. [24]. Kato [26] has earlier suggested that the WS of steel specimens
can vary over a wide range, from 10−5 to 10−13 mm3 N−1 mm−1 , and the exact magnitude of WS are influenced by several statistical variables involved in the wear tests. The major variables that determine the magnitude of WS are macro-material structures, surface roughness, flash temperature, local contamination, adhesive transfers, free wear particles and tribochemical reactions on the micro-scale on the contact surfaces [26]. Thus, the differences in the magnitudes of WS as obtained in the present investigation and that reported by Wang et al. [24] for demarcating the mild and the severe wear regimes may be attributed to the differences in the microstructures and properties of the selected materials, employed wear test methodologies and the chosen test conditions. Wang et al. [24] have studied the wear behaviour of 52100 and 1080 steels employing pin-on-ring testing method, whereas the material and test methodology of the present study are AISI D2 steel and pin-ondisc.
Table 3 Ratio of specific wear rates (WS ) and the operative modes and mechanisms of wear. Sliding velocity (m s−1 )
Specimen codes
Wear rate ratio, a
Mode and mechanism of wear Mode
Mechanism
1.00
CONT CT SCT DCT
1 6.5 7.3 7.9
Severe Mild Mild Mild
Delaminative Oxidative Oxidative Oxidative
1.25
CONT CT SCT DCT
1 1.8 103 125
Severe Severe Mild Mild
Delaminative Delaminative Oxidative Oxidative
1.50
CONT CT SCT DCT
1 1.2 1.3 1.5
Severe Severe Severe Severe
Delaminative Delaminative Delaminative Delaminative
a
is ratio of WS of CONT specimen to that of CT, SCT and DCT specimens.
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Fig. 9. Representative SEM micrographs of generated wear debris corresponds to the steady-state wear regime of (a) CONT, (b) CT, (c) SCT and (d) DCT specimens tested at SV = 1.25 m s−1 . All micrographs taken at the same magnification of 250×, whereas the insets in (c) and (d) at 1500× are detailing the features of the same micrographs. (e) and (f) are representative EDX profiles taken from the rubbed surfaces of the wear debris as marked by area-1 in (a) and (b), and by area-2 in the inset of (c) and (d), respectively.
Fig. 10. Representative SEM micrographs of worn surfaces at the end of wear tests of (a) SCT and (b) DCT specimens tested at SV = 1.50 m s−1 . Insets represent the generated wear debris corresponds to the steady-state wear regime for the same samples.
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Fig. 11. Representative back scattered SEM subsurface micrographs of worn surfaces at the end of wear tests of (a) CONT, (b) CT, (c) SCT and (b) DCT specimens tested at SV = 1.25 m s−1 . WL: while layer; PDL: plastically deformed layer.
3.5. Influence of temperature of sub-zero treatments on wear behaviour In order to assess the influence of temperature of sub-zero treatments on the degree of improvement in wear resistance, a new parameter has been incorporated here which is considered as the ratio of WS of CONT specimen to that of the different sub-zero treated specimens at a given SV . The estimated values of for different SV are compiled in Table 3 along with the operative modes and mechanisms of wear. The results in Table 3 reveal that the improvement in wear resistance by sub-zero treatment is marginal when the operative mode and mechanism between CONT and sub-zero treated specimens are similar (severe-delaminative). Under this condition, the value of is 1.8 for CT specimen at SV = 1.25 m s−1 and it varies from 1.2 to 1.5 at SV = 1.50 m s−1 for all the sub-zero treated specimens. In contrast, the magnitude of varies over two orders of magnitude when the operative mode and mechanism of wear for CONT specimens is severe-delaminative and that for the sub-zero treated specimens is mild-oxidative. Such wide variation in the magnitudes of is due to the similar or dissimilar operative modes and mechanisms of wear amongst the specimens considered (Table 3). These findings are in excellent agreement with the recent reports by the present authors [8,10] on the deep cryogenically treated and conventionally treated tool steels. However, comparison of values for SCT and DCT specimens at SV = 1.00 and 1.25 m s−1 indicates that the increases significantly at higher SV for these specimens, though the operative mode and mechanism of wear for the CONT (severe-delaminative) and that for the SCT and DCT (mild-oxidative) remains unchanged for above-mentioned SV values (Table 3). Such variation in occurs due to increase in WS for the CONT specimens and reduction of the same for the SCT and DCT specimens when SV increases from 1.00 to 1.25 m s−1 (Fig. 5). This apparent contradiction, however, is in agreement with earlier suggestions by Welsh [18,19]. He elucidated that the WS of steels increases with increasing SV, if the SV lies between the characteristics critical SV related to the mildto-severe and the severe-to-mild wear transitions. Identification of operative wear mechanism as delaminative for CONT specimen for
all the tested SV (1.00–1.50 m s−1 ) confirms that these SV indeed lie in between the two transitions SV for this specimen. Thus, according to this explanation of Welsh [18,19] it is natural that WS would increase with the increase in SV from 1.00 to 1.25 m s−1 as observed for the CONT specimens (Fig. 5). Welsh [18,19] and Sullivan and Hodgson [25] have earlier also shown that the WS of steels tend to decrease with increasing SV , if the applied SV is lower than the transition SV at which mildto-severe transition takes place. For both SCT and DCT specimens, mode and mechanism of wear are ‘mild-oxidative’ at SV = 1.00 and 1.25 m s−1 and ‘severe-delaminative’ at SV = 1.50 ms−1 . Thus, SV = 1.00 and 1.25 ms−1 are below the transition SV for these specimens. Therefore, it is natural that WS of these specimens would decrease with the increase of SV from 1.00 to 1.25 m s−1 (Fig. 5). This behaviour is related to the rate of tribochemical reaction at the contact surfaces as the increase in SV increases the flash temperature, which accelerates the formation of oxide layer at the contact junctions. At higher SV , development of more compacted oxide layer at the contact is expected to reduce the magnitude of WS as observed in the present study for the SCT and DCT specimens when SV increased from 1.00 to 1.25 m s−1 . This is well supported from the observations that (i) the area of worn surface covered by the compacted oxide layer increases with increasing SV (Fig. 6c and d vis-à-vis Fig. 8c and d) and (ii) the generated oxide debris are much thicker at higher SV (Fig. 7e vis-à-vis insets in Fig. 9c and d). An attempt to correlate the results of wear tests in Table 3 with the results of microstructural analyses in Figs. 1 and 2 reveal that TLQ has considerable influence on generating the nature of the microstructures and hence on the resulted wear properties of the selected steel. In order to get insights on the role of microstructural constituents on the variation of wear behaviour, the subsurfaces of all types of specimens tested at SV = 1.25 m s−1 were examined under SEM. Typical micrographs of the subsurfaces are presented in Fig. 11. The features beneath the worn surface of CONT and CT specimens are markedly different than that for SCT and DCT specimens (Fig. 11). The SCT and DCT specimens show only plastically deformed layer; in contrast, the CONT and CT specimens exhibit a well-defined ‘heavily deformed layer’, termed popularly
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in the literature as white (non-etching) layer [27], and followed by a plastically deformed layer. The presence of the white layer is the signature of deformation induced delaminative wear [21] which is indeed the governing wear mechanism for CONT and CT specimens at SV = 1.25 m s−1 . The presence of only plastically deformed layer and the absence of any white layer in the subsurfaces of SCT and DCT specimens are indicative of mild-oxidative wear for these specimens. The white layer of CONT specimen (Fig. 11a) shows presence of cracks that originated at a given depth beneath the worn surface and propagated parallel to the sliding direction; whereas, the cracking of white layer is less prominent for CT specimen (Fig. 11b). This observation is in agreement with the lower estimated WS of the CT specimen than that of CONT specimen at identical test conditions (Fig. 5). Severe-delaminative wear of steels occurs by continued plastic deformation at the contact surfaces of a test specimen leading to nucleation of cracks followed by its propagation as suggested by Suh [21]. The presence of soft ␥R associated with lower amount of hard SCs in the CONT and CT specimens as compared to those in the SCT and DCT specimens (Fig. 2) suggests that the former specimens are more prone to plastic deformation during wear at a lower SV for a constant FN or at lower FN at a constant SV . Therefore, specimens having less resistance to plastic deformation are expected to undergo mild-to-severe wear transition at a lower FN and/or SV . Thus, the present observation of higher SV for mild-tosevere wear transition for SCT and DCT specimens as compared to that for CONT and CT specimens is natural. Moreover, the fact that at SV = 1.00 m s−1 , the operative wear mechanisms for CONT and CT specimens are delaminative and oxidative wear, respectively, also supports the above argument; since CT specimen possesses higher amount of SCs and lower amount of ␥R than those in CONT specimens (Fig. 2). In summary, this investigation highlights the influence of temperature of sub-zero treatments on the microstructure and wear behaviour of die steel. 4. Conclusions The experimental results related to the influence of temperature of sub-zero treatments on the wear behaviour of selected AISI D2 steel and their pertinent analyses assist to infer the following major conclusions: (i) All types of sub-zero treatments appreciably improve the wear resistance of the die steels compared to the conventional treated (CONT) ones. However, the improvement in wear resistance by shallow (SCT) and deep cryogenic treatment (DCT) is significantly higher than that achieved by CT, and the maximum improvement is obtained by DCT. This is attributed to the decrease in the retained austenite content associated with the increasing amount of secondary carbide particles with lowering of temperature of the sub-zero treatments. (ii) The obtained hardness of AISI D2 steel for CONT and DCT are 759 and 791 VHN, respectively and typical values of their specific wear rates (WS ) are 1.03 × 10−6 and 8.26 × 10−9 mm3 N−1 mm−1 , respectively, at sliding velocity (SV ) of 1.25 m s−1 . (iii) Estimation of WS and detailed characterization of worn surfaces, generated debris and subsurfaces of all types of specimens assist to reveal that the severe mode of wear is associated with the deformation induced delaminative wear mechanism, whereas the mild mode of wear is associated with the predominantly oxidative wear mechanism coupled with occasional cracking and pull-out of primary carbide particles. (iv) The modes and mechanisms of wear are synergistically affected by the temperature of sub-zero treatments and the
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