Tribol Lett (2009) 34:155–166 DOI 10.1007/s11249-009-9421-y

ORIGINAL PAPER

Wear Mechanisms at High Temperatures. Part 1: Wear Mechanisms of Different Fe-Based Alloys at Elevated Temperatures H. Winkelmann Æ E. Badisch Æ M. Kirchgaßner Æ H. Danninger

Received: 30 December 2008 / Accepted: 16 February 2009 / Published online: 5 March 2009 Ó Springer Science+Business Media, LLC 2009

Abstract To extend the lifetime of the sinter grate used to crush the sinter cake into smaller pieces for steel fabrication, a study was undertaken to investigate which wear processes are primarily responsible for limiting the lifetime of the sinter grate. Several wear processes could be identified. The sinter temperature which is up to 800 °C causes temperature-induced material ageing and oxidation. The falling of the sinter cake onto the sinter grate causes high impacts, erosion and abrasive wear. There is enormous economic pressure, which makes the most cost-efficient solution the most attractive one, not the technically ‘‘best’’ coating material; thus, Fe–Cr–C hardfacing alloys are mostly used. In view of the above, four different alloys which are promising for this application were studied with regard to their wear resistance. Each wear mechanism was investigated in a special test tribometer. Fatigue wear caused by multiple impacts and abrasion was tested in the high-temperature continuous impact abrasion test. Materials behaviour in heavy single impacts was evaluated in the single impact test. Characterisation of microstructure and wear behaviour was performed by optical microscopy and scanning electron microscopy. The results obtained with the help of the different measurement techniques were linked and set into comparison to calculate the volumetric

H. Winkelmann (&)
E. Badisch AC2T Research GmbH, Viktor Kaplan-Straße 2, 2700 Wiener Neustadt, Austria e-mail: [email protected] M. Kirchgaßner Castolin GmbH, Brunner Straße 69, 1230 Vienna, Austria H. Danninger Institute of Chemical Technologies and Analytics, TU Wien, Getreidemarkt 9, 1060 Vienna, Austria

wear of the specimen. Aim of this work was to investigate the influence of the material parameters such as macrohardness, hard phase content, microstructure coarseness on the wear resistance in impact loading and abrasive applications at high temperatures. Results also indicate that the matrix ability to bind carbides at high temperature as well as the matrix hardness at high temperatures strongly influence the wear resistance in the different tests. Those material parameters get correlated to the wear rates in different material demands. The test results indicate that at higher temperatures material fatigue becomes a major wear-determining factor which makes the matrix hardness and the matrix ability to bind carbides at high temperatures very important. Especially, in abrasive wear, a certain content of hard phases is also necessary to keep the wear to a lower level. It could also be shown that in impact loading applications, a coarse microstructure is a disadvantage. Keywords Abrasion

High temperature
Wear
Impact


1 Introduction Iron ore sintering has stood at the heart of the ferrous metallurgical processes for over half a century. The cooled sinter is crushed to a pre-determined maximum particle size by using a sinter grate (Fig. 1). Undersized sinter (\15–25 mm) that is not suitable for the blast furnace is recycled to the return fines bin and re-circulated to the sinter machine. The lifetime of a sinter grate is very important for the smooth running of the sinter plant. The sinter grate encounters various types of wear loading, especially abrasion, impact and high-temperature erosion. An optical image of a sinter grate and the temperatures

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Fig. 1 Sinter grate and sinter crusher. a Optical image. b Thermograph study of real temperatures at work (°C)

attained in service are shown in Fig. 1. The sinter cake falls 4 m deep onto the sinter grate and the rotating sinter crusher pushes the sinter through the sinter grate. This leads to three different demands on the materials which are used in the sinter grate. First, impact resistance to high single loading is necessary because of the high impact energy of the 3 tonnes of sinter falling onto the sinter grate every 30 s. Second, abrasion resistance is required because the sinter crushed through the sinter grate causes constrained abrasion. And, third, temperature resistance and oxidation resistance are required for extended lifetime because of the high temperatures prevailing in the sinter grate (Fig. 1b). Various materials have been studied to increase the lifetime of the sinter grate. In general, various iron-based alloys have been used for this purpose. Mechanical and abrasive wear properties of the alloys depend on their microstructure and chemical composition [1–12]. To achieve impact resistance, a good ductility and good interfacial carbide–matrix bonding is necessary [13]. Therefore, for high-impact application, martensitic materials are best suited [14]. For abrasion resistance, hard phases and high hardness are important and, especially, it is important to ensure that the hardness of the hard phases and/or the hardness of the matrix is higher than the hardness of the abrasive [15–18]. For temperature resistance and oxidation resistance, a high alloyed matrix, especially, austenites are best suited [19, 20]. In view of the above, various materials have been tested to suite the conditions of the sinter grate and to find correlations between high temperature wear, hard phase content, different kinds of matrix and coarseness of the microstructure.

2 Experimental 2.1 Materials and Characterisation Within this study, an austenitic stainless steel (material A), a standard ingot metallurgy M2 tool steel (HSS) (material

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B), a complex Fe–Cr–C–Nb–Mo–W–B alloy with fine microstructure (material C) and a hypereutectic Fe–Cr–C– Mo–Nb alloy with coarse microstructure (material D) were investigated. All the materials and their chemical composition as well as their hardness and hard phase content are given in Table 1. Typical microstructures of the alloys are shown in Fig. 2. Materials A and B in opposite to materials C and D have been chosen to compare the influence of high to low hard phase content. By comparing materials A and B, the influence of the martensitic and austenitic matrix was compared. Materials C and D have been compared to know how the coarseness of the microstructure influences the wear. It was expected that materials A and B behave ‘‘better’’, meaning less break outs in the single impact test (SIT), while materials C and D behave ‘‘better’’, meaning lower wear rates in the high-temperature continuous impact abrasion test (HT-CIAT). Characterisation of microstructure was performed by optical microscopy (OM) after etching and scanning electron microscopy (SEM ? EDS). Quantitative analysis of the microstructure was carried out with the help of Intronic Image C Software. Hardness measurements were carried out with a standard Vickers hardness technique HV5. To determine the hardness of each phase in the microstructure, e.g. hard particles and metallic matrix, microhardness HV0.1 was used. Quantitative wear characterisation was done by gravimetric mass loss of the testing specimens during wear testing. Qualitative characterisation of worn surfaces was carried out by evaluating macroscopic and cross-section images, as described above, and by SEM investigations. Alloy A (Fig. 2a) is a heat-resistant austenite steel with a C content of 0.08%, Cr content of *25% and *20% Ni. Hardness of this alloy was determined as 175 HV5. Austenitic stainless steels have high ductility, low yield stress and relatively high ultimate tensile strength, when compared to typical carbon steel. Austenitic steels showing fcc atomic structure, which provides more planes for the flow of dislocations, when combined with the low level of interstitial elements gives this material a good ductility.

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Table 1 Chemical composition and hardness of the Fe-based alloys investigated

Carbide content (%)

Low hard phase content

High hard phase content

A Austenite 1.4841

B Tool steel 1.3343

C Complex Fe–Cr–C–Nb–Mo–W–B alloy; fine microstructure

D Hypereutectic Fe–Cr–C–Mo–Nb alloy; coarse microstructure

0

15–20

52

57

Chemical composition (wt%) Fe

Base

Base

Base

Base

C

0.08

0.9

1.3

5.5

Cr

24.8

4.1

15.4

21.0

Ni

19.8

–

–

–

Si

1.7

0.25

0.5

0.8

Mn

1.2

0.3

0.2

0.2

Nb

–

–

4.2

7.0

B

–

–

4.2

–

Others (Mo, V, W)

–

13.2

11.5

10.0

175

880

1020

880

Hardness (HV5)

Fig. 2 Microstructure of Febased alloys investigated. a A: austenite. b B: M2 tool steel. c C: complex Fe–Cr–C–Nb– Mo–W–B. d D: hypereutectic Fe–Cr–C–Mo–Nb

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Alloy B is a standard M2 high-speed steel (HSS). Figure 2b shows a martensitic microstructure with fine primary hard phases. Total content of carbides is approximately 15–20%. The hardness of this alloy is measured as 880 HV5. Alloy C shows a dense and uniform distribution of very hard complex carbides and carbo-borides (Fig. 2c) with hardness values between 1200 and 1900 HV0.1. The hard phases were identified as Fe/Cr carbo-borides with a volume content of 52% and a size of 10–100 lm, Nb carbides and Mo/W carbo-borides with a volume content of approximately 5% in blocky shape [14]. In addition, the hardness of the matrix is also very high which is reflected by a hardness of 1020 HV5. Alloy D consists of primary Fe/Cr carbides with a microhardness of roughly 1600 HV0.1 in a ledeburitic matrix (Fig. 2d). The content of Fe/Cr carbides is 57.1% with size of 30–200 lm [14]. The chemistry of the Fe/Cr carbides is reported for hypereutectic Fe–Cr–C alloys in literature to be M7C3 structure [21–24]. The hardness values of the ledeburitic matrix, which are determined to be about 800 HV0.1, are close to the results given by Fischer and Buytoz [25, 26]. In addition, small and evenly distributed primary Nb carbides with a volume content of approximately 5% can be detected. These are supposed to be of major importance for increasing the resistance against erosion and abrasion due to their high hardness. The hardness of material D is 880 HV5. 2.2 Single Impact Test The SIT was developed at AC2T (Austrian Center of Competence for Tribology) to characterise the impact resistance of materials to single impacts with high energies up to 80 J. The test principle is based on a falling ‘‘hammer’’ (m = 10 kg) with defined kinetic energy hitting the sample with a sharp edge 5° inclined to the sample’s surface (Fig. 3). This impact causes an impact mark on the sample and this deformation is analysed. The impact energy is regulated by the free falling height; in this experimental series, the energy was set to 7 J for all tests. The samples are fixed, so that they can not dodge during impact and thereby distort the results. Every material is tested using three specimens with three impacts each to enable some statistics. The impact marks are quantified primarily by measuring their length in the optical microscope. Additionally, the depth of the impact mark is measured to calculate the angle of the resulting impact. This angle provides information about the relationship between elastic and plastic deformation of the material. These analyses have been made by means of OM using stack images to get focused images of the entire depth of the impacts. The length is measured by the parallel distance from the intact

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sample edge to the end of the impact mark. The depth was equally quantified by measuring from the intact sample edge to the peak of the impact mark. Resulting impacts on ductile materials show a bulge which distorts the calculated impact angle. To calculate an accurate angle, collaterally the size of this bulge is measured and considered in the calculation. In every brittle materials break outs may occur, which make the measurement of the impact length impossible. These materials can not be compared with the others through this method; however, critical impact energies can be detected where the cracks start occurring. Thereto additionally impact energies where increased from 1 to 30 J to detect the crack initiation energy. 2.3 High-Temperature Continuous Impact Abrasion Test The HT-CIAT was developed at AC2T to determine the behaviour of the materials in continuous impact abrasive environment at elevated temperatures (Fig. 4a). Test principle is simply based on the potential energy which is cyclic turned into kinetic energy by free fall. The samples are fixed in 45° and get continuously hit by the plunger, while a constant abrasive flow is running between the sample and the plunger as shown in Fig. 4b and c. The testing parameters are summarised in Table 2. Impact energy, angle of impact and frequency were chosen as 0.8 J, 45° and 2 Hz, respectively. The total number of testing cycles was fixed to 7,200 which correlate to a testing duration of 1 h. The abrasive material used for three-body contact was silica sand with a particle size of 0.4–0.9 mm with angular shape and a flow rate of 3 g/s. Experiments were carried out at room temperature as well as at 600 °C. Sample dimension was chosen as 8 9 25 9 35 mm3 and the sample surface was ground with a 120 lm grinding disc. The plunger material used in these tests was EnDOtec DO*70 (48 HRC, Table 3). Characterisation of wear behaviour was performed by measuring the weight loss of the samples, by standard OM, stereo microscopy (SM) and SEM. The topography of the worn samples was evaluated with a confocal microscope. By comparing the area of a worn sample with an original sample, the volumetric loss of the worn area can be calculated. Through this method the volumetric wear can be measured directly, which is sometimes advisable. Due to oxidation effects, at higher temperatures, the samples or the sample substrate can gain weight, and so the wear cannot be measured in terms of mass loss directly when high testing temperatures were used. So, when there are no austenite (or other non-oxidising) substrates/materials available (or possible due to sample preparation), two solutions are possible. One possible solution is to parallel

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Fig. 3 a View of the SIT. b Testing principle. c Measuring of the impact mark length. d Measuring of the impact mark depth

run oxidation dummies and subtract the oxidation weight gain of the dummy from the tested sample, and calculate the volume loss via the density of the sample afterwards. But, in this case, the already mentioned approach via measuring the volume loss directly in the confocal microscope often is advisable. Also, cross-sectional images of the worn sample area were used to analyse the carbide breaking, cold work hardening, composite layer formation and changes in the matrix caused by the temperature influence.

3 Results and Discussion 3.1 SIT Results The impact marks on the investigated specimens were analysed by means of OM. Also by means of SEM and OM, analyses was carried out to get qualitative view of the impact process. It was found that, brittle materials containing a high percentage of hard phases like M7C3

carbides or tungsten carbides tend to form cracks and break outs as shown in Fig. 5. More ductile materials tend to deform plastically or cleft when the impact load is higher as shown in Fig. 6. The mean values of the impact mark length of the materials investigated can be found in Fig. 7. It can clearly be seen that the impact mark lengths of the materials containing hard phases tend to be shorter; on the other hand, break outs have occurred frequently. In material A, the austenitic steel, the impact mark length is decreased as the temperature increased from 300 to 600 °C. This could be due to the precipitation of additional Cr-carbides at higher temperatures which is usual for unstabilised austenitic steels [27]. As it can be seen from Fig. 6a and b, material A does not tend to crack or cleft at impact energies below 10 J. Material B, with a martensitic microstructure, shows normal linear softening at higher temperatures. At temperatures higher than 650 °C, the microstructure would change significantly since the martensitic matrix becomes spheriodise annealed, so it is expected that the behaviour

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Fig. 4 a View of the HT-CIAT. b Testing principle (45° impact angle with abrasives tested). c Testing chamber

Table 2 Testing parameters used in HT-CIAT Parameter

Value

Impact energy

0.8 J

Impact angle

45°

Frequency

2 Hz

Testing cycles

7,200

Abrasive material

Silica sand

Abrasive flow

3 g/s

Abrasive size, shape

0.4–0.9 mm, angular

Abrasive hardness

1000–1200 HV

Test temperatures

RT, 600 °C

Table 3 Chemical composition of the plunger material (wt%) C

Si

Mn

S

Cr

Ni

W

Fe

Co

1.7

1

0.5

0.01

29

0.1

8

2.5

Bal.

would diverge from linearity. The significant low standard deviation of material B is an indicator for a high homogeneity of the material. It can be seen from Fig. 6c and d that material B does not crack, but can generate small clefts at an impact energy of 7 J. Material C, complex Fe–Cr–C–Nb–Mo–W–B alloy, shows high resistance to single impacts at room temperature. Also with the increase in temperature, the impact mark length increases, and at temperatures above 300 °C, the matrix softens and thus the impact mark length increases nonlinearly beyond 300 °C. It can be seen from Fig. 5a and b

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that for material C, 7 J impact energy is already sufficient for crack formation. In material D, hypereutectic Fe–Cr–C, the impact mark length increases from RT to 300 °C, but at temperatures from 300 to 600 °C, the impact length remains constant. This could be due to the high percentage of large primary carbides which do not change their hardness and wear behaviour up to 600 °C. While the martensitic matrix softens with increasing temperature, the ductility of the material is increased and the hardness decreases, but the high hard phase content prevents the reduction in hardness (HV5) to lower than 770 HV, a value that was measured in a thermal ageing experiment for this material. It can be seen from Fig. 5c and d that material D tends to form break outs at 7 J, and the breakouts tend to be bigger than the cracks occurring in material C, which could be due to the coarser microstructure of material D. The correlation coefficients between the impact mark length and the hardness (HV5) of the materials are shown in Table 4. All the correlation coefficient values are negative which means that as hardness increases the impact mark length decreases, also the correlation coefficient varies just between -0.84 and -0.95, which indicates good correlation (in terms of linearity, i.e., if the correlation coefficient would be -1, the impact mark length would be linear decreasing with increasing hardness). Material C, which is the hardest material investigated (Table 1), shows the shortest impact mark length, and the softest material A shows the longest impact mark length.

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Fig. 5 OM of SIT wear marks. a, b Cracks developed in SIT for material C at 7 J impact at 600 °C. c, d Cracks developed in SIT for material D at 7 J impact at 600 °C

Fig. 6 OM of SIT wear marks. a, b Plastic stage for material A in SIT at 7 J impact at 600 °C. c, d Plastic stage for material B in SIT at 7 J impact at 600 °C

The depths of the impact marks generally correlate with the length, which is geometrically obvious; nevertheless, information on the plastic or elastic tendencies of the materials can be distinguished. With the length and the depth of the impact mark, the impact angle can be calculated, which in case of an ideally plastic impact should have 5° as the angle of the hitting edge. The smaller the

impact angle, the more elastic is the impact behaviour and more elastic is the material. The calculated impact angles can be found in Table 5. At room temperature, material A is the most elastic material investigated (Table 5). At 600 °C, the impact angle of material A is still approximately the same as at room temperature, indicating that the plastic behaviour

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Fig. 7 Variations of impact mark length with the different temperatures at 7 J impact energy

Table 4 Correlation coefficient of impact mark length vs. hardness at various temperatures calculated from the data form Fig. 7 Correlating parameters

Correlation coefficient Room temperature 300 °C 600 °C

Impact length rate vs. hardness -0.88

-0.95

-0.84

Table 5 Impact angle for SIT Materials

Impact angle (°) Room temperature

600 °C

A

4.16

4.16

B

4.41

3.71

C

4.81

4.07

D

4.45

3.39

does not change much. Material C, complex Fe–Cr–C–Nb– Mo–W–B alloy, shows plastic behaviour up to 300 °C before it breaks out; the impact angle is almost exactly 5°. From 300 °C to higher temperatures, the matrix behaviour as well as the impact angle changes, indicating a more elastic behaviour. Materials B and D get more elastic at 300 °C, but do not change their plastic behaviour up to 600 °C. It can be seen from materials A and C that materials which have a good matrix stability at higher temperatures have higher impact angles at 600 °C, and in materials B and C, where the matrix is spheroidise annealed, the impact angle is significantly lower. The hardness of the materials investigated seems to be the main influencing factor on the length of the impact

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marks caused by single impacts with high energy. In materials with high hard phase content, the cracking or breakout behaviour has a major influence on the wear performance in high impact wear, and the breakouts tend to be bigger with coarser microstructure. The critical impact energy where the material starts breaking out is temperature and microstructure dependent and not easily predictable. The critical impact energies where the materials start breaking out at room temperature can be found in Fig. 8. 3.2 HT-CIAT Results and Discussion The materials investigated show very different wear behaviour in the HT-CIAT, mainly influenced by the hard phase content. In Fig. 9, the worn surfaces of material A as a representative for materials with a low hard phase content and material C as representative for materials with a high hard phase content can be seen after testing at 600 °C. It was observed that in materials with high hard phase content (Fig. 9b) the grooving is low. The materials with less hard phases have deeper grooving (Fig. 9a) and form a composite layer (abrasive particles were mechanically mixed with the sample material). The composite layer can be seen in detail in Fig. 10 as an example for material B. By comparing the brittle materials C and D, it can be seen that the wear mechanism is significantly different, because in the brittle materials breakouts occur and cause the main wear, while in the ductile materials the main wear is due to ploughing and cutting. Furthermore, the ductile materials show plastic deformations, while the brittle ones do not.

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Fig. 8 Critical impact energies for the tested materials at room temperature

The different wear mechanisms can be better understood from SEM analyses, which can be seen for material A as an example for ductile materials in Fig. 11, and in Fig. 12 for material C as an example for hard phase-containing materials. Ductile materials generally become more plastic at high temperatures which causes more grooving and a thicker composite layer. In materials with high hard phase content, the matrix also becomes more plastic at high temperatures and can more easily be eroded. The protruding carbides also break at higher temperatures due to fatigue and by the worse matrix backing. The quantitative results of the materials tested in the HTCIAT at room temperature and at high temperatures can be found in Fig. 13. Material C, complex Fe–Cr–C–Nb–Mo– W–B alloy, shows the best wear resistance in the HT-CIAT. At room temperature, both the hard phase-containing materials C and D clearly perform better than the materials A and B, with low hard phase content, which is obvious, since a certain amount of hard phases is necessary to resist abrasive wear. The abrasion resistance at room temperature is quite correlated to the hardness (HV5) of the materials. Surprisingly, material D with the highest hard phase content does not perform better than material B at higher temperatures; the worse matrix backup at high temperatures leads to bigger breakouts, which increases the wear. Due to a finer microstructure and better matrix–carbide interfacial bonding, investigated earlier [13], material C does not suffer from this wear increase at higher temperatures that much. Both the ductile materials A and B form composite layers at high temperatures (Fig. 10) and have higher wear rates than material C, due to insufficient big hard phases, which are important for wear resistance in abrasive environments

especially at high temperatures. At high temperatures, when the matrix is softened, the carbides are important to stall the grooving, contrary to room temperature where the matrix is still sufficiently hard to resist the abrasive wear itself. As shown in Fig. 14, the tested alloys are correlated in terms of material parameters like macrohardness, matrix stability and hard phase coarseness with the stress fields abrasion, impacts load and high temperatures. Alloy C which has high hardness do have very good abrasion resistance and medium resistance to single impacts. Alloy D, in contrast, which also has high hardness, is quite good in abrasion resistance at room temperature, but loses much of its wear resistance at high temperatures, what could be caused by worse matrix–carbide interfacial bonding. Material D also performs disadvantageously in SIT; due to its coarse carbide structure, the break outs tend to be big. Material A performs well in the SIT but, at the same time, has the highest wear rate in the HT-CIAT, since without carbides nothing can resist the grooving. Material B has some small hard phases and higher hardness so performs a little better in the abrasion test at room temperature, but due to the martensitic matrix, once the application temperature would increase to 750 °C, the matrix is spheroidise annealed and the alloy loses its wear resistance. Material A with an austenitic matrix has the best matrix stability and so retains its positive properties at high temperatures. Material C’s matrix is also temperature resistant and also performs well at temperatures up to 650 °C. So, regarding the combined wear mechanisms of abrasion, impact loads and high temperature influence, material C is best suited to be used in the sinter grate, followed by material A which would be even better suited, if the abrasive wear would be less.

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Fig. 9 Micrographs of worn surfaces. a Material A after 600 °C testing. b Material C after 600 °C testing

Fig. 10 Cross-section of plastic stage with composite layer for material B after CIAT

Fig. 12 SEM image of material C after HT-CIAT at 600 °C

4 Conclusion

Fig. 11 SEM image of material A after HT-CIAT at 600 °C

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Aim of this work was to investigate the influence of the material parameters such as macrohardness and hard phase content, coarseness on the wear resistance in impact loading and abrasive applications at high temperatures. It was shown that the impact mark length of materials hit by single impacts with high energy is correlated to the hardness of the materials. The harder the material, the smaller will be the impact mark length. On the other hand, impact energies above a material-dependent critical energy can

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Fig. 13 Quantitative HT-CIAT wear rates of the tested materials at room temperature and at 600 °C

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film where the abrasive SiO2 particles are worked into the material during the testing. The wear rate in the materials with low hard phase content is higher compared to the complex Fe–Cr–C–Nb–Mo–W–B alloy, due to the lack of big hard phases which are necessary to stall grooving. The material parameter hardness, which is increased with higher hard phase content, improves the abrasion resistance, while the material parameter microstructure coarseness reduces the critical impact energy necessary to form breakouts. Good matrix behaviour is important for low wear rates at high temperature, it is expected that especially the bonding hard phase-matrix is important due to paradontose effect, meaning that broken or badly bound carbides simply fall out of the matrix. After comparing all the conditions it can be concluded that, from the materials tested here, the complex Fe–Cr–C– Nb–Mo–W–B alloy is best suited for the use as sinter grate for steel production due to its wear resistance at all conditions as well as its economic viability. Acknowledgements This work was funded by the ‘‘Austrian KplusProgram’’ and was carried out within the ‘‘Austrian Center of Competence for Tribology’’. The authors are also grateful to Bo¨hler-Edelstahl for providing starting materials and performing heat-treatment procedures in different steel types and to Castolin Eutectic for helpful work in manufacturing of welding samples. Markus Varga, BSc is acknowledged for helpful work in performing SIT tests and wear quantification.

References Fig. 14 Correlation diagram of material properties to material demands

create breakouts. These breakouts lead to increased wear in the real system. Below this material-dependent critical impact energy, the wear is expected to be significantly lower than above this threshold. The critical impact energy depends on the hardness of the hard phases, the size of the carbides, the hardness of the matrix, matrix–hard phase interfacial bonding and the total hard phase volume content. Therefore, it is very difficult to predict the critical energy through a model without any testing. Furthermore, especially the matrix parameters change with temperature. On the other hand, hard phases retain its hardness also at temperatures up to 800 °C and so it increases the materials abrasion resistance. Since the hard phase distribution and the hardness of the hard phases is hardly affected by temperatures up to 800 °C, the matrix-dependent parameters change with temperature. In case of the HT-CIAT test, the complex Fe–Cr–C–Nb–Mo–W–B alloy shows the best wear resistance compared to other materials. Here, ductile materials such as tool steel and austenite form a composite

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