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The effect of mechanical properties on erosion
307
Wear, 121 (1988)
307 - 324
THE EFFECT
OF MECHANICAL
PROPERTIES
ON EROSION*
ANDREW NINHAM? Lawrence
Berkeley
Laboratory,
University
of California, Berkeley,
CA 94720 (U.S.A.)
Summary A number of alloys, commonly used either for their abrasion resistance or their elevated temperature properties (high strength and corrosion resistance), were studied to investigate their erosion behavior. Angular quartz particles (70 - 200 pm diameter) and silicon carbide particles (250 pm average diameter) were used as erodents. Several of the alloys were tested in two conditions; annealed and in either a work-hardened or aged condition. The aging or work-hardening treatments carried out on the high temperature alloys had essentially no effect on erosion rates. Further, there were only small differences in erosion rate between the different high temperature alloys; they included cobalt-based, nickel-based, and iron-based alloys. The abrasion-resistant alloys generally eroded at higher rates than the high temperature alloys with the matrix composition being of greater importance than the morphology of carbides within the alloy. Erosion rates were not strongly dependent upon impingement angle for any of the alloys tested and in many cases the erosion rate was higher at 60” impingement than at either 30” or 90”. The carbide particles which confer abrasion resistance to alloys were found to disrupt plastic flow around the indentation, causing void formation and fracture. The carbides were thus deleterious to erosion performance. It is proposed that the high strain rate, short contact time, and essentially adiabatic nature of erodent particle impacts are responsible for the near-constant erosion rates of the high temperature alloys. Their mechanical properties (measured quasi-statically) covered a very wide range. Estimates of the local temperature rise at indentations are calculated for different alloys and are shown to support the proposition.
1. Introduction A number of investigations have been made into the effect that an alloy’s microstructure and mechanical properties have on its erosion resis*Paper presented at the International Conference on Wear of Materials, Houston, TX, U.S.A., April 5 - 9,1987. TPresent address: Department of Materials Science and Metallurgy, Cambridge University, Cambridge, U.K. 0043-1648/88/$3.50
@ Elsevier Sequoia/Printed in The Netherlands
308
tance, e.g. Gulden (1979) [ 11, Green et al. (1981) [2], and Foley and Levy (1983) [3]. Gulden (1979) [l] found that increasing the hardness of a given steel by quenching to form martensite increased its erosion rate which then fell as the steel was tempered to successively higher temperatures. The variation in erosion rate with heat treatment was far more marked at normal impact than at an impingement angle of 30” and for most of the steels tested moderate tempering (typically to maximum K,,) produced a rather flat erosion rate-impingement angle plot, often with the maximum in erosion rate near 60” impact angle. Green et al. (1981) [2] restricted their testing to 30” impact angle but found results consistent with those of Gulden. For a series of plain carbon steels (AISI 1020 to AISI 1080) an increase in erosion rate was measured as the carbon content of the steel increased. This was the case whether the steels were normalized or quenched, although the trend was stronger for steels in the martensitic condition. A rapid drop in erosion rate occurred for AISI 1080 steel tempered at 650 “C for times up to 30 min as compared with the untempered condition; longer tempering times had essentially no further effect on the erosion rate. The erosion rate of heavily tempered 1080 steel was close to that of normalized 1080 steel. Foley and Levy (1983) [3] found that heat treatment of AISI 4340 steel (0.4 wt.% C, 0.8 wt.% Cr, 1.8 wt.% Ni, 0.25 wt.% MO) had only a very small effect on its 30” impact angle erosion rate. Erosion tests on hot-rolled 1020 and cold-rolled 1020 steels, and annealed 304 (19Cr-1ONi) and wrought 304 stainless steels, indicated that ductility was an important parameter. There were exceptions, however; annealed AISI 4340 steel had twice the measured reduction in area of as-quenched 4340 steel during tensile tests with negligible differences in erosion rate between the two heat treatment conditions. In abrasion tests where metals are rubbed against abrasive papers, an increase in hardness results in a decrease in wear rate. A similarity with erosion exists, however, in that the effect is more pronounced for different annealed metals than for differences in hardness of a single alloy due to heat treatment (Khruschov (1957) [4]). Abrasion resistance may also be improved by employing alloys which contain very large hard particles, typically carbides, in a more ductile matrix: this is the basis of the commercially successful Stellite alloys. A systematic study by Aptekar and Kosel (1985) [5] of white cast irons with different volume fractions of carbide showed that the erosion behavior of these materials was sensitive to the properties of the erodent. Using alumina as the erodent (harder than the carbides in the iron), higher carbide volume fractions were detrimental to the erosion resistance, whereas quartz erodent (softer than the carbides) showed higher carbide volume fractions to be beneficial. However, even for quartz abrasive the white irons eroded faster than annealed 1020 steel, indicating the approach of using large second phase carbides to be ineffective in erosion despite its utility in abrasion.
309
The aim of this work is to examine the erosion behavior of a number of alloys normally used for their abrasive wear resistance or their high strength at elevated temperatures. The selection of the alloys, however, was based upon their having large differences in certain mechanical properties as the result of thermal or mechanical treatments. They were used as test vehicles to study the effects of large property changes in a single alloy on its erosion behavior. Thus several of the alloys were tested in two conditions; these were generally the annealed condition and the strengthened condition, by either mechanical or thermal treatment, as appropriate for the alloy. Other alloy pairs were chosen with only one property differing between them, so the effect of that single property on erosion behavior could be ascertained. 2. Experimental
procedure
The alloys tested are listed in Table 1 together with their nominal compositions and the mechanical properties that were of particular interest. Room temperature erosion tests were carried out on these alloys using angular 250 - 300 pm (60 mesh) silicon carbide abrasive and 75 - 200 pm quartz abrasive in an air-blast-type rig. The equipment has been described in some detail elsewhere (Lapides and Levy (1980) [6]). The particle velocity was 60 m s-r in each case, as determined by the computer model of Ninham and Hutchings (1983) [7], and tests were conducted at impingement angles of 30”, 60”, and 90”. TABLE Alloys
(1) (2)
(3)
(4)
(5) (6) (7)
(8)
1 tested
and nominal
compositions
Haynes 6B (wrought) Stellite number 6 (Oxyacetylene deposit). 6B and number 6 have similar compositions (Co-3OCr-4W-l.lC) but very different microstructures. 6B has large discrete carbides and an elongation of ll%, number 6 has a cast microstructure and 1% elongation Tristelle alloys TS-1, 2, 3. (Fe-35Cr-12Co20Ni-1, 2, or 3C). Series of iron-based wear-resistant alloys containing up to 3 wt.% C. Oxyacetylene-deposited TS-1 may be compared with Stellite number 6 - similar carbide volume, microstructure, and elongation Ferralium 255 - 30% cold reduced. High strength duplex stainless steel (Fe-26Cr-5Ni-3Mo) that has high strainhardening coefficient Haynes 188 -annealed and 30% cold reduced. (Co-22Ni-22Cr-14W). Cold reduction reduces elongation from 50% to 10% Hastelloy C-276 - annealed and 30% cold reduced. (Ni-16Cr-16Mo-5Fe-4W). Cold reduction reduces elongation from 65% to 15% Cabot 7 18 - annealed and aged. (Ni-19Ck-19Fe-3Mo-5(Nb + Ta)). Aging doubles yield strength and reduces elongation from 30% to 12% Berylco 25 - annealed and aged. (Cu-1.85Be). Aging doubles yield strength and reduces elongation from 38% to 4%
310
The testing was carried out incremen~ly, the specimen being cleaned by a high pressure air jet after each period of erosion and before weighing. Erosive mass losses were determined by weighing on a balance reading to 0.1 mg. 20 or 40 g increments of abrasive were used until the specimen had lost at least 15 mg. Steady state erosion rates were computed from least-squares linear regression of cumulative mass loss-cumulative mass of erodent plots. The mass loss after the first increment of erodent in all cases lay on the calculated straight line through the points, indicating that incubation effects were small; they will not be considered further here. A typical cumulative plot is shown in Fig. 1; the linear nature of steady state erosion following an initially low rate is evident. The steady state erosion rates were converted to volume rates, there being significant density differences between the various alloys. OnIy one erosion test was carried out on each of the be~llium copper specimens owing to the extreme toxicity of berylliumcontaining dust.
Fig. 1. Typical cumulative mass loss us. mass of abrasive plot from which erosion rates were calculated: 60 mesh silicon carbide; 30” impact angle; 60 m s’-‘; *, Haynes 188 cold reduced; m, Haynes 188 annealed.
After erosion testing, all of the specimens were examined by scanning electron microscopy. A number of specimens were sectioned through the eroded area and, in particular, a series of taper section specimens was prepared from samples eroded at 60” impact angle by quartz. The technique follows that of Mutton and Watson (1978) [S]; a cut is made through the erosion area using a silicon carbide slitting wheel, the cut lying at 45” to the specimen surface. The wide-angled piece of specimen is then metallographically mounted and prepared such that the cut surface is ground, polished, and etched. Careful removal of the mounting medium then allows simultaneous observation in the scanning electron microscope of the eroded surface and the microstructure immediately below the surface. The influence
311
of the microstructure examined.
on erosion
behavior
may
then
be more
readily
3. Results and discussion 3.1. Erosion rates Table 2 gives the steady state volume erosion rates of the alloys tested, for silicon carbide and for quartz erodents. Elongation and Vickers’ hardness values for the alloys are also included in this table. Figure 2 and Fig. 3 show erosion rate us. impingement angle plots for silicon carbide and quartz abrasives respectively. Only one data point is shown at each angle to represent both the annealed and the strengthened versions of each alloy. In the case of silicon carbide abrasive, the strengthening treatments made negligibly small differences to the erosion rates, see Table 2. For quartz abrasive the differences were somewhat larger, but still rather small, and mean values have been plotted. Figures 2 and 3 are also in two parts, for clarity. The lower plots show results for the alloys that contain carbides. These alloys are commonly used for abrasive wear resistance. Alloys shown in the upper plots are predominantly used for their corrosion resistance and high strength at elevated temperatures. It is clear from Table 2 and Figs. 2 and 3 that not only are erosion rates higher when silicon carbide abrasive is used (the impact velocities were TABLE 2 Steady state volume erosion rates of the alloys testeda Alloy
Haynes 6B (wrought) Stellite number 6 (deposit) Tristelle TS-1 (deposit) Tristelle TS-2 (wrought) Tristelle TS-3 (wrought) Ferralium 255 (cold rolled) Haynes 188 annealed Haynes 188 cold rolled Hastelloy C-276 annealed Hastelloy C-27 6 cold rolled Cabot 718 annealed Cabot 718 aged Berylco 25 annealed Berylco 25 aged
Elongation @)
11 1 1 22 53 9 67 15 30 12 38 4
Hardness WV)
494 465 330 430 460 407 320 465 229 408 243 476 199 410
Erosion rate (cm3 g-r
x
10e5)
Silicon carbide
Quartz
30’
60”
90”
30’
60”
90”
2.29 2.32 2.74 2.31 2.07 2.18 2.10 2.13 2.15 2.28 2.81 3.07
2.26 2.36 2.55 1.81 1.89 1.89 1.78 1.77 1.75 1.93 -
1.94 1.96 2.10 1.45 1.54 1.52 1.46 1.52 1.54 1.59 -
1.14 1.08 1.50 1.46 1.60 1.38 1.07 1.18 1.20 1.17 1.22 1.32 -
1.28 1.48 1.60 2.11 2.45 1.31 1.41 1.38 1.22 1.32 1.29 1.40 -
1.29 1.34 1.47 1.87 1.91 1.08 1.16 1.16 1.07 1.11 1.10 1.17 -
aRoom temperature erosion; air carrier gas; 60 m s-r impingement; carbide; 135 pm quartz.
250
pm silicon
312 14
3Or
1.6-
29
14" 20Fm
12-
"ilO? 0 0.9 X0 ^,
- l-Sa "E E ':
:i&+ 4'
L 30
CI
60
90
15
10j
0’
30 60 Impingement AngleIDegreesl
90
OLd
30 60 Impingement AngleCOegreesi
90
Fig. 2. Volume erosion rates vs. impingement angle for tests employing silicon carbide erodent: 250 pm silicon carbide; 60 m s-l; 0, Ferralium 256; 0, Cabot 718; 0, Haynes 188; 0, Hastelloy C-276; 0, Tristelle T&l; a, Stellite number 6, +, Haynes 6B. Fig. 3. Volume erosion rates us. impingement angfe for tests employing quartz as erodent: 75 - 200 pm quartz; 60 m s-l ;o, Ferralium 255; 0, Cabot 718; 0, Haynes 188; n, Hastelloy C-276; 0, Tristelle TS-1; 0, Tristelle TS-2; 0, Tristelle TS-3; I, Stellite 6; +, Haynes 6B.
similar) but that the response with respect to impingement angle is closer to that expected for ductile materials than when quartz abrasive is used. For silicon carbide abrasive, the erosion rate at normal impact was always less than it was at 30” impact angle, although the difference in rates between these angles was small and for Haynes 6B and Stellite number 6 a maximum existed near 60” impingement angle. It is noteworthy that at 30” impingement angle the alloys designed for high temperature service eroded at similar rates to the wear-resistant alloys and at steeper angles they eroded less than the wear-resistant alloys. The results of tests employing quartz as erodent show a trend of maximum erosion occurring at a higher angle than for silicon carbide. With the exception of Ferralium 255 which showed a monotonic decrease in erosion rate with angle (for the limited number of angles tested) and Haynes 6B which showed a monotonic increase, all of the alloys eroded more at 60” than at either 30’ or 90’ impact angle. As Aptekar and Kosel
313
(1985) [5] point out, this should not be interpreted as being the result of the sum of a “ductile” erosion curve due to the matrix and a “brittle” curve due to carbides; for most of the alloys tested here that contain carbides, the carbides are smaller than the damage zone caused by individual particle impacts. Also, several of the single-phase alloys show similar peaks in erosion near 60”. For these high strength materials, particularly when eroded by quartz, the erosion mechanism itself is intermediate between ideally ductiIe and ideally brittle mechanisms. The volume loss data in Table 2 show that the erosion rates for these alloys, which cover a wide range of mechanical properties, are all quite similar and that the variation in erosion rate with angle is small. There are no great differences between iron-based, cobalt-based, or nickel-based alloys. With the exception of the higher carbon content Tristelles, the maximum erosion rates seen here are no more than roughly 40% greater than the rates at normal impact; classically ductile materials typically exhibit a factor of three between maximum and 90” impact rates (e.g. Rickerby and Macmillan (1981) [ 91) while classically brittle materials display a monotonically increasing erosion rate with angle. The relationships between inden~tion hardness and erosion rate, and between ductility and erosion rate, are shown in Figs. 4 and 5. The erosion rates used for these figures are those resulting from impact by quartz at 60” impact angle. Where the hardness or ductility has been changed by thermal or mechanical treatment, points have been plotted to show the erosion rate for the material m both the annealed and the treated conditions and those points joined by a solid line. The lines are for clarity of presentation only; they are not meant to imply an attempt to interpolate between pairs of data points. However, the slopes of the lines do indicate the trends in erosion rate that changes in elongation or hardness have on particular alloys, given that the aging or cold-working treatments change a number of mechanical prop-
I
lo
zoo
250
300 BY Vickers Hardness
600 (kgf mm+
450
)
SO0
Fig. 4. Plot of erosion rate us. Vickers’ Hardness number. The erosion rates used were those resulting from erosion by 75 - 200 pm quartz impacting at 60 m s-l and 60’ impact angle.
11 t 10'
0
10
20
30 40 Elongation(%I
50
60
10
Fig. 5. Plot of erosion rate us. percentage elongation in a tensile test for the alloys tested. The erosion rates used were those resulting from erosion by 75 - 200 pm quartz impacting at 60 m s-l and 60” impact angle.
erties; the hardness and elongation are not independently altered. By considering the percentage change in erosion rate and the percentage change in hardness or elongation, the slopes of the lines may be calculated, the result also being a percentage, i.e. the percentage change in erosion rate for 100% increase in indentation hardness or elongation. The data in Figs. 4 and 5 give slopes of order lo%, with Haynes 188 showing a trend in the opposite direction to that shown by Hastelloy C-276 and Cabot 718. Thus only small changes in erosion rate occur for large changes in elongation and hardness. Elongation and hardness have a roughly inverse relationship during thermomechanical treatments; increasing one almost invariably reduces the other. A consideration of the plasticity and material displacement which occurs during a particle impact leads to the intuitive expectation that both high strength and high ductility are desirable for erosion resistance. This has also been predicted theoretically by Hutchings (1981) [lo]. Thus to see erosion performance changed so little by large changes in hardness and ductility is most likely a manifestation of their inverse relationship. The results of Figs. 4 and 5 also indicate that neither elongation nor indentation hardness dominate erosion behavior at the expense of the other. Naim and Bahadur (1985) [ll] studied the erosion behavior of an 18 wt.% Ni maraging steel whose hardness could be varied (over a limited range) with negligible change in ductility. They found erosion increased as the hardness increased while the ductility was constant. This is in conflict with the work of Finnie et al. (1967) [12], who showed that erosion decreased with increasing hardness, albeit for pure annealed metals with different ductilities. Naim and Bahadur (1985) [ll] also noted an increase in erosion rate as the ductility was reduced by cold working, at essentially constant strength. It should be noted that ductility as measured in a quasistatic tensile test may be very different from the ductility that pertains during the very high strain rate of a particle impact.
3.2. Eroded surfaces Examination of eroded surfaces in the scanning electron microscope confirms the indications of the mass loss data: for a given abrasive the erosion behavior is very similar for a wide variety of alloys which have large variations in their mechanical properties. No differences were visible between the eroded surfaces of annealed alloys and those that had been aged or cold worked, even for erosion by quartz. It was hoped that quartz, being a milder abrasive than silicon carbide, would better differentiate between the alloys. While this was true for the mass loss measurements, the differences were too small to be reflected in the morphology of the eroded surfaces. A typical example is shown in Fig. 6 which compares cold-worked and annealed Hastelloy C-276.
Fig. 6. Scanning electron micrographs of (a) annealed and (b) cold-reduced C-276 eroded by 75 - 200 (urnquartz at 60 m 6-l.
Hastelloy
Figure 7 compares the surfaces of Stellite number 6, Haynes 6B, and Tristelle TS-1 eroded by silicon carbide at 30’ impingement angle. Stellite number 6 (cobalt based) and Tristelle TS-1 (iron based) were both oxyacetylene deposited, resulting in the formation of interdendritic carbides during solidification, while Haynes 6B is a wrought alloy having a composition close to that of number 6 but with much more rounded carbides dispersed uniformly throughout the matrix; its tensile elongation is ten times that of the deposited alloys. Although the eroded surfaces show many similar features, it is evident that TS-1 has more deformed material attached to its surface, the cobalt-based alloys having a “cleaner” surface. Also, the difference in erosion rates between 6B and number 6 is very small, much smaller than the difference between TS-1 and number 6. These factors indicate that the morphology of the carbides has a much smaller effect on erosion performance than does the composition of the matrix containing them; the intrinsic hardness and ductility of the matrix is more important than the hardness and ductility of the whole alloy. The matrix of -Stellite number 6 is a Co-23Cr-4.5W solid solution (Silence (1977) [13]). At normal impact by quartz, a sample of pure
316
Fig. I. Scanning electron micrographs of the surfaces of the carbide-containing alloys eroded by 250 pm silicon carbide at 60 m s-l and 60’: (a) Stellite 6; (b) Haynes 6B; (c) Tristelle TS-1.
Co-25Cr eroded only 2% more than Stellite number 6 tested under identical conditions. Although the difference in erosion rates was greater at 30” impingement angle, this result supports the assertion that carbides do not greatly improve impact erosion performance and that the matrix properties dominate erosion behavior. Results of erosion tests on model alloys and on engineering alloys (Aptekar and Kosel(lQ85) [ 51; Kosel and Aptekar (1986) [14]) that contain hard second phase particles are in agreement with this. The effect that carbides have on plastic flow at an impact site, and below the eroded surface in general, is particularly well illustrated by the taper-sectioned specimens. Two examples are given in Fig. 8 which shows Haynes 6B and Stellite number 6 eroded by quartz at 60” impact angle. Slip bands are evident in grains close to the eroded surface but the main points to note are the void formation and the disruption of plastic flow caused by carbides. For the larger more rounded carbides (Haynes 6B), voiding at the carbide-matrix interface and carbide cracking are common; it should be noted that these carbides have not been struck by abrasive particles but have cracked because of stresses imposed by the matrix. The interdendritic carbides (Stellite number 6), in addition to disrupting plastic flow, cause displaced material to be only weakly attached to the surface owing to the discontinuous nature of the more ductile matrix in a region of
Fig. 8. Scanning electron micrographs showing tapered sections for (a) Haynes 6B and (b) Stellite 6. Erosion conditions of 75 - 200 pm quartz impacting at 60 m s-l and 60’.
carbides. In either case, the carbides do little to resist erosion. It is instructive here to consider the size of the erodent particles: even quartz particles toward the bottom of the size range used are some ten times larger than the carbides present in Haynes 6B (see Fig. 8). 4. Erosion mechanism The eroded surfaces of all the alloys tested in this study have many features in common; this is the case even when considering together surfaces impacted by the two different erodents. Micrographs of the erodents are presented in Fig. 9. Both silicon carbide and quartz are quite angular and thus the similarity in appearance between surfaces eroded by the two types of particles is to be expected. The features seen in Figs. 6 and 7 are also common to many eroded surfaces presented in the literature, e.g. by Ives and Ruff (1979) f15] who studied copper, by Naim and Bahadur (1985) [II] during erosion studies of a maraging steel and by Levy (1986) [ 161 who eroded aluminum and steel alloys.
Fig. 9. SiIicon carbide (250 - 300 pm) and (b) quartz (75 - 200 pm) erodents.
318
The general mech~ism is thus seen to be one of gouging (type I cutting, as defined by Hutchings (1979) 1171) in which material is raised above the mean surface, followed by deformation during subsequent impacts and ultimately removal of the deformed volume. There will be considerable variation in the number of impacts required to remove any given volume of material, according to the orientation and shape of the particle and its contact point relative to the displaced material. Final removal may involve cutting, strain to failure or some combination of these two processes. Gulden (1979) [l] noted that deformation around single-impact craters was highly plastic and localized, even in steels heat treated so as to exhibit an erosion with angle signature characteristic of “brittle” materials. Thus similar processes occur during the erosion of widely differing metallic materials under a variety of conditions. Attempts to improve erosion resistance by alloying and/or heat treatment have generally proven unsuccessful even when based on approaches that have been fruitful in reducing other types of wear, e.g. Green et aZ. (1981) [2]. It would appear that there are unidentified factors which cause erosion rates to differ from those which might be expected from consideration of abrasive wear test results and that render heat treatments ineffective. A trade-off that has been cited is that between hardness and ductility (e.g. Finnie et al. (1967) [12]); based on quasi-static data, a hard surface should suffer a smaller indentation than a softer surface, for a given impact, leading to a lower erosion rate. Since erosion rates are not usually reduced by hardening t~atmen~, and often increase, the reduction in ductility attendant with most hardening treatments is invoked to explain the constant erosion rate, since under most circumstances displaced material on the eroding surface is deformed to the failure strain. However, different materials lose ductility at very different rates for given strength increases, as has been seen in this study, and yet erosion rates still show little variation. Much of this inconsistency is based upon the use of quasi-static data for both hardness and ductility values, although this is frequently the only data available. Results from some simple dynamic hardness tests carried out on the alloys used in this study illustrate the problem. Steel ball-bearings, 1.588 mm in diameter, were fired at the alloys, using the air blast erosion rig mentioned earlier but with a nozzle, 2 mm in inner diameter. The exit velocity was estimated to be roughly 100 m s-I. While the exact velocity is not known, the reproducibi~ty of results on a single specimen indicates that the velocity itself was reproducible. The dynamic hardness of the alloys cannot be calculated without a precise knowledge of the particle velocity (and even then there are problems in measuring the correct crater volume, Tabor (1951) 1181) but the ratios of dynamic hardnesses may be estimated for pairs of specimens of the same composition, one of which has undergone a hardening treatment. It is informative to compare these dynamic ratios with quasi-static hardness ratios. Some results are presented in Table 3.
319 TABLE 3 Hardness ratios
Dynamic Quasi-static Vickers’
Cabot 718 aged-annealed
BeCu aged-annealed
Hastelloy C-276 cold reduced-annealed
Tristelle TS-3-TS-1
1.353 1.838 1.959
1.440 2.007 2.060
1.202 1.742 1.782
1.044 1.0 1.4
aVery inhomogeneous microstructure
a
for TS-3.
It is evident that quasi-static hardness tests overestimate the increase in hardness introduced into these specimens by the hardening treatment, compared with dynamic hardness values. Also, since the strain rate during a particle impact varies inversely with particle size and as the square root of particle velocity (Hutchings (1977) [ 19]), it is likely that the strain rate experienced during the contact between an asperity on a particle and the target during an erosion test is some two orders of magnitude greater than the 5 X lo4 s-l which occurred during the ball-bearing impacts. In a review of dynamic plasticity, Campbell (1972) [20] reports that for f.c.c. metals the increase in flow stress which occurs on increasing the strain rate becomes smaller as the quasi-static flow stress is increased by alloying or heat treatment. Thus, at the very high strain rates typical of erosive impacts, high strength materials have a greatly reduced strength advantage over low strength materials. The simple experiments described above are in agreement with this. It is not clear, however, how the ductility of different strength materials will vary at high strain rates. Also associated with the high strain rates is a temperature increase within the volume of material being deformed. Such temperature increases reduce the flow stress, although there is evidence that at strain rates approaching 10’ s-i the flow stress is dependent upon strain rate but not temperature (Campbell (1972) [20]). Nevertheless, in many situations a reduction in flow stress will occur. The reductions in flow stress are likely to be due to thermal activation of dislocations only, since the duration of temperature transients is very small and there is insufficient time at elevated temperature for such metallurgical changes as overaging of precipitationhardened alloys or recrystallization to occur; the latter is particularly due to the concurrent severe plastic flow. Only metallurgical changes that do not require diffusion can take place during the temperature transient around an impact. Thus, if high enough temperatures are reached, ferrite might transform to austenite in iron-based systems, subsequently re-transforming to martensite on cooling. These hypotheses were investigated by experiments on AISI 1096 steel in which a fully martensitic sample and a spheroidized sample were eroded at normal impact using 250 pm silicon carbide at 60 m s-i. The spheroidized
320
sample eroded at one-third the rate of the martensitic sample, Shallow angle taper sections through the eroded areas allowed microhardness profiles to be obtained which included reliable data from indentations made very close to the eroded surface. The microstructure below the eroded surface was also visible on an expanded scale. The spheroidized sample, containing carbides of roughly 1 pm diameter, had undergone very large strains around the impact craters with many examples of “folding over” of raised lips of material. The untempered specimen exhibited much more localized deformation and a light etching layer at the eroded surface. Microhardness profiles for the two specimens are shown in Fig. 10. It should be noted that the distance scale along the ordinate corresponds to distance along the taper surface; this is an expansion of approximately five times compared with the depth below the eroded surface. The increase in hardness near to the surface of the spheroidized specimen can only be ,att~buted to work hardening while that for the ~-quench~ specimen could be the result of reaustenitizing and quenching in combination with plastic deformation or plastic deformation alone. The altered microstructure in the vicinity of the eroded surface (as-quenched specimen) is compared with the microstructure of the bulk in Fig. 11; the difference in mi~rost~ct~e implies that changes more drastic than work hardening alone have taken place. Extremely rapid quenching of austenite is required to form martensite that does not etch heavily, as is seen here, since this is the result of the carbon rem~ning in solid solution; quenching of minute volumes of austenite into the bulk metal will produce such a quench rate. Conventional water quenching following removal from a furnace allows a degree of autotempering to occur, resulting in rapid etching. Identical microstructural effects were seen by Gulden (1979) [l] when examining material displaced by single impacts on AISI 1095 steel tempered
VtiN
1100
900
100 200 Ofstance From Waded
300
LOO
Edge Ipml
Fig. 10. Microhardness profiles from the eroded surface of quenched, and quenched and tempered 1096 steel.
into the bulk on taper
sections
Fig. 11. Microstructure of quenched 1096 steel: (a) from the bulk;(b) immediately under the eroded surface. Erosion conditions: 250 /.fm silicon carbide impacting normally at 60 m s-l.
to maximum hardness, i.e. fine-grained martensite in the bulk but a featureless surface on the crater lips following etching. An estimate of the local temperature rise during a particle impact may be made if the event is assumed to be adiabatic. This assumption is justified when the time for thermal diffusion over a characteristic distance is large compared with the time of impact. Table 4 shows estimated values of the thermal diffusion time, the impact time, and the adiabatic temperature rise for impacts on to a variety of metals by quartz, 100 I.rrnin diameter, travelling at 60 m s-I. Since these values are strongly dependent upon the hardness of the metal, the values have been calculated for two different indentation hardnesses, which correspond approximately to values for the annealed metal and a harder alloy based on the particular metal. In calculating values for this table it has been assumed that 50% of the particle kinetic energy is converted into heat, the crater volume is inversely proportional to the quasi-static indentation pressure (hardness), and that the characteristic diffusion distance for the heat generated is equal to the depth of the crater formed. The contact times were calculated from an equation given by Tabor (1951) [18]. While the calculations have involved numerous simplifying assumptions, an estimate of 50% for the fraction of energy converted to heat is conservative and thus the calculated temperature increases should not be unduly oversized. It is evident that for the higher hardness values of each metal considerable temperature rises may occur, particularly for the metals with low thermal diffusivities such as iron and titanium. Conversely, for high conductivity metals, the impact cannot be considered adiabatic and the temperature increases will be smaller than indicated in Table 4. At lower strain rates such as lo3 - lo4 s-l this temperature rise in the immediate vicinity of the impact will markedly reduce the yield strength. The highest strength alloys undergo the greatest temperature rise and therefore the largest fractional reduction in strength. At higher strain rates the flow stress depends only
322 TABLE
Metal
Al cu Fe Ti co Ni NiCr
4
Hardness
R.m.s. diffusion time (ps)
Contact time (P)
AT (adiabatic)
@@‘a)
300 1200 500 4000 1250 7000 1200 5000 1000 4600 900 4000 1200 4700
0.13 0.05 0.08 0.02 0.26 0.08 0.59 0.23 0.26 0.09 0.31 0.11 1.2 0.49
0.183 0.092 0.142 0.050 0.090 0.038 0.092 0.045 0.100 0.047 0.106 0.050 0.092 0.046
55 220 70 560 147 825 230 960 122 560 100 450 140 550
(K)
tdiffusion
tcontact 0.71 0.54 0.56 0.40 2.9 2.1 6.4 5.1 2.6 1.9 2.9 2.2 13 11
upon the strain rate (Campbell (1972) [20]), corresponding to viscous flow, and in this condition prior metallurgical treatments will have a negligible effect. The combined effects of localized adiabatic (or nearly adiabatic) heating and sensitivity, predominantly, to the erosion conditions (strain rate) therefore renders control of erosion by alloy manipulation implausible; the more an alloy is strengthened, the larger is the temperature rise in the deformed material, defeating the strength increase. The very similar erosion rates seen for the alloys tested in this work is evidence of these effects. The implication of this is that in order to resist erosion it is insufficient to reduce the amount of plasticity on the surface; plasticity must be prevented entirely in the target. Impacts must therefore be either entirely elastic or plasticity limited to the erodent, i.e. the target must be significantly harder than the erodent, since impulsive pressures are generated during impacts which may be considerably greater than the yield pressure of either erodent or target.
5. Conclusions (1) The erosion rates (for a given abrasive) of several different alloys designed for elevated temperature service are very similar and are essentially unaffected by prior cold work or precipitation-hardening treatments. (2) Abrasion-resistant alloys containing second phase carbides also show similar erosion rates despite differences in matrix composition and carbide morphology. Further, these abrasion-resistant alloys erode at a somewhat higher rate than alloys designed for elevated temperature use.
323
(3) Results from a series of wear-resistant alloys of similar composition (Tristelles) indicate that an increasing carbide volume fraction causes an increase in erosion rate. (4) The erosion response of high strength materials, particularly when eroded by quartz, is only weakly dependent upon impact angle. Most materials give rise to a rather flat erosion us. angle plot, with a peak at roughly 60” impact. This peak does become more marked, however, for carbidecontaining alloys, particularly as the carbide content is increased. (5) Erosion rates are higher for silicon carbide abrasive than for angular quartz abrasive in tests conducted at the same impact velocity. It is not clear from the results of this work if this is due to differences in hardness and shape between the two erodents or to a size effect. The quartz used was somewhat smaller than the silicon carbide abrasive. (6) Experimental evidence and theoretical considerations indicate that very high temperature excursions of very small duration occur on eroding surfaces; temperature rises are greater for high strength alloys than for low strength alloys, The high temperature of a deforming volume, in combination with extremely high strain rates, leads to viscous flow which is unaffected by prior treatment of the alloy. The flow is also only weakly sensitive to the alloy’s physical properties (composition). (7) The carbides which confer abrasion resistance to alloys cause a disruption of plastic flow during erosion. The inhomogeneous nature of the plastic flow results in very high strain gradients and, consequently, to void formation near to and cracking of the carbides. These processes facilitate material removal. (8) In view of the above conclusions, it is unlikely that metallic alloys can be modified to reduce significantly their erosion rates, except in situations involving very soft abrasives.
Acknowledgment This work was sponsored by the U.S. Department of Energy under DOE/FE AA 15 10 10 0, Advanced Research and Technical Development, Fossil Energy Materials Program, work breakdown structure elements LBL-3 and under contract number DE-AC03-76SF00098.
References 1 M. E. Golden, in J. E. Field (ed.), Proc. 5th Znt. Conf. on Erosion by Liquid and Solid Impact, Cambridge, September 1979, Cavendish Laboratory, Cambridge, 1979, Paper 31. 2 G. M. Green, R. Taggart and D. H. Polonis, Metallography, 14 (1981) 191 - 212. 3 T. Foley and A. Levy, Wear, 91 (1983) 45 - 64. 4 M. M. Khruschov, Proc. Conf. on Lubrication and Wear, London, 1957, Institution of Mechanical Engineers, London, 1957, pp. 657 - 659.
324 5 S. S. Aptekar and T. H. Kosel, in K. C. Ludema (ed.), Proc. Int. Conf. on Wear of MaterioZs, Vancouver, April 14 - 18, Z985, American Society of Mechanical Engineers, New York, 1985, pp. 677 - 686. 6 L. Lapides and A. Levy, Wear, 58 (1980) 301 - 312. 7 A. J. Ninham and I. M. Hutchings, in J. E. Field (ed.), Proc. 6th Znt. Conf on Erosion by Liquid and Solid impact, Cambridge, September 1983, Cavendish Laboratory, Cambridge, 1979, Paper 50. 8 P. J. Mutton and J. D. Watson, Wear, 48 (1978) 385 - 389. 9 D. G. Rickerby and N. H. Macmi~an, in S. K. Rhee, A. W. Ruff and K. C. Ludema (eds.), Proc. Znt. Conf. on Wear of Materials, San Francisco, CA, March 30 -April 1, 1981, American Society of Mechanical Engineers, New York, 1981, pp. 548 - 563. 10 1. M. Hutchings, Wear, 70 (1981) 269 - 282. 11 M. Naim and S. Bahadur, in K. C. Ludema (ed.), Proc. fnt. Conf on Wear of Materials, Vancouver, April 14 - 18, 1985, American Society of Mechanical Engineers, New York, 1985, pp. 586 - 594. 12 I. Finnie, J. Wolak and Y. Kabil, J. Mater., 2 (1967) 682 - 700. 13 W. L. Silence, in W. A. Glaeser, K. C. Ludema and S. K. Rhee (eds.), Proc. Znt. Conf. on Wear of Mater~~s, St. Louis, MO, April 25 - 28, 1977, American Society of Mechanical Engineers, New York, 197 7. 14 T. H. Kosel and S. S. Aptekar, Corrosion ‘86, National Association of Corrosion Engineers, Houston, TX, 1986, Paper 113. 15 L. K. Ives and A. W. Ruff, in W. F. Adler (ed.), Erosion: Prevention and Useful Applications, ASTM Spec. Tech. Publ, 664, 19’79, pp. 5 - 35. 16 A. Levy, Wear, 108 (1) (1986) 1 - 22. 17 I. M. Hutchings, in W. F. Adler (ed.), Erosion: Prevention and Useful Applications, ASTM Spec. Tech. Publ. 664, 1979, pp. 59 - 75. 18 D. Tabor, The Gardner of Metals, Clarendon, Oxford, 1951. 19 I. M. Hutchings, J. Phys. r), 10 (1977) L179 - L184. 20 J. D. Campbell, Mater. Sci. Eng., 12 (1973) 3 - 21.
Wear, 121 (1988)
307 - 324
THE EFFECT
OF MECHANICAL
PROPERTIES
ON EROSION*
ANDREW NINHAM? Lawrence
Berkeley
Laboratory,
University
of California, Berkeley,
CA 94720 (U.S.A.)
Summary A number of alloys, commonly used either for their abrasion resistance or their elevated temperature properties (high strength and corrosion resistance), were studied to investigate their erosion behavior. Angular quartz particles (70 - 200 pm diameter) and silicon carbide particles (250 pm average diameter) were used as erodents. Several of the alloys were tested in two conditions; annealed and in either a work-hardened or aged condition. The aging or work-hardening treatments carried out on the high temperature alloys had essentially no effect on erosion rates. Further, there were only small differences in erosion rate between the different high temperature alloys; they included cobalt-based, nickel-based, and iron-based alloys. The abrasion-resistant alloys generally eroded at higher rates than the high temperature alloys with the matrix composition being of greater importance than the morphology of carbides within the alloy. Erosion rates were not strongly dependent upon impingement angle for any of the alloys tested and in many cases the erosion rate was higher at 60” impingement than at either 30” or 90”. The carbide particles which confer abrasion resistance to alloys were found to disrupt plastic flow around the indentation, causing void formation and fracture. The carbides were thus deleterious to erosion performance. It is proposed that the high strain rate, short contact time, and essentially adiabatic nature of erodent particle impacts are responsible for the near-constant erosion rates of the high temperature alloys. Their mechanical properties (measured quasi-statically) covered a very wide range. Estimates of the local temperature rise at indentations are calculated for different alloys and are shown to support the proposition.
1. Introduction A number of investigations have been made into the effect that an alloy’s microstructure and mechanical properties have on its erosion resis*Paper presented at the International Conference on Wear of Materials, Houston, TX, U.S.A., April 5 - 9,1987. TPresent address: Department of Materials Science and Metallurgy, Cambridge University, Cambridge, U.K. 0043-1648/88/$3.50
@ Elsevier Sequoia/Printed in The Netherlands
308
tance, e.g. Gulden (1979) [ 11, Green et al. (1981) [2], and Foley and Levy (1983) [3]. Gulden (1979) [l] found that increasing the hardness of a given steel by quenching to form martensite increased its erosion rate which then fell as the steel was tempered to successively higher temperatures. The variation in erosion rate with heat treatment was far more marked at normal impact than at an impingement angle of 30” and for most of the steels tested moderate tempering (typically to maximum K,,) produced a rather flat erosion rate-impingement angle plot, often with the maximum in erosion rate near 60” impact angle. Green et al. (1981) [2] restricted their testing to 30” impact angle but found results consistent with those of Gulden. For a series of plain carbon steels (AISI 1020 to AISI 1080) an increase in erosion rate was measured as the carbon content of the steel increased. This was the case whether the steels were normalized or quenched, although the trend was stronger for steels in the martensitic condition. A rapid drop in erosion rate occurred for AISI 1080 steel tempered at 650 “C for times up to 30 min as compared with the untempered condition; longer tempering times had essentially no further effect on the erosion rate. The erosion rate of heavily tempered 1080 steel was close to that of normalized 1080 steel. Foley and Levy (1983) [3] found that heat treatment of AISI 4340 steel (0.4 wt.% C, 0.8 wt.% Cr, 1.8 wt.% Ni, 0.25 wt.% MO) had only a very small effect on its 30” impact angle erosion rate. Erosion tests on hot-rolled 1020 and cold-rolled 1020 steels, and annealed 304 (19Cr-1ONi) and wrought 304 stainless steels, indicated that ductility was an important parameter. There were exceptions, however; annealed AISI 4340 steel had twice the measured reduction in area of as-quenched 4340 steel during tensile tests with negligible differences in erosion rate between the two heat treatment conditions. In abrasion tests where metals are rubbed against abrasive papers, an increase in hardness results in a decrease in wear rate. A similarity with erosion exists, however, in that the effect is more pronounced for different annealed metals than for differences in hardness of a single alloy due to heat treatment (Khruschov (1957) [4]). Abrasion resistance may also be improved by employing alloys which contain very large hard particles, typically carbides, in a more ductile matrix: this is the basis of the commercially successful Stellite alloys. A systematic study by Aptekar and Kosel (1985) [5] of white cast irons with different volume fractions of carbide showed that the erosion behavior of these materials was sensitive to the properties of the erodent. Using alumina as the erodent (harder than the carbides in the iron), higher carbide volume fractions were detrimental to the erosion resistance, whereas quartz erodent (softer than the carbides) showed higher carbide volume fractions to be beneficial. However, even for quartz abrasive the white irons eroded faster than annealed 1020 steel, indicating the approach of using large second phase carbides to be ineffective in erosion despite its utility in abrasion.
309
The aim of this work is to examine the erosion behavior of a number of alloys normally used for their abrasive wear resistance or their high strength at elevated temperatures. The selection of the alloys, however, was based upon their having large differences in certain mechanical properties as the result of thermal or mechanical treatments. They were used as test vehicles to study the effects of large property changes in a single alloy on its erosion behavior. Thus several of the alloys were tested in two conditions; these were generally the annealed condition and the strengthened condition, by either mechanical or thermal treatment, as appropriate for the alloy. Other alloy pairs were chosen with only one property differing between them, so the effect of that single property on erosion behavior could be ascertained. 2. Experimental
procedure
The alloys tested are listed in Table 1 together with their nominal compositions and the mechanical properties that were of particular interest. Room temperature erosion tests were carried out on these alloys using angular 250 - 300 pm (60 mesh) silicon carbide abrasive and 75 - 200 pm quartz abrasive in an air-blast-type rig. The equipment has been described in some detail elsewhere (Lapides and Levy (1980) [6]). The particle velocity was 60 m s-r in each case, as determined by the computer model of Ninham and Hutchings (1983) [7], and tests were conducted at impingement angles of 30”, 60”, and 90”. TABLE Alloys
(1) (2)
(3)
(4)
(5) (6) (7)
(8)
1 tested
and nominal
compositions
Haynes 6B (wrought) Stellite number 6 (Oxyacetylene deposit). 6B and number 6 have similar compositions (Co-3OCr-4W-l.lC) but very different microstructures. 6B has large discrete carbides and an elongation of ll%, number 6 has a cast microstructure and 1% elongation Tristelle alloys TS-1, 2, 3. (Fe-35Cr-12Co20Ni-1, 2, or 3C). Series of iron-based wear-resistant alloys containing up to 3 wt.% C. Oxyacetylene-deposited TS-1 may be compared with Stellite number 6 - similar carbide volume, microstructure, and elongation Ferralium 255 - 30% cold reduced. High strength duplex stainless steel (Fe-26Cr-5Ni-3Mo) that has high strainhardening coefficient Haynes 188 -annealed and 30% cold reduced. (Co-22Ni-22Cr-14W). Cold reduction reduces elongation from 50% to 10% Hastelloy C-276 - annealed and 30% cold reduced. (Ni-16Cr-16Mo-5Fe-4W). Cold reduction reduces elongation from 65% to 15% Cabot 7 18 - annealed and aged. (Ni-19Ck-19Fe-3Mo-5(Nb + Ta)). Aging doubles yield strength and reduces elongation from 30% to 12% Berylco 25 - annealed and aged. (Cu-1.85Be). Aging doubles yield strength and reduces elongation from 38% to 4%
310
The testing was carried out incremen~ly, the specimen being cleaned by a high pressure air jet after each period of erosion and before weighing. Erosive mass losses were determined by weighing on a balance reading to 0.1 mg. 20 or 40 g increments of abrasive were used until the specimen had lost at least 15 mg. Steady state erosion rates were computed from least-squares linear regression of cumulative mass loss-cumulative mass of erodent plots. The mass loss after the first increment of erodent in all cases lay on the calculated straight line through the points, indicating that incubation effects were small; they will not be considered further here. A typical cumulative plot is shown in Fig. 1; the linear nature of steady state erosion following an initially low rate is evident. The steady state erosion rates were converted to volume rates, there being significant density differences between the various alloys. OnIy one erosion test was carried out on each of the be~llium copper specimens owing to the extreme toxicity of berylliumcontaining dust.
Fig. 1. Typical cumulative mass loss us. mass of abrasive plot from which erosion rates were calculated: 60 mesh silicon carbide; 30” impact angle; 60 m s’-‘; *, Haynes 188 cold reduced; m, Haynes 188 annealed.
After erosion testing, all of the specimens were examined by scanning electron microscopy. A number of specimens were sectioned through the eroded area and, in particular, a series of taper section specimens was prepared from samples eroded at 60” impact angle by quartz. The technique follows that of Mutton and Watson (1978) [S]; a cut is made through the erosion area using a silicon carbide slitting wheel, the cut lying at 45” to the specimen surface. The wide-angled piece of specimen is then metallographically mounted and prepared such that the cut surface is ground, polished, and etched. Careful removal of the mounting medium then allows simultaneous observation in the scanning electron microscope of the eroded surface and the microstructure immediately below the surface. The influence
311
of the microstructure examined.
on erosion
behavior
may
then
be more
readily
3. Results and discussion 3.1. Erosion rates Table 2 gives the steady state volume erosion rates of the alloys tested, for silicon carbide and for quartz erodents. Elongation and Vickers’ hardness values for the alloys are also included in this table. Figure 2 and Fig. 3 show erosion rate us. impingement angle plots for silicon carbide and quartz abrasives respectively. Only one data point is shown at each angle to represent both the annealed and the strengthened versions of each alloy. In the case of silicon carbide abrasive, the strengthening treatments made negligibly small differences to the erosion rates, see Table 2. For quartz abrasive the differences were somewhat larger, but still rather small, and mean values have been plotted. Figures 2 and 3 are also in two parts, for clarity. The lower plots show results for the alloys that contain carbides. These alloys are commonly used for abrasive wear resistance. Alloys shown in the upper plots are predominantly used for their corrosion resistance and high strength at elevated temperatures. It is clear from Table 2 and Figs. 2 and 3 that not only are erosion rates higher when silicon carbide abrasive is used (the impact velocities were TABLE 2 Steady state volume erosion rates of the alloys testeda Alloy
Haynes 6B (wrought) Stellite number 6 (deposit) Tristelle TS-1 (deposit) Tristelle TS-2 (wrought) Tristelle TS-3 (wrought) Ferralium 255 (cold rolled) Haynes 188 annealed Haynes 188 cold rolled Hastelloy C-276 annealed Hastelloy C-27 6 cold rolled Cabot 718 annealed Cabot 718 aged Berylco 25 annealed Berylco 25 aged
Elongation @)
11 1 1 22 53 9 67 15 30 12 38 4
Hardness WV)
494 465 330 430 460 407 320 465 229 408 243 476 199 410
Erosion rate (cm3 g-r
x
10e5)
Silicon carbide
Quartz
30’
60”
90”
30’
60”
90”
2.29 2.32 2.74 2.31 2.07 2.18 2.10 2.13 2.15 2.28 2.81 3.07
2.26 2.36 2.55 1.81 1.89 1.89 1.78 1.77 1.75 1.93 -
1.94 1.96 2.10 1.45 1.54 1.52 1.46 1.52 1.54 1.59 -
1.14 1.08 1.50 1.46 1.60 1.38 1.07 1.18 1.20 1.17 1.22 1.32 -
1.28 1.48 1.60 2.11 2.45 1.31 1.41 1.38 1.22 1.32 1.29 1.40 -
1.29 1.34 1.47 1.87 1.91 1.08 1.16 1.16 1.07 1.11 1.10 1.17 -
aRoom temperature erosion; air carrier gas; 60 m s-r impingement; carbide; 135 pm quartz.
250
pm silicon
312 14
3Or
1.6-
29
14" 20Fm
12-
"ilO? 0 0.9 X0 ^,
- l-Sa "E E ':
:i&+ 4'
L 30
CI
60
90
15
10j
0’
30 60 Impingement AngleIDegreesl
90
OLd
30 60 Impingement AngleCOegreesi
90
Fig. 2. Volume erosion rates vs. impingement angle for tests employing silicon carbide erodent: 250 pm silicon carbide; 60 m s-l; 0, Ferralium 256; 0, Cabot 718; 0, Haynes 188; 0, Hastelloy C-276; 0, Tristelle T&l; a, Stellite number 6, +, Haynes 6B. Fig. 3. Volume erosion rates us. impingement angfe for tests employing quartz as erodent: 75 - 200 pm quartz; 60 m s-l ;o, Ferralium 255; 0, Cabot 718; 0, Haynes 188; n, Hastelloy C-276; 0, Tristelle TS-1; 0, Tristelle TS-2; 0, Tristelle TS-3; I, Stellite 6; +, Haynes 6B.
similar) but that the response with respect to impingement angle is closer to that expected for ductile materials than when quartz abrasive is used. For silicon carbide abrasive, the erosion rate at normal impact was always less than it was at 30” impact angle, although the difference in rates between these angles was small and for Haynes 6B and Stellite number 6 a maximum existed near 60” impingement angle. It is noteworthy that at 30” impingement angle the alloys designed for high temperature service eroded at similar rates to the wear-resistant alloys and at steeper angles they eroded less than the wear-resistant alloys. The results of tests employing quartz as erodent show a trend of maximum erosion occurring at a higher angle than for silicon carbide. With the exception of Ferralium 255 which showed a monotonic decrease in erosion rate with angle (for the limited number of angles tested) and Haynes 6B which showed a monotonic increase, all of the alloys eroded more at 60” than at either 30’ or 90’ impact angle. As Aptekar and Kosel
313
(1985) [5] point out, this should not be interpreted as being the result of the sum of a “ductile” erosion curve due to the matrix and a “brittle” curve due to carbides; for most of the alloys tested here that contain carbides, the carbides are smaller than the damage zone caused by individual particle impacts. Also, several of the single-phase alloys show similar peaks in erosion near 60”. For these high strength materials, particularly when eroded by quartz, the erosion mechanism itself is intermediate between ideally ductiIe and ideally brittle mechanisms. The volume loss data in Table 2 show that the erosion rates for these alloys, which cover a wide range of mechanical properties, are all quite similar and that the variation in erosion rate with angle is small. There are no great differences between iron-based, cobalt-based, or nickel-based alloys. With the exception of the higher carbon content Tristelles, the maximum erosion rates seen here are no more than roughly 40% greater than the rates at normal impact; classically ductile materials typically exhibit a factor of three between maximum and 90” impact rates (e.g. Rickerby and Macmillan (1981) [ 91) while classically brittle materials display a monotonically increasing erosion rate with angle. The relationships between inden~tion hardness and erosion rate, and between ductility and erosion rate, are shown in Figs. 4 and 5. The erosion rates used for these figures are those resulting from impact by quartz at 60” impact angle. Where the hardness or ductility has been changed by thermal or mechanical treatment, points have been plotted to show the erosion rate for the material m both the annealed and the treated conditions and those points joined by a solid line. The lines are for clarity of presentation only; they are not meant to imply an attempt to interpolate between pairs of data points. However, the slopes of the lines do indicate the trends in erosion rate that changes in elongation or hardness have on particular alloys, given that the aging or cold-working treatments change a number of mechanical prop-
I
lo
zoo
250
300 BY Vickers Hardness
600 (kgf mm+
450
)
SO0
Fig. 4. Plot of erosion rate us. Vickers’ Hardness number. The erosion rates used were those resulting from erosion by 75 - 200 pm quartz impacting at 60 m s-l and 60’ impact angle.
11 t 10'
0
10
20
30 40 Elongation(%I
50
60
10
Fig. 5. Plot of erosion rate us. percentage elongation in a tensile test for the alloys tested. The erosion rates used were those resulting from erosion by 75 - 200 pm quartz impacting at 60 m s-l and 60” impact angle.
erties; the hardness and elongation are not independently altered. By considering the percentage change in erosion rate and the percentage change in hardness or elongation, the slopes of the lines may be calculated, the result also being a percentage, i.e. the percentage change in erosion rate for 100% increase in indentation hardness or elongation. The data in Figs. 4 and 5 give slopes of order lo%, with Haynes 188 showing a trend in the opposite direction to that shown by Hastelloy C-276 and Cabot 718. Thus only small changes in erosion rate occur for large changes in elongation and hardness. Elongation and hardness have a roughly inverse relationship during thermomechanical treatments; increasing one almost invariably reduces the other. A consideration of the plasticity and material displacement which occurs during a particle impact leads to the intuitive expectation that both high strength and high ductility are desirable for erosion resistance. This has also been predicted theoretically by Hutchings (1981) [lo]. Thus to see erosion performance changed so little by large changes in hardness and ductility is most likely a manifestation of their inverse relationship. The results of Figs. 4 and 5 also indicate that neither elongation nor indentation hardness dominate erosion behavior at the expense of the other. Naim and Bahadur (1985) [ll] studied the erosion behavior of an 18 wt.% Ni maraging steel whose hardness could be varied (over a limited range) with negligible change in ductility. They found erosion increased as the hardness increased while the ductility was constant. This is in conflict with the work of Finnie et al. (1967) [12], who showed that erosion decreased with increasing hardness, albeit for pure annealed metals with different ductilities. Naim and Bahadur (1985) [ll] also noted an increase in erosion rate as the ductility was reduced by cold working, at essentially constant strength. It should be noted that ductility as measured in a quasistatic tensile test may be very different from the ductility that pertains during the very high strain rate of a particle impact.
3.2. Eroded surfaces Examination of eroded surfaces in the scanning electron microscope confirms the indications of the mass loss data: for a given abrasive the erosion behavior is very similar for a wide variety of alloys which have large variations in their mechanical properties. No differences were visible between the eroded surfaces of annealed alloys and those that had been aged or cold worked, even for erosion by quartz. It was hoped that quartz, being a milder abrasive than silicon carbide, would better differentiate between the alloys. While this was true for the mass loss measurements, the differences were too small to be reflected in the morphology of the eroded surfaces. A typical example is shown in Fig. 6 which compares cold-worked and annealed Hastelloy C-276.
Fig. 6. Scanning electron micrographs of (a) annealed and (b) cold-reduced C-276 eroded by 75 - 200 (urnquartz at 60 m 6-l.
Hastelloy
Figure 7 compares the surfaces of Stellite number 6, Haynes 6B, and Tristelle TS-1 eroded by silicon carbide at 30’ impingement angle. Stellite number 6 (cobalt based) and Tristelle TS-1 (iron based) were both oxyacetylene deposited, resulting in the formation of interdendritic carbides during solidification, while Haynes 6B is a wrought alloy having a composition close to that of number 6 but with much more rounded carbides dispersed uniformly throughout the matrix; its tensile elongation is ten times that of the deposited alloys. Although the eroded surfaces show many similar features, it is evident that TS-1 has more deformed material attached to its surface, the cobalt-based alloys having a “cleaner” surface. Also, the difference in erosion rates between 6B and number 6 is very small, much smaller than the difference between TS-1 and number 6. These factors indicate that the morphology of the carbides has a much smaller effect on erosion performance than does the composition of the matrix containing them; the intrinsic hardness and ductility of the matrix is more important than the hardness and ductility of the whole alloy. The matrix of -Stellite number 6 is a Co-23Cr-4.5W solid solution (Silence (1977) [13]). At normal impact by quartz, a sample of pure
316
Fig. I. Scanning electron micrographs of the surfaces of the carbide-containing alloys eroded by 250 pm silicon carbide at 60 m s-l and 60’: (a) Stellite 6; (b) Haynes 6B; (c) Tristelle TS-1.
Co-25Cr eroded only 2% more than Stellite number 6 tested under identical conditions. Although the difference in erosion rates was greater at 30” impingement angle, this result supports the assertion that carbides do not greatly improve impact erosion performance and that the matrix properties dominate erosion behavior. Results of erosion tests on model alloys and on engineering alloys (Aptekar and Kosel(lQ85) [ 51; Kosel and Aptekar (1986) [14]) that contain hard second phase particles are in agreement with this. The effect that carbides have on plastic flow at an impact site, and below the eroded surface in general, is particularly well illustrated by the taper-sectioned specimens. Two examples are given in Fig. 8 which shows Haynes 6B and Stellite number 6 eroded by quartz at 60” impact angle. Slip bands are evident in grains close to the eroded surface but the main points to note are the void formation and the disruption of plastic flow caused by carbides. For the larger more rounded carbides (Haynes 6B), voiding at the carbide-matrix interface and carbide cracking are common; it should be noted that these carbides have not been struck by abrasive particles but have cracked because of stresses imposed by the matrix. The interdendritic carbides (Stellite number 6), in addition to disrupting plastic flow, cause displaced material to be only weakly attached to the surface owing to the discontinuous nature of the more ductile matrix in a region of
Fig. 8. Scanning electron micrographs showing tapered sections for (a) Haynes 6B and (b) Stellite 6. Erosion conditions of 75 - 200 pm quartz impacting at 60 m s-l and 60’.
carbides. In either case, the carbides do little to resist erosion. It is instructive here to consider the size of the erodent particles: even quartz particles toward the bottom of the size range used are some ten times larger than the carbides present in Haynes 6B (see Fig. 8). 4. Erosion mechanism The eroded surfaces of all the alloys tested in this study have many features in common; this is the case even when considering together surfaces impacted by the two different erodents. Micrographs of the erodents are presented in Fig. 9. Both silicon carbide and quartz are quite angular and thus the similarity in appearance between surfaces eroded by the two types of particles is to be expected. The features seen in Figs. 6 and 7 are also common to many eroded surfaces presented in the literature, e.g. by Ives and Ruff (1979) f15] who studied copper, by Naim and Bahadur (1985) [II] during erosion studies of a maraging steel and by Levy (1986) [ 161 who eroded aluminum and steel alloys.
Fig. 9. SiIicon carbide (250 - 300 pm) and (b) quartz (75 - 200 pm) erodents.
318
The general mech~ism is thus seen to be one of gouging (type I cutting, as defined by Hutchings (1979) 1171) in which material is raised above the mean surface, followed by deformation during subsequent impacts and ultimately removal of the deformed volume. There will be considerable variation in the number of impacts required to remove any given volume of material, according to the orientation and shape of the particle and its contact point relative to the displaced material. Final removal may involve cutting, strain to failure or some combination of these two processes. Gulden (1979) [l] noted that deformation around single-impact craters was highly plastic and localized, even in steels heat treated so as to exhibit an erosion with angle signature characteristic of “brittle” materials. Thus similar processes occur during the erosion of widely differing metallic materials under a variety of conditions. Attempts to improve erosion resistance by alloying and/or heat treatment have generally proven unsuccessful even when based on approaches that have been fruitful in reducing other types of wear, e.g. Green et aZ. (1981) [2]. It would appear that there are unidentified factors which cause erosion rates to differ from those which might be expected from consideration of abrasive wear test results and that render heat treatments ineffective. A trade-off that has been cited is that between hardness and ductility (e.g. Finnie et al. (1967) [12]); based on quasi-static data, a hard surface should suffer a smaller indentation than a softer surface, for a given impact, leading to a lower erosion rate. Since erosion rates are not usually reduced by hardening t~atmen~, and often increase, the reduction in ductility attendant with most hardening treatments is invoked to explain the constant erosion rate, since under most circumstances displaced material on the eroding surface is deformed to the failure strain. However, different materials lose ductility at very different rates for given strength increases, as has been seen in this study, and yet erosion rates still show little variation. Much of this inconsistency is based upon the use of quasi-static data for both hardness and ductility values, although this is frequently the only data available. Results from some simple dynamic hardness tests carried out on the alloys used in this study illustrate the problem. Steel ball-bearings, 1.588 mm in diameter, were fired at the alloys, using the air blast erosion rig mentioned earlier but with a nozzle, 2 mm in inner diameter. The exit velocity was estimated to be roughly 100 m s-I. While the exact velocity is not known, the reproducibi~ty of results on a single specimen indicates that the velocity itself was reproducible. The dynamic hardness of the alloys cannot be calculated without a precise knowledge of the particle velocity (and even then there are problems in measuring the correct crater volume, Tabor (1951) 1181) but the ratios of dynamic hardnesses may be estimated for pairs of specimens of the same composition, one of which has undergone a hardening treatment. It is informative to compare these dynamic ratios with quasi-static hardness ratios. Some results are presented in Table 3.
319 TABLE 3 Hardness ratios
Dynamic Quasi-static Vickers’
Cabot 718 aged-annealed
BeCu aged-annealed
Hastelloy C-276 cold reduced-annealed
Tristelle TS-3-TS-1
1.353 1.838 1.959
1.440 2.007 2.060
1.202 1.742 1.782
1.044 1.0 1.4
aVery inhomogeneous microstructure
a
for TS-3.
It is evident that quasi-static hardness tests overestimate the increase in hardness introduced into these specimens by the hardening treatment, compared with dynamic hardness values. Also, since the strain rate during a particle impact varies inversely with particle size and as the square root of particle velocity (Hutchings (1977) [ 19]), it is likely that the strain rate experienced during the contact between an asperity on a particle and the target during an erosion test is some two orders of magnitude greater than the 5 X lo4 s-l which occurred during the ball-bearing impacts. In a review of dynamic plasticity, Campbell (1972) [20] reports that for f.c.c. metals the increase in flow stress which occurs on increasing the strain rate becomes smaller as the quasi-static flow stress is increased by alloying or heat treatment. Thus, at the very high strain rates typical of erosive impacts, high strength materials have a greatly reduced strength advantage over low strength materials. The simple experiments described above are in agreement with this. It is not clear, however, how the ductility of different strength materials will vary at high strain rates. Also associated with the high strain rates is a temperature increase within the volume of material being deformed. Such temperature increases reduce the flow stress, although there is evidence that at strain rates approaching 10’ s-i the flow stress is dependent upon strain rate but not temperature (Campbell (1972) [20]). Nevertheless, in many situations a reduction in flow stress will occur. The reductions in flow stress are likely to be due to thermal activation of dislocations only, since the duration of temperature transients is very small and there is insufficient time at elevated temperature for such metallurgical changes as overaging of precipitationhardened alloys or recrystallization to occur; the latter is particularly due to the concurrent severe plastic flow. Only metallurgical changes that do not require diffusion can take place during the temperature transient around an impact. Thus, if high enough temperatures are reached, ferrite might transform to austenite in iron-based systems, subsequently re-transforming to martensite on cooling. These hypotheses were investigated by experiments on AISI 1096 steel in which a fully martensitic sample and a spheroidized sample were eroded at normal impact using 250 pm silicon carbide at 60 m s-i. The spheroidized
320
sample eroded at one-third the rate of the martensitic sample, Shallow angle taper sections through the eroded areas allowed microhardness profiles to be obtained which included reliable data from indentations made very close to the eroded surface. The microstructure below the eroded surface was also visible on an expanded scale. The spheroidized sample, containing carbides of roughly 1 pm diameter, had undergone very large strains around the impact craters with many examples of “folding over” of raised lips of material. The untempered specimen exhibited much more localized deformation and a light etching layer at the eroded surface. Microhardness profiles for the two specimens are shown in Fig. 10. It should be noted that the distance scale along the ordinate corresponds to distance along the taper surface; this is an expansion of approximately five times compared with the depth below the eroded surface. The increase in hardness near to the surface of the spheroidized specimen can only be ,att~buted to work hardening while that for the ~-quench~ specimen could be the result of reaustenitizing and quenching in combination with plastic deformation or plastic deformation alone. The altered microstructure in the vicinity of the eroded surface (as-quenched specimen) is compared with the microstructure of the bulk in Fig. 11; the difference in mi~rost~ct~e implies that changes more drastic than work hardening alone have taken place. Extremely rapid quenching of austenite is required to form martensite that does not etch heavily, as is seen here, since this is the result of the carbon rem~ning in solid solution; quenching of minute volumes of austenite into the bulk metal will produce such a quench rate. Conventional water quenching following removal from a furnace allows a degree of autotempering to occur, resulting in rapid etching. Identical microstructural effects were seen by Gulden (1979) [l] when examining material displaced by single impacts on AISI 1095 steel tempered
VtiN
1100
900
100 200 Ofstance From Waded
300
LOO
Edge Ipml
Fig. 10. Microhardness profiles from the eroded surface of quenched, and quenched and tempered 1096 steel.
into the bulk on taper
sections
Fig. 11. Microstructure of quenched 1096 steel: (a) from the bulk;(b) immediately under the eroded surface. Erosion conditions: 250 /.fm silicon carbide impacting normally at 60 m s-l.
to maximum hardness, i.e. fine-grained martensite in the bulk but a featureless surface on the crater lips following etching. An estimate of the local temperature rise during a particle impact may be made if the event is assumed to be adiabatic. This assumption is justified when the time for thermal diffusion over a characteristic distance is large compared with the time of impact. Table 4 shows estimated values of the thermal diffusion time, the impact time, and the adiabatic temperature rise for impacts on to a variety of metals by quartz, 100 I.rrnin diameter, travelling at 60 m s-I. Since these values are strongly dependent upon the hardness of the metal, the values have been calculated for two different indentation hardnesses, which correspond approximately to values for the annealed metal and a harder alloy based on the particular metal. In calculating values for this table it has been assumed that 50% of the particle kinetic energy is converted into heat, the crater volume is inversely proportional to the quasi-static indentation pressure (hardness), and that the characteristic diffusion distance for the heat generated is equal to the depth of the crater formed. The contact times were calculated from an equation given by Tabor (1951) [18]. While the calculations have involved numerous simplifying assumptions, an estimate of 50% for the fraction of energy converted to heat is conservative and thus the calculated temperature increases should not be unduly oversized. It is evident that for the higher hardness values of each metal considerable temperature rises may occur, particularly for the metals with low thermal diffusivities such as iron and titanium. Conversely, for high conductivity metals, the impact cannot be considered adiabatic and the temperature increases will be smaller than indicated in Table 4. At lower strain rates such as lo3 - lo4 s-l this temperature rise in the immediate vicinity of the impact will markedly reduce the yield strength. The highest strength alloys undergo the greatest temperature rise and therefore the largest fractional reduction in strength. At higher strain rates the flow stress depends only
322 TABLE
Metal
Al cu Fe Ti co Ni NiCr
4
Hardness
R.m.s. diffusion time (ps)
Contact time (P)
AT (adiabatic)
@@‘a)
300 1200 500 4000 1250 7000 1200 5000 1000 4600 900 4000 1200 4700
0.13 0.05 0.08 0.02 0.26 0.08 0.59 0.23 0.26 0.09 0.31 0.11 1.2 0.49
0.183 0.092 0.142 0.050 0.090 0.038 0.092 0.045 0.100 0.047 0.106 0.050 0.092 0.046
55 220 70 560 147 825 230 960 122 560 100 450 140 550
(K)
tdiffusion
tcontact 0.71 0.54 0.56 0.40 2.9 2.1 6.4 5.1 2.6 1.9 2.9 2.2 13 11
upon the strain rate (Campbell (1972) [20]), corresponding to viscous flow, and in this condition prior metallurgical treatments will have a negligible effect. The combined effects of localized adiabatic (or nearly adiabatic) heating and sensitivity, predominantly, to the erosion conditions (strain rate) therefore renders control of erosion by alloy manipulation implausible; the more an alloy is strengthened, the larger is the temperature rise in the deformed material, defeating the strength increase. The very similar erosion rates seen for the alloys tested in this work is evidence of these effects. The implication of this is that in order to resist erosion it is insufficient to reduce the amount of plasticity on the surface; plasticity must be prevented entirely in the target. Impacts must therefore be either entirely elastic or plasticity limited to the erodent, i.e. the target must be significantly harder than the erodent, since impulsive pressures are generated during impacts which may be considerably greater than the yield pressure of either erodent or target.
5. Conclusions (1) The erosion rates (for a given abrasive) of several different alloys designed for elevated temperature service are very similar and are essentially unaffected by prior cold work or precipitation-hardening treatments. (2) Abrasion-resistant alloys containing second phase carbides also show similar erosion rates despite differences in matrix composition and carbide morphology. Further, these abrasion-resistant alloys erode at a somewhat higher rate than alloys designed for elevated temperature use.
323
(3) Results from a series of wear-resistant alloys of similar composition (Tristelles) indicate that an increasing carbide volume fraction causes an increase in erosion rate. (4) The erosion response of high strength materials, particularly when eroded by quartz, is only weakly dependent upon impact angle. Most materials give rise to a rather flat erosion us. angle plot, with a peak at roughly 60” impact. This peak does become more marked, however, for carbidecontaining alloys, particularly as the carbide content is increased. (5) Erosion rates are higher for silicon carbide abrasive than for angular quartz abrasive in tests conducted at the same impact velocity. It is not clear from the results of this work if this is due to differences in hardness and shape between the two erodents or to a size effect. The quartz used was somewhat smaller than the silicon carbide abrasive. (6) Experimental evidence and theoretical considerations indicate that very high temperature excursions of very small duration occur on eroding surfaces; temperature rises are greater for high strength alloys than for low strength alloys, The high temperature of a deforming volume, in combination with extremely high strain rates, leads to viscous flow which is unaffected by prior treatment of the alloy. The flow is also only weakly sensitive to the alloy’s physical properties (composition). (7) The carbides which confer abrasion resistance to alloys cause a disruption of plastic flow during erosion. The inhomogeneous nature of the plastic flow results in very high strain gradients and, consequently, to void formation near to and cracking of the carbides. These processes facilitate material removal. (8) In view of the above conclusions, it is unlikely that metallic alloys can be modified to reduce significantly their erosion rates, except in situations involving very soft abrasives.
Acknowledgment This work was sponsored by the U.S. Department of Energy under DOE/FE AA 15 10 10 0, Advanced Research and Technical Development, Fossil Energy Materials Program, work breakdown structure elements LBL-3 and under contract number DE-AC03-76SF00098.
References 1 M. E. Golden, in J. E. Field (ed.), Proc. 5th Znt. Conf. on Erosion by Liquid and Solid Impact, Cambridge, September 1979, Cavendish Laboratory, Cambridge, 1979, Paper 31. 2 G. M. Green, R. Taggart and D. H. Polonis, Metallography, 14 (1981) 191 - 212. 3 T. Foley and A. Levy, Wear, 91 (1983) 45 - 64. 4 M. M. Khruschov, Proc. Conf. on Lubrication and Wear, London, 1957, Institution of Mechanical Engineers, London, 1957, pp. 657 - 659.
324 5 S. S. Aptekar and T. H. Kosel, in K. C. Ludema (ed.), Proc. Int. Conf. on Wear of MaterioZs, Vancouver, April 14 - 18, Z985, American Society of Mechanical Engineers, New York, 1985, pp. 677 - 686. 6 L. Lapides and A. Levy, Wear, 58 (1980) 301 - 312. 7 A. J. Ninham and I. M. Hutchings, in J. E. Field (ed.), Proc. 6th Znt. Conf on Erosion by Liquid and Solid impact, Cambridge, September 1983, Cavendish Laboratory, Cambridge, 1979, Paper 50. 8 P. J. Mutton and J. D. Watson, Wear, 48 (1978) 385 - 389. 9 D. G. Rickerby and N. H. Macmi~an, in S. K. Rhee, A. W. Ruff and K. C. Ludema (eds.), Proc. Znt. Conf. on Wear of Materials, San Francisco, CA, March 30 -April 1, 1981, American Society of Mechanical Engineers, New York, 1981, pp. 548 - 563. 10 1. M. Hutchings, Wear, 70 (1981) 269 - 282. 11 M. Naim and S. Bahadur, in K. C. Ludema (ed.), Proc. fnt. Conf on Wear of Materials, Vancouver, April 14 - 18, 1985, American Society of Mechanical Engineers, New York, 1985, pp. 586 - 594. 12 I. Finnie, J. Wolak and Y. Kabil, J. Mater., 2 (1967) 682 - 700. 13 W. L. Silence, in W. A. Glaeser, K. C. Ludema and S. K. Rhee (eds.), Proc. Znt. Conf. on Wear of Mater~~s, St. Louis, MO, April 25 - 28, 1977, American Society of Mechanical Engineers, New York, 197 7. 14 T. H. Kosel and S. S. Aptekar, Corrosion ‘86, National Association of Corrosion Engineers, Houston, TX, 1986, Paper 113. 15 L. K. Ives and A. W. Ruff, in W. F. Adler (ed.), Erosion: Prevention and Useful Applications, ASTM Spec. Tech. Publ, 664, 19’79, pp. 5 - 35. 16 A. Levy, Wear, 108 (1) (1986) 1 - 22. 17 I. M. Hutchings, in W. F. Adler (ed.), Erosion: Prevention and Useful Applications, ASTM Spec. Tech. Publ. 664, 1979, pp. 59 - 75. 18 D. Tabor, The Gardner of Metals, Clarendon, Oxford, 1951. 19 I. M. Hutchings, J. Phys. r), 10 (1977) L179 - L184. 20 J. D. Campbell, Mater. Sci. Eng., 12 (1973) 3 - 21.