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Liquid phase sintering of metal matrix composites containing solid lubricants
Journal of
ELSEVIER
Journal of Materials Processing Technology 56 (1996) 136-147
Materials Processing Technology
LIQUID PHASE SINTERING OF METAL MATRIX COMPOSITES CONTAINING SOLID LUBRICANTS
J.D. Bolton and A.J. Gant Department of Mechanical Eng., University of Bradford, U.K.
ABSTRACT Attempts to improve the wear properties of sintered high speed steels have been made by the simultaneous addition of hard ceramic particles (TIC) or (NbC) alongside particles with solid lubricative properties such as (MnS) or (CaF2). Preliminary investigations carried out to study any interactions between such particles and a sintered high speed steel matrix indicated that the steel matrix carbides were modified by a chemical reaction which occurred between the added ceramic particle and its high speed steel matrix; the result of which was to give a substantial improvement in hardness. The solid lubricant remained chemically unaltered by the sintering process, and tended to reduce hardness. Both types of particulate addition raised the sintering temperature needed to achieve full density due to their effect on solidus/liquidus temperatures. Studies of the sintering kinetics in such materials confirmed that densification was due to a liquid phase sintering mechanism and that solution/re-precipitation of carbide phases played a major role in the densification process. Coarsening of the solid lubricant particles into large agglomerates occurred during sintering by a process of the transportation of smaller particles into pore interstices by the liquid phase.
1. INTRODUCTION
Preliminary investigations were carried out to study any interactions between particulate additions of TiC/NbC hard ceramic carbide, MnS/CaF 2 solid lubricants to a sintered high speed steel matrix in the hope of forming a metal matrix composite with enhanced wear properties. Emphasis was placed on determining any chemical reactions between the added particle and its high speed steel matrix plus any subsequent effects on the microstructure and sintering kinetics of the high speed steel. The chemical stability and distribution of the particulate addition in relation to the steel matrix were also studied.
2. EXPERIMENTAL WORK
Sintered composites were all basically prepared by mixing a water atomised M3 class 2 high speed steel (45~tm grain size) powder of the composition shown in Table 1, with one or more of the following powder ceramic carbides or solid lubricants 0924-0136/96/$15.00 © 1996 Elsevier Science S.A. All rights reserved SSD10924-0136 ( 95 ) 01829-4
.I.D. Bolton, A.J. Gant / Journal ¢~fMaterials Processing Technology 56 (1996) 136-147
Table 1. Composition of M3/2 high speed steel powder, %C %Cr %W %Mo %V 1.10 4.02 6.05 5.80 2.91 Manganese sulphide (MnS); low oxygen grade. Manganese sulphide (MnS); average particle size 10gm. Calcium fluoride (CaF2); average particle size 30gin. Titanium carbide (TIC); average particle size 51am. Titanium carbide (TIC); average particle size 151am. Niobium carbide (NbC); average grain size 41am. Graphite. Composite mixtures were made, to which 5 wt% additions of either solid lubricant or hard ceramic particles were blended by dry mixing in a Y cone blender for 15 rains: 0.25 wt% of graphite was also added to all batches to assist the sintering process, [3]. Compacts 4ram thick by 15 mm square were die pressed to green densities of approximately 70% theoretical, at pressure of 800 M.Pa and were sintered by holding for a fixed period under a vacuum (0.15 labar) after heating at 10OC/min to a choice of different sintering times of between 2 mins and four hours at a range of temperatures between 1180 and 1340°C. The as-sintered densities of all specimens were measured by weighing and by determining the volume by water displacement, but specimens were first coated with lacquer to prevent water penetration into surface pores during the weighing process. An appropriate correction was made for the weight and volume of this lacquer coating. Characterisation of microstruetures and the identification of phases present was performed by both optical and manning electron microscopy, assisted by the use of X ray energy dispersive analysis (E.D.A), back scattered electron image contrast (BE.I), and by some X-ray diffraction data. Quantitative measurements of microstructural features such as particle size and volume fraction for both the solid lubricant, ceramic particles, and of the high speed steel carbide phases were carried out from back scattered electron images using the "Magicscan, J Loebel image analysis system" on at least four randomly chosen areas for each sample. At least 1000 individual particles were counted for each size or volume fraction measurement. Reaction temperatures such as the determination ofsolidus-liquidus points within the sintering temperature range were determined by differential thermal analysis (DTA). Hardness tests were conducted on the various composites in the as sintered condition and also after the sintered specimen had been subjected to a double sintering tempering treatment each for one hour at 550°C. The purpose of this treatment was to transform any retained austenite and to develop any secondary hardening that might be available from the solution of alloy carbides during the sintering treatment: 550°C was found to be about the optimum tempering temperature for this type of high speed steel by previous experiment.
3. RESULTS 3.1 Sintering behaviour The effects of sintering for one hour at different sintering temperatures on the density of various composites containing 5 wt% of either the individual particle or combined particulate additions are shown in Figure 1. These results were normalised to represent a fraction of the full/maximum density attainable by each composite which was virtually free from porosity. In all cases ceramic/solid lubricant particle additions raised the sintering temperatures required to reach full density, particularly for the case of TiC and for the combined TiC/MnS additions. Actual maximum density values obtained for each composite compared with a theoretical density calculated by "a simple rule of mixtures" are shown in Table 2. Theoretical density p = 1/(Ca/Pa + Cb/Pb + Cc/Pc ) p = theoretical density: C = concentration (components a, b, c): P ab.c = component density
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The theoretical density values calculated by rule of mixtures agreed well the measured maximum density values, see Table 2.
i.O
.9
1-7] M3/2 + 0.25% C b a s e l i n e Z~ M3/2 + 5% T i C [5um] X7 M3/2 + 5% MnS
M3/2 + 5% CaF Jr M3/2 + 5% TiC [15um]
X M3/2 + 51 lq~oC [4um] i
1180
1220
i
1260 Sintering Temperature C
i
1300
Figure 1. Effects ofsintering temperature on the density of M3/2 HSS containing various ceramic carbide and solid lubricant additions, densities normalised to a fraction of the maximum full density achieved, see Table 2.
Table 2. Actual densities and theoretical densities for fially sintered M3/2 HSS/cea".~e/solid lubricant composites. Composite: Maximum Theoretical density g/co density, g/cc M3/2 HSS. 8.15 8.15 M3/2 HSS + 5% TiC 7.88 7.89 /v13/2 HSS + 5% NbC 8.09 8.12 M3/2 HSS + 5% MnS 7.78 7.75 M3/2 HSS + 5% CaF 2 7.58 7.56 M3/2 HSS + 5% TiC 7.42 7.51 + 5% MnS. M3/2 HSS + 5% TiC 7.19 7.34 + 5% CaF2. M3/2 HSS + 5% NbC 7.76 7.72. + 5% MnS. Isothermal studies of the effects of sintering time and temperature on the densification in the various composite materials also indicated that densification rates were extremely slow in all materials at temperatures below approximately 1225oc but that extremely rapid densification occurred at temperatures above this. Density versus time results, from which it was possible to determine an "Activation Energy" ' for densification, indicated a significant slope increase and hence increase in the "Activation Energy" at sintering temperatures below 1226 to 1235oc, see Figure 2. This temperature range was shown by DTA, to coincide with the "solidus" temperature and the start of liquid phase formation in the composites studied and gave rise to dramatic increases in sintering rates which were greatly influenced by sintering temperature: only a slight increase in temperature significantly increased the sintering rate. The "activation energy" for sintering at this higher temperature was approximately 1040 kJ/mol.
J.D. Bolto,I. A.J. G~,lt I Jo,,',~al qf Materials Processing Technology 56 (1996) 136-147
139
0 L
[~ M3/2 HSS baseline
n
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/I M3/2 HSS + 5% MnS
I -4
~7 M3/2 HSS + 5% TiC
[15um l
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M3/2 HSS + 5% TiC
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.
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. 6.50
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in M6C carbide fraction . 6.55
.
.
.
. '~
6.60 6.65 6.70 I/T x I0": T Temp: K
6.75
6.80
6.85
Figure 2. Arrhenius plot, Ln t [t = time for 50% densification] versus l/T, [T = sintering temperature OK] 3.1 Microstructure
3.1.1 M3/2 HSS + TiC Major changes occurred to the basic microstructure of the M3/2 high speed steel matrix as a result of introducing either ceramic or solid lubricant particles, especially for alloys containing TiC and the combination of TiC/MnS particles. The structures of sintered M3/2 HSS baseline material were typically composed of a martensitie matrix, plus retained austenite, together with both M6C and MC type carbides, both as rounded carbide particles within the grains and as angular or continuous carbides at prior austenite grain boundaries, Coarse M6C eutectie grain boundary carbides and rod like M2C carbides were produced by over sintering at excessively high sintering temperatures. With the addition of titanium carbide (TIC) significant changes to the relative proportions of M6C and MC carbides within the steel matrix occurred, and it was evident that M6C carbides progressively disappeared from the microstructure, to be replaced by an MC type carbide, see Figure 3. A 16 ~7 -~ X
MC carbide in M3/2 HSS + TiC [15urn] MC carbide in M3/2 HSS + TiC [5urn] MGC carbide in M3/2 HSS + TiC [15um] M~C carbide in M3/2 HS8 + TiC [Sum] MC carbide in M3/2 HSS 12 "~ M~C carbide in M3/2 HSS
J /x
1180
1220
1260
Sintering Temperature
130c
C
Figure 3. The effect of sintering temperature on the volume ti-action of MC and M6C carbide particles retained after sintering into an M3/2 HSS plus TiC composite.
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The replacement of MrC carbides was more noticeable at higher sintering temperatures and in the materials containing the finer [5pm] grade of titanium carbide. The identities of the MC carbide phases formed by reaction between the M3/2 HSS and the TiC carbides was determined by X ray diffraction and by X ray energy dispersive analysis and were shown in fact to grow at the steel matrix/TiC particle interface as two distinctive types with different composition and lattice parameters to each other. MC carbides were also present as isolated carbide particles within the prior austenite grain boundaries. Typical compositions for the two MC carbide types are shown in Table 3 alongside analysis data for the MC carbide formed inside the steel matrix. Table 3 Xrav enerzv dispersive analysis data for the MC carbide phases formed by reactions between M3/2 HSS and TiC additions. Element ato.mic % normalised, excluding carbon content. Carbide location Fe W Mo Cr V Ti TiC/HSS interface 8.5 10.5 19 3.5 42 17 TiC/HSS interface 31 8.5 15 4 32.5 10.5 HSS matrix 9.5 10 19.5 4 44 12.5 Energy dispersive analysis of this MC carbide phase also confirmed the existence of essentially different MC carbides at different sintering temperatures. MC carbides formed below the solidus temperature were essentially rich in both iron and vanadium, but also contained detectable quantities of titanium As sintering temperatures were increased, and liquid phases were formed, the titanium concentration in these MC carbides increased considerably, mainly at the expense of the chromium, tungsten, molybdenum, and vanadium normally present within the MC carbide phase. Quantitative image ~nslysis also indicated significant changes to the size and volume fraction of the TiC particles retained by the microstructure aider sintering at successively higher sintering temperatures and it became obvious that the finer of the two TiC additions, namely the [5~tm] grade, suffered some reduction both in its volume fi'action and in its average size at the higher sintering temperatures used. This indicated that some chemical dissolution of the TiC into the steel matrix had occurred, but that the effect was less pronounced with the coarser [151am] grade TiC carbide. The finer TiC composite also contained clusters of small TiC particles surrounded by a layer of MC carbide containing both the iron rich and the titanium rich carbide variety shown in table 3, see Figure 4, whereas the coarser 151am TiC particles were retained as separate particles surrounded by an interracial MC layer with occasional evidence that the MC carbide had grown into cracks formed in the TiC during compacting.
Figure 4. Growth of MC carbides around clusters of fine TiC carbide particles due to interaction between the TiC and the steel matrix., backscattered electron image.
JD. Bolton, A.J. Gant / Journal (~fMaterials Processing Technology 56 (1996) 136-147
X ray energy dispersive analysis also showed that only fractional amounts of titanium were retained in solution in either the steel matrix or in any & t h e remaining M6C carbides formed in the sintered structure.
3.1.2 M3/2 + NbC Niobium carbide additions resulted in the formation of large carbide agglomerates and the presence of carbide clusters/networks in the sintered structure, see Fig 5.
Figure 5. Clusters of niobium rich MC carbides formed by reaction with the steel matrix during sintering. Clusters originated because of poor mixing between the high speed steel powder and its NbC carbide additive, backscattered electron image.
These clusters originated from the NbC carbide particle addition but had undergone considerable change in composition as a result of the sintering process. X ray energy dispersive analysis indicated the gradual replacement of niobium with vanadium and also by molybdenum as sintering.temperatures were increased and that this replacement began at temperatures below those at which liquid phases were formed during sintering, see Table 4. Table 4. Typical X ray energy dispersive analyses of the NbC carbide phase for different sintering temperatures. Sintering Temp:-°C E.lement atomic % normalised, excluding carbon Fe W Mo Cr V Nb. 1210 3 3 11 1 12 71 @1220 1250
3 4
3 5
12 14
1 1
14 17
68 59
® below solidus temperature. Changes to the steel matrix microstructure were less pronounced for the NbC containing composites than were shown for the TiC type composites and no evidence of preferential replacement of the M6C matrix carbide phase by niobium rich MC carbides in the M3/2 steel matrix was found. The matrix carbide population of the steel matrix was virtually unchanged from that found in a normal M3/2 high speed steel and there was little evidence of niobium dissolving either in the steel matrix or in the M6C carbide phase, e.g. Table 5.
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Table 5. Composition of steel matrix and M C c.arbides in the M3/2 HSS/NbC composite. Element, atomic % normalised, excludingcarbon. Phase Fe W Mo Cr V Nb steel matrix 93 1 1 4 ] 0 M6C 5] 16 22 6 4 1 3. 1.3 M3/2 + MnS/CaF 2 Less pronounced changes occurred to the microstructure of the M3/2 HSS matrix when either MnS or CaF2 solid lubricants were individually introduced into the sintered material. The relative proportions of either MC or M6C carbides formed at different sintering temperatures remained comparable with that produced in the basic M3/2 HSS steel. No evidence was found for any diffusional layers or of new phase formation at the interface between the highspeed steel matrix and either its MnS or CaF2 particle addition. Only slight wetting between the high speedsteel and the solid lubricant particle seemed to have occurred. Energy dispersive analysis also confumed that little chemical interaction between either MnS or CaF2 and the steel occurred except for some slight diffusion of iron from the matrix into the CaF2 and some diffusion of molybdenum and chromium into the MnS solid lubricant particles, see Table 6. Table 6. X ray energy dispersive.analyses for the. solid lubricant. Solid lubrieant Element, at.omit % Fe W Mo Cr V Ca CaF 2 before sintering. 0.1 0 0.1 0 0.1 99.8 CaF2 after sintering. 4.2 0.1 0.1 0.2 0.1 95.3 MnS before sirrtering. 0.01 0.06 0 0.01 0.0 0.69 MnS after sintering. 0.3 0 5.6 2.4 0.7
Mn
S
0 0 49.1 43.1
0 0 49.8 48.1
Despite this apparent lack of reaction between the steel and either of the solid lubricants, noticeable changes to the size and shape of both types of solid lubricant were produced by sintering at high temperatures. Both types of particle grew in size and became more rounded in shape as a result of this high temperatures sintering, see Figure 6. 18 P a
r
15
I"7 M3/2 A M3/2 ~7 M3/2
HSS + 5Z MnS [high purity] HSS + 5% MnS [impure] HSS + 5X CaF
t i
12
c
1 e
0
1180
!
1200
i
i
~
1220 1240 1260 Sintering Temperature C
i
1280
1300
Figure 6. The effect of sintering temperature on the size of MnS/CaF2 particles dispersed in the sintered high speed steel matrix
J.D. Bolton, A.J. Gant / Journal oJ Materials Processing Technology 56 (1996) 136-147
3.1.4 M3/2 + TiC or NbC + CaF 2 or MnS Composites which contained both TiC and CaF2 additions produced similar microstructures to those found in the M3/2 + TiC composites; the major effect being that M6C carbides were replaced by the MC carbide phases in the same manner as that previously described for ~he M3/2 + TiC composite. The CaF 2 showed no reaction other than to undergo the coarsening already described. This was not true of the composites to which both TiC and MnS were added and several new microstructural effects were observed, especially when high sintering temperatures, necessary for producing full density, were used. The principal effects on microstructure was the appearance of a eutectoid-like structure composed of M6C carbide and austenite (martensite on cooling to room temperature) which was probably the result of the eutectoid decomposition of delta ferrite into austenite and M6C carbides, [4]. A needle like phase was also shown to occur in specimens sintered above approximately 1290oc, see Figure 7, which was analysed by E.D.A as containing the composition shown in Table 7.
Figure 7. Formation of Ti4C2S 2 as a result o f sintering at above 1280oc in an M3/2 HSS + TiC + MnS composite, backsca~ered electron image.
Table 7 Typical composition of the Ti4C2SS2 phase formed in the M3/2 HSS/TiC/MnS composites. Element, atomic % Fe V Ti Mn %S 3.1 10.3 52.2 0.5 33.1 Sintering o f the at a high temperature of 1320oc produced no evidence o f any similar reaction in the M3/2 + NbC + M_nS composite and no carbo-sulphide phase was formed by reaction between the NbC and MnS particles. The high sintering temperature did, however, produce significant quantities of M6C carbide + austenite grain boundary eutectic. Although it was not possible to carry out quantitative analysis of the carbon content in this needle like structure a significant presence of carbon was found by using a windowless E.D.A. detector.. X ray diffraction patterns also confirmed the possibility that the identity of this phase could be the Ti4C2S 2 phase also reported in some titanitnn stabilised stainless steels,[5]. Such a compound is thermodynamically more stable than either MnS or TiC at all temperatures, [6,7,8].
3.2 Hardness Typical hardness values for a variety of ceramic carbide/solid lubricant composites sintered to full density are shown in Table 8.
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Table 8. Hardness of M3/2 HSS/ceramic carbide/solid lubricant composites. Composite: Sinter Treatment: Hardness H.v30_Q_k~k. As simered temnered 550-°_C M3/2 HSS + 0.25%C. 1240oc l hr 555. 663. M3/2 HSS + 0.25%C 1260°C lhr 584. 661 + 5% TiC[5~tm]. M3/2 HSS + 0.25%C 1260°C lhr 593. 695. + 5% TIC[151am]. M3/2 HSS + 0.25%C 1280oc lhr 532. 580. + 5% TiC[51am] + 5% MnS. M3/2 HSS + 0.25%C 1280°C 1hr 492 574. + 5% TIC[151am] + 5% MnS. M3/2 HSS + 0.25%C 1280oc lhr 484. 530. + 5% TiC[51am]. + 5% CaF2. M3/2 HSS + 0.25%C 1280oc lhr 479. 541. + 5% TiC[151am] + 5% CaF2. M/2 HSS + 0.25%C 1255oc lhr 600. 720. + 5% NbC[4tam]. M3/2 HSS + 0.25%C 1320oc lhr 574. 660. + 5% NbC[4~tm] + 5% MnS.
4. DISCUSSION 4.1 Sintering Mechanisms The sintering kinetics data support the premise that the major densification process in both the baseline M3/2 HSS and its composite derivatives occurred by a "super-solidus liquid phase sintering", (SSLPS), mechanism,[9]. Densification rates at temperatures below approximately 1225oc were reliant on solid state diffusion and were consequently extremely slow. Once liquid phases began to appear at temperatures above the "solidus" the densification rate became extremely rapid and showed a marked dependence on sintering temperature. Each of the ceramic/solid lubricant particulate additions also slightly raised the temperature at which SSLPS sintering began, but for apparently different reasons. The effects of MnS in raising the sintering temperature was almost certainly related to the high oxygen concentration of the original MnS powder, e.g., 5.9 wt%, which caused the solidus temperature of the M3/2 high speed steel to be raised by the decarburisation of the steel during sintering. Use of a purer grade of MnS powder containing only 1.1 wt% of oxygen was successful in lowering the sintering temperature for this M3/2 HSS/MnS composite and confirmed that decarburisation was the major cause of high sintering temperatures in this type of composite. Micro-hardness tests on the martensitic matrix structure for the various composites together with metaUographic evidence that the steel matrix contained low carbon lath type martensite in place of the more acicular high carbon martensite formed in the other composites further supported the idea that decarburisation was responsible for raising the sintering temperature in composites containing the impure grade of MnS The effects of TiC and NbC additions in raising the sintering temperature were probably caused by the dissolution of titanium or niobium into the steel matrix during the period spent in heating up to the sintering temperature, which also raised the solidus temperature.
.I.D. Bolton, A.J. Gant / Joun~al t~fMaterials Processing Technology 56 (1996) 136-147
4.2 Sintering kinetics A valuable insight into the possible mechanism of sintering was also gained by applying the Kingery liquid phase sintering model,[10] to the isothermal data used to determine sintering kinetics. At sintering temperatures below those at which liquid phase was present, e.g. 1225°C, Kingery slopes of well below 1/5th were obtained and supported the proposalof that solid state diffusion was the major cause of densification. Similar shallow Kingery slopes were also obtained at very high sintering temperatures, where densification was so rapid that full densities were achieved even during the stage when samples were still being heated up to their sintering temperature. Thus complete densification had occurred before the effects of isothermal holding on density could be properly measured. Only within an intermediate sintering range, just above the solidus temperature, was the true slope formed by a liquid phase sintering reaction shown. Within this sintering temperature range Kingery slopes of around l/5th were obtained and suggested that the predominant mechanism for densification consisted of the solution re-precipitation between some solid and its surrounding liquid phase. Quantitative measurements of the effects of sintering time/temperature on the relative proportions of MC and M6C carbides together with the metallographic evidence of a gradual transition from M6C into MC carbides at high sintering temperatures prompted the suggestion that this solution re-precipitation reaction could be the dissolution of M6C carbide into the liquid followed by re-preclpitation as an MC carbide. Such a process was prevalent amongst the TiC containing composites due to the natural encouragement given to MC carbide formation by the solution of titanium into the steel, and activation energy determinations for the changes in volume fraction of the MC and M6C carbides relative to sintering time for the M3/2 + TiC composites supported this idea. This activation energy was remarkably close to that determined for densification, see Figure 2, i.e. 1080 kJ/mol. 4.3 Chemical stability of the ceramic/solid lubricant particles Both types of solid lubricant, namely MnS and CaF 2 were shown to be chemically stable when sintered into a high speed steel matrix with little evidence that they dissolved in the steel. In the presence of TiC, however, there were indications that the MnS was thermodynamically unstable and that reaction between MnS and TiC could produce a new Ti4C2S 2 phase, [7]. Although such a reaction was thermodynamically possible at all temperatures the reaction kinetics were such that the Ti4C2S 2 phase was only observed in any significant quantities in samples sintered above 1280oc. Little is yet known concerning the possible effects of this phase on the wear/friction properties of the high speed steel. Both of-the ceramic carbides of TiC and NbC were found to react with the high speed steel matrix but the extent of this reactivity differed between the two types of carbide and was also a function of their particle size. Significant reactions were also only evident at sintering temperatures above those at which liquid phases were present so that it was possible to speculate that any new phases formed were as a result of reaction between the TiC or NbC carbide and the liquid phase formed from the steel matrix. In the case of TiC, reaction with this liquid phase at the surface of the TiC particles was naturally more evident with the finer 5gm grade of TiC particle and caused the formation of an interfacial layer of cubic MC mono-carbide phase which surrounded the TiC particles. Two distinctive varieties, either an iron rich or a vanadium rich form of this MC carbide were produced. The tendencies for titanium, vanadium, or niobium to form an isomorphous series of cubic MC carbides, mutually soluble in one another, is well known, [ 11 ], and would of course explain why MC alloy carbides were formed at the TiC/steel interface. Less clear, however, was the fact that the MC carbides did not form as a complete range of compositions involving a gradual change in their titanium and vanadium composition. Instead only two general compositions for this MC carbide arose with fairly specific metal atom/carbon ratios, i.e., M2X3C 5 (lattice parameter a o ~ 0.427 rlm)or M3X2C5, (lattice parameter a o ~ 0.429 qm); where M denotes the combined total of metallic elements capable of forming cubic MC mono-carbides, e.g., Ti Nb, and V, X denotes the combined total of metallic elements which do not normally form cubic MC carbides, e.g., Fe, Cr,W,Mo, and C is the total carbon, see Table 3. No such interfacial layers were produced in the composites containing NbC and unlike the TiC composites little or none of the original carbide addition was retained after sintering. Niobium carbide particles were totally converted into an alloy MC carbide by the diffusion of alloying elements from the high speed steel into the NbC with no interfacial layer or diffusion gradient from the centre to the outside of each particle. This possibly reflected a greater reactivity for the NbC particle with its surrounding steel matrix, compared with the
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reactivity of TiC, but also demonstrated the role of finer particle size in speeding up any reaction. Progressively more replacement of the niobium by both vanadium and by molybdenum occurred in the carbide as sintering temperatures were increased, but a ratio ofM4XC 5 was maintained between the MC forming and the non MC carbide forming metallic elements. Differences in behaviour between TiC and NbC also existed for the manner in which solution and diffusion of the ceramic carbide addition into the high speed steel affected the microstructure steel matrix. Slight solution of titanium into the steel caused the disappearance of the usual M6C carbides from the steel structure by encouraging MC carbides to form in their place. This did not occur in the NbC materials either because niobium has such limited solubility in the austenite at the sintering temperatures employed,[12] or because its diffusion rates are so much slower than those of titanium in the austenite phase. 4.4 Particle distribution The agglomeration and clustering effects observed with the finer TiC or NbC particles were partially due to the agglomeration of fine carbide particles during the powder mixing stage but this was not the only cause. Dramatic increases in size for the MnS and CaF 2 particles could also be explained by the presence of a liquid phase at the higher sintering temperatures. The smaller ceramic and solid lubricant particles, some of which originated from fragmentation during die compacting stage, were picked up by the liquid phase as it flowed via any interconnecting pores through the compact at the earlier sintering stage. These particles became trapped in the larger pore interstices to become agglomerates which subsequently either coalesced into large particles, as in the case of the MnS and CaF 2 additions, or were formed into clusters of small particles surrounded by a second phase due to their chemical reaction with the liquid phase. Such was the case for the fine grained TiC [5 !am] carbide additions. 4.5 Hardness The addition of hard TiC or NbC ceramic carbides to the high speed steel produced only a small increase in hardness compared with that of the baseline high speed steel both in the as-sintered and in the as-sintered plus double tempered condition, see Table 8, and the double tempering treatment produced higher hardness values than those of the as- sintered materials because of the removal of retained austenite and the development of secondary hardening by carbide precipitation. Significant reductions in hardness were produced by the addition ofMnS/CaF 2 solid lubricant to the various composite mixtures. The small increases in hardness, associated with the addition of TiC/NbC, were presumably due to the small volume fractions of carbide added and somewhat surprisingly the NbC was the slightly more effective of the two carbides, despite the fact that the volume fraction of NbC carbide added to the steel was smaller for the NbC than for the TiC carbide, e.g., 5 wt% ~ 8.3 vol% TiC and 5.4 vol% NbC. Higher hardness values for the NbC composite were also at odds with the observed changes in microstructure, namely that much of the TiC additions were retained in the structure and that TiC encouraged the formation of the harder MC carbides in place of the relatively soft M6C carbides, [4], in the steel matrix: pure NbC particles were not retained and no preferential MC carbide fo.rmation occurred within the steel matrix of the NbC composite. Poor hardness values were achieved in composites containing the solid lubricant and this was partially due to the decarburisation of the steel caused by the introduction of oxygen as an impurity, at least in the case of MnS additions. The more important reason for poor hardness was that both types of solid lubricant tended to collapse under the hardness indentor giving similar effect to that of porosity in under-sintered materials. Thus hardness readings in the composites containing solid lubricants were to some extent not a true value.
5. CONCLUSIONS [1] Metal matrix composites based on high speed steel and containing particulate additions of hard TiC or NbC ceramic plus MnS/CaF 2 solid lubricant particles were successfully sintered to full density. TiC particles additions were essentially preserved as a dispersion in their original chemical form, but the NbC addition was changed into an alloy MC carbide by the diffusion of alloying elements from the high speed steel. This was possibly encouraged by the fine grain size of the NbC used.
J.D. Bolton, A.J. Gant / Journal of Materials Processing Technology 56 (1996) 136-147
[2] Both types of ceramic carbide were chemically bonded into the high speed steel matrix, either as a result of forming an interracial vanadium rich MC carbide layer at the TiC/steel interface, or by diffusion from the steel into the ceramic carbide, as in the case of NbC carbide. [3] Densification was achieved by a liquid phase sintering reaction which appeared to involve solution reprecipitation as the predominant sintering process. Solution of the matrix M6C carbides together with the added TiC/NbC carbides, into a liquid phase, followed by a re-precipitation to form MC carbides was a possible reaction for the solution re-precipitation reaction involved. [4] Although the solid lubricant particles did not react chemically with the steel matrix during sintering they underwent considerable coarsening in size and did not appear to be all that well wetted by the steel matrix. [5] Double tempering at 550°C atter the sintering treatment raised hardness levels but only slight increases in hardness were brought about by the addition of hard ceramic carbide, and considerable, but apparent, reductions in hardness were caused by the addition ofMnS/CaF 2 solid lubricant.
6. REFERENCES [1] Trudel.Y, Cliloglu.C, Tremblay.S, Mod. Dev. in P/M. MPIF Princeton, NY, 15, 1985, 775. [2] Yuko. T, Kazunori U, Lubrication Engineering, 31, 8, (1976), 402. [3] Price.W.J.S, Rebbeck.M.M, Wronski.A.S, Amen.A.S, Powder Metallurgy, 28 (1985), 1, [4] Hoyle.G, "High Speed Steels", Butterworth, London, (1988). [5] BaI1.C.J, Metal Science, 18, (1984), 577. [6] Bandi.W.R, Krapf.G, Analyst, 104, (1979), 812. [7] Lui.W.J, Jonas.J.J, Bouchard.D, Bale.C.W, I.S.I.J Int., 30, (1990), 985. [8] Kubachewski.O, Alcock.C.B, Metallurgical Thermochemistry, 5th edition, 1979, Pergamon, London. [9] Wright.C.S, Powder Metallurgy, 32, (1989), 114. [10] Kingery.W.D, J. Appl. Physics, 30, (1959), 301. [11] Goldschmidt.H.J, J. Iron & Steel Inst. 160, 1948, 345. [12] Narita.K, Trans.I S I J, 15, 1975, 145.
147
ELSEVIER
Journal of Materials Processing Technology 56 (1996) 136-147
Materials Processing Technology
LIQUID PHASE SINTERING OF METAL MATRIX COMPOSITES CONTAINING SOLID LUBRICANTS
J.D. Bolton and A.J. Gant Department of Mechanical Eng., University of Bradford, U.K.
ABSTRACT Attempts to improve the wear properties of sintered high speed steels have been made by the simultaneous addition of hard ceramic particles (TIC) or (NbC) alongside particles with solid lubricative properties such as (MnS) or (CaF2). Preliminary investigations carried out to study any interactions between such particles and a sintered high speed steel matrix indicated that the steel matrix carbides were modified by a chemical reaction which occurred between the added ceramic particle and its high speed steel matrix; the result of which was to give a substantial improvement in hardness. The solid lubricant remained chemically unaltered by the sintering process, and tended to reduce hardness. Both types of particulate addition raised the sintering temperature needed to achieve full density due to their effect on solidus/liquidus temperatures. Studies of the sintering kinetics in such materials confirmed that densification was due to a liquid phase sintering mechanism and that solution/re-precipitation of carbide phases played a major role in the densification process. Coarsening of the solid lubricant particles into large agglomerates occurred during sintering by a process of the transportation of smaller particles into pore interstices by the liquid phase.
1. INTRODUCTION
Preliminary investigations were carried out to study any interactions between particulate additions of TiC/NbC hard ceramic carbide, MnS/CaF 2 solid lubricants to a sintered high speed steel matrix in the hope of forming a metal matrix composite with enhanced wear properties. Emphasis was placed on determining any chemical reactions between the added particle and its high speed steel matrix plus any subsequent effects on the microstructure and sintering kinetics of the high speed steel. The chemical stability and distribution of the particulate addition in relation to the steel matrix were also studied.
2. EXPERIMENTAL WORK
Sintered composites were all basically prepared by mixing a water atomised M3 class 2 high speed steel (45~tm grain size) powder of the composition shown in Table 1, with one or more of the following powder ceramic carbides or solid lubricants 0924-0136/96/$15.00 © 1996 Elsevier Science S.A. All rights reserved SSD10924-0136 ( 95 ) 01829-4
.I.D. Bolton, A.J. Gant / Journal ¢~fMaterials Processing Technology 56 (1996) 136-147
Table 1. Composition of M3/2 high speed steel powder, %C %Cr %W %Mo %V 1.10 4.02 6.05 5.80 2.91 Manganese sulphide (MnS); low oxygen grade. Manganese sulphide (MnS); average particle size 10gm. Calcium fluoride (CaF2); average particle size 30gin. Titanium carbide (TIC); average particle size 51am. Titanium carbide (TIC); average particle size 151am. Niobium carbide (NbC); average grain size 41am. Graphite. Composite mixtures were made, to which 5 wt% additions of either solid lubricant or hard ceramic particles were blended by dry mixing in a Y cone blender for 15 rains: 0.25 wt% of graphite was also added to all batches to assist the sintering process, [3]. Compacts 4ram thick by 15 mm square were die pressed to green densities of approximately 70% theoretical, at pressure of 800 M.Pa and were sintered by holding for a fixed period under a vacuum (0.15 labar) after heating at 10OC/min to a choice of different sintering times of between 2 mins and four hours at a range of temperatures between 1180 and 1340°C. The as-sintered densities of all specimens were measured by weighing and by determining the volume by water displacement, but specimens were first coated with lacquer to prevent water penetration into surface pores during the weighing process. An appropriate correction was made for the weight and volume of this lacquer coating. Characterisation of microstruetures and the identification of phases present was performed by both optical and manning electron microscopy, assisted by the use of X ray energy dispersive analysis (E.D.A), back scattered electron image contrast (BE.I), and by some X-ray diffraction data. Quantitative measurements of microstructural features such as particle size and volume fraction for both the solid lubricant, ceramic particles, and of the high speed steel carbide phases were carried out from back scattered electron images using the "Magicscan, J Loebel image analysis system" on at least four randomly chosen areas for each sample. At least 1000 individual particles were counted for each size or volume fraction measurement. Reaction temperatures such as the determination ofsolidus-liquidus points within the sintering temperature range were determined by differential thermal analysis (DTA). Hardness tests were conducted on the various composites in the as sintered condition and also after the sintered specimen had been subjected to a double sintering tempering treatment each for one hour at 550°C. The purpose of this treatment was to transform any retained austenite and to develop any secondary hardening that might be available from the solution of alloy carbides during the sintering treatment: 550°C was found to be about the optimum tempering temperature for this type of high speed steel by previous experiment.
3. RESULTS 3.1 Sintering behaviour The effects of sintering for one hour at different sintering temperatures on the density of various composites containing 5 wt% of either the individual particle or combined particulate additions are shown in Figure 1. These results were normalised to represent a fraction of the full/maximum density attainable by each composite which was virtually free from porosity. In all cases ceramic/solid lubricant particle additions raised the sintering temperatures required to reach full density, particularly for the case of TiC and for the combined TiC/MnS additions. Actual maximum density values obtained for each composite compared with a theoretical density calculated by "a simple rule of mixtures" are shown in Table 2. Theoretical density p = 1/(Ca/Pa + Cb/Pb + Cc/Pc ) p = theoretical density: C = concentration (components a, b, c): P ab.c = component density
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The theoretical density values calculated by rule of mixtures agreed well the measured maximum density values, see Table 2.
i.O
.9
1-7] M3/2 + 0.25% C b a s e l i n e Z~ M3/2 + 5% T i C [5um] X7 M3/2 + 5% MnS
M3/2 + 5% CaF Jr M3/2 + 5% TiC [15um]
X M3/2 + 51 lq~oC [4um] i
1180
1220
i
1260 Sintering Temperature C
i
1300
Figure 1. Effects ofsintering temperature on the density of M3/2 HSS containing various ceramic carbide and solid lubricant additions, densities normalised to a fraction of the maximum full density achieved, see Table 2.
Table 2. Actual densities and theoretical densities for fially sintered M3/2 HSS/cea".~e/solid lubricant composites. Composite: Maximum Theoretical density g/co density, g/cc M3/2 HSS. 8.15 8.15 M3/2 HSS + 5% TiC 7.88 7.89 /v13/2 HSS + 5% NbC 8.09 8.12 M3/2 HSS + 5% MnS 7.78 7.75 M3/2 HSS + 5% CaF 2 7.58 7.56 M3/2 HSS + 5% TiC 7.42 7.51 + 5% MnS. M3/2 HSS + 5% TiC 7.19 7.34 + 5% CaF2. M3/2 HSS + 5% NbC 7.76 7.72. + 5% MnS. Isothermal studies of the effects of sintering time and temperature on the densification in the various composite materials also indicated that densification rates were extremely slow in all materials at temperatures below approximately 1225oc but that extremely rapid densification occurred at temperatures above this. Density versus time results, from which it was possible to determine an "Activation Energy" ' for densification, indicated a significant slope increase and hence increase in the "Activation Energy" at sintering temperatures below 1226 to 1235oc, see Figure 2. This temperature range was shown by DTA, to coincide with the "solidus" temperature and the start of liquid phase formation in the composites studied and gave rise to dramatic increases in sintering rates which were greatly influenced by sintering temperature: only a slight increase in temperature significantly increased the sintering rate. The "activation energy" for sintering at this higher temperature was approximately 1040 kJ/mol.
J.D. Bolto,I. A.J. G~,lt I Jo,,',~al qf Materials Processing Technology 56 (1996) 136-147
139
0 L
[~ M3/2 HSS baseline
n
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/I M3/2 HSS + 5% MnS
I -4
~7 M3/2 HSS + 5% TiC
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0
[5um]
~
/
n
t
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M3/2 HSS + 5% TiC
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s
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~- change -16
.
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. 6.50
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\
in M6C carbide fraction . 6.55
.
.
.
. '~
6.60 6.65 6.70 I/T x I0": T Temp: K
6.75
6.80
6.85
Figure 2. Arrhenius plot, Ln t [t = time for 50% densification] versus l/T, [T = sintering temperature OK] 3.1 Microstructure
3.1.1 M3/2 HSS + TiC Major changes occurred to the basic microstructure of the M3/2 high speed steel matrix as a result of introducing either ceramic or solid lubricant particles, especially for alloys containing TiC and the combination of TiC/MnS particles. The structures of sintered M3/2 HSS baseline material were typically composed of a martensitie matrix, plus retained austenite, together with both M6C and MC type carbides, both as rounded carbide particles within the grains and as angular or continuous carbides at prior austenite grain boundaries, Coarse M6C eutectie grain boundary carbides and rod like M2C carbides were produced by over sintering at excessively high sintering temperatures. With the addition of titanium carbide (TIC) significant changes to the relative proportions of M6C and MC carbides within the steel matrix occurred, and it was evident that M6C carbides progressively disappeared from the microstructure, to be replaced by an MC type carbide, see Figure 3. A 16 ~7 -~ X
MC carbide in M3/2 HSS + TiC [15urn] MC carbide in M3/2 HSS + TiC [5urn] MGC carbide in M3/2 HSS + TiC [15um] M~C carbide in M3/2 HS8 + TiC [Sum] MC carbide in M3/2 HSS 12 "~ M~C carbide in M3/2 HSS
J /x
1180
1220
1260
Sintering Temperature
130c
C
Figure 3. The effect of sintering temperature on the volume ti-action of MC and M6C carbide particles retained after sintering into an M3/2 HSS plus TiC composite.
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The replacement of MrC carbides was more noticeable at higher sintering temperatures and in the materials containing the finer [5pm] grade of titanium carbide. The identities of the MC carbide phases formed by reaction between the M3/2 HSS and the TiC carbides was determined by X ray diffraction and by X ray energy dispersive analysis and were shown in fact to grow at the steel matrix/TiC particle interface as two distinctive types with different composition and lattice parameters to each other. MC carbides were also present as isolated carbide particles within the prior austenite grain boundaries. Typical compositions for the two MC carbide types are shown in Table 3 alongside analysis data for the MC carbide formed inside the steel matrix. Table 3 Xrav enerzv dispersive analysis data for the MC carbide phases formed by reactions between M3/2 HSS and TiC additions. Element ato.mic % normalised, excluding carbon content. Carbide location Fe W Mo Cr V Ti TiC/HSS interface 8.5 10.5 19 3.5 42 17 TiC/HSS interface 31 8.5 15 4 32.5 10.5 HSS matrix 9.5 10 19.5 4 44 12.5 Energy dispersive analysis of this MC carbide phase also confirmed the existence of essentially different MC carbides at different sintering temperatures. MC carbides formed below the solidus temperature were essentially rich in both iron and vanadium, but also contained detectable quantities of titanium As sintering temperatures were increased, and liquid phases were formed, the titanium concentration in these MC carbides increased considerably, mainly at the expense of the chromium, tungsten, molybdenum, and vanadium normally present within the MC carbide phase. Quantitative image ~nslysis also indicated significant changes to the size and volume fraction of the TiC particles retained by the microstructure aider sintering at successively higher sintering temperatures and it became obvious that the finer of the two TiC additions, namely the [5~tm] grade, suffered some reduction both in its volume fi'action and in its average size at the higher sintering temperatures used. This indicated that some chemical dissolution of the TiC into the steel matrix had occurred, but that the effect was less pronounced with the coarser [151am] grade TiC carbide. The finer TiC composite also contained clusters of small TiC particles surrounded by a layer of MC carbide containing both the iron rich and the titanium rich carbide variety shown in table 3, see Figure 4, whereas the coarser 151am TiC particles were retained as separate particles surrounded by an interracial MC layer with occasional evidence that the MC carbide had grown into cracks formed in the TiC during compacting.
Figure 4. Growth of MC carbides around clusters of fine TiC carbide particles due to interaction between the TiC and the steel matrix., backscattered electron image.
JD. Bolton, A.J. Gant / Journal (~fMaterials Processing Technology 56 (1996) 136-147
X ray energy dispersive analysis also showed that only fractional amounts of titanium were retained in solution in either the steel matrix or in any & t h e remaining M6C carbides formed in the sintered structure.
3.1.2 M3/2 + NbC Niobium carbide additions resulted in the formation of large carbide agglomerates and the presence of carbide clusters/networks in the sintered structure, see Fig 5.
Figure 5. Clusters of niobium rich MC carbides formed by reaction with the steel matrix during sintering. Clusters originated because of poor mixing between the high speed steel powder and its NbC carbide additive, backscattered electron image.
These clusters originated from the NbC carbide particle addition but had undergone considerable change in composition as a result of the sintering process. X ray energy dispersive analysis indicated the gradual replacement of niobium with vanadium and also by molybdenum as sintering.temperatures were increased and that this replacement began at temperatures below those at which liquid phases were formed during sintering, see Table 4. Table 4. Typical X ray energy dispersive analyses of the NbC carbide phase for different sintering temperatures. Sintering Temp:-°C E.lement atomic % normalised, excluding carbon Fe W Mo Cr V Nb. 1210 3 3 11 1 12 71 @1220 1250
3 4
3 5
12 14
1 1
14 17
68 59
® below solidus temperature. Changes to the steel matrix microstructure were less pronounced for the NbC containing composites than were shown for the TiC type composites and no evidence of preferential replacement of the M6C matrix carbide phase by niobium rich MC carbides in the M3/2 steel matrix was found. The matrix carbide population of the steel matrix was virtually unchanged from that found in a normal M3/2 high speed steel and there was little evidence of niobium dissolving either in the steel matrix or in the M6C carbide phase, e.g. Table 5.
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Table 5. Composition of steel matrix and M C c.arbides in the M3/2 HSS/NbC composite. Element, atomic % normalised, excludingcarbon. Phase Fe W Mo Cr V Nb steel matrix 93 1 1 4 ] 0 M6C 5] 16 22 6 4 1 3. 1.3 M3/2 + MnS/CaF 2 Less pronounced changes occurred to the microstructure of the M3/2 HSS matrix when either MnS or CaF2 solid lubricants were individually introduced into the sintered material. The relative proportions of either MC or M6C carbides formed at different sintering temperatures remained comparable with that produced in the basic M3/2 HSS steel. No evidence was found for any diffusional layers or of new phase formation at the interface between the highspeed steel matrix and either its MnS or CaF2 particle addition. Only slight wetting between the high speedsteel and the solid lubricant particle seemed to have occurred. Energy dispersive analysis also confumed that little chemical interaction between either MnS or CaF2 and the steel occurred except for some slight diffusion of iron from the matrix into the CaF2 and some diffusion of molybdenum and chromium into the MnS solid lubricant particles, see Table 6. Table 6. X ray energy dispersive.analyses for the. solid lubricant. Solid lubrieant Element, at.omit % Fe W Mo Cr V Ca CaF 2 before sintering. 0.1 0 0.1 0 0.1 99.8 CaF2 after sintering. 4.2 0.1 0.1 0.2 0.1 95.3 MnS before sirrtering. 0.01 0.06 0 0.01 0.0 0.69 MnS after sintering. 0.3 0 5.6 2.4 0.7
Mn
S
0 0 49.1 43.1
0 0 49.8 48.1
Despite this apparent lack of reaction between the steel and either of the solid lubricants, noticeable changes to the size and shape of both types of solid lubricant were produced by sintering at high temperatures. Both types of particle grew in size and became more rounded in shape as a result of this high temperatures sintering, see Figure 6. 18 P a
r
15
I"7 M3/2 A M3/2 ~7 M3/2
HSS + 5Z MnS [high purity] HSS + 5% MnS [impure] HSS + 5X CaF
t i
12
c
1 e
0
1180
!
1200
i
i
~
1220 1240 1260 Sintering Temperature C
i
1280
1300
Figure 6. The effect of sintering temperature on the size of MnS/CaF2 particles dispersed in the sintered high speed steel matrix
J.D. Bolton, A.J. Gant / Journal oJ Materials Processing Technology 56 (1996) 136-147
3.1.4 M3/2 + TiC or NbC + CaF 2 or MnS Composites which contained both TiC and CaF2 additions produced similar microstructures to those found in the M3/2 + TiC composites; the major effect being that M6C carbides were replaced by the MC carbide phases in the same manner as that previously described for ~he M3/2 + TiC composite. The CaF 2 showed no reaction other than to undergo the coarsening already described. This was not true of the composites to which both TiC and MnS were added and several new microstructural effects were observed, especially when high sintering temperatures, necessary for producing full density, were used. The principal effects on microstructure was the appearance of a eutectoid-like structure composed of M6C carbide and austenite (martensite on cooling to room temperature) which was probably the result of the eutectoid decomposition of delta ferrite into austenite and M6C carbides, [4]. A needle like phase was also shown to occur in specimens sintered above approximately 1290oc, see Figure 7, which was analysed by E.D.A as containing the composition shown in Table 7.
Figure 7. Formation of Ti4C2S 2 as a result o f sintering at above 1280oc in an M3/2 HSS + TiC + MnS composite, backsca~ered electron image.
Table 7 Typical composition of the Ti4C2SS2 phase formed in the M3/2 HSS/TiC/MnS composites. Element, atomic % Fe V Ti Mn %S 3.1 10.3 52.2 0.5 33.1 Sintering o f the at a high temperature of 1320oc produced no evidence o f any similar reaction in the M3/2 + NbC + M_nS composite and no carbo-sulphide phase was formed by reaction between the NbC and MnS particles. The high sintering temperature did, however, produce significant quantities of M6C carbide + austenite grain boundary eutectic. Although it was not possible to carry out quantitative analysis of the carbon content in this needle like structure a significant presence of carbon was found by using a windowless E.D.A. detector.. X ray diffraction patterns also confirmed the possibility that the identity of this phase could be the Ti4C2S 2 phase also reported in some titanitnn stabilised stainless steels,[5]. Such a compound is thermodynamically more stable than either MnS or TiC at all temperatures, [6,7,8].
3.2 Hardness Typical hardness values for a variety of ceramic carbide/solid lubricant composites sintered to full density are shown in Table 8.
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Table 8. Hardness of M3/2 HSS/ceramic carbide/solid lubricant composites. Composite: Sinter Treatment: Hardness H.v30_Q_k~k. As simered temnered 550-°_C M3/2 HSS + 0.25%C. 1240oc l hr 555. 663. M3/2 HSS + 0.25%C 1260°C lhr 584. 661 + 5% TiC[5~tm]. M3/2 HSS + 0.25%C 1260°C lhr 593. 695. + 5% TIC[151am]. M3/2 HSS + 0.25%C 1280oc lhr 532. 580. + 5% TiC[51am] + 5% MnS. M3/2 HSS + 0.25%C 1280°C 1hr 492 574. + 5% TIC[151am] + 5% MnS. M3/2 HSS + 0.25%C 1280oc lhr 484. 530. + 5% TiC[51am]. + 5% CaF2. M3/2 HSS + 0.25%C 1280oc lhr 479. 541. + 5% TiC[151am] + 5% CaF2. M/2 HSS + 0.25%C 1255oc lhr 600. 720. + 5% NbC[4tam]. M3/2 HSS + 0.25%C 1320oc lhr 574. 660. + 5% NbC[4~tm] + 5% MnS.
4. DISCUSSION 4.1 Sintering Mechanisms The sintering kinetics data support the premise that the major densification process in both the baseline M3/2 HSS and its composite derivatives occurred by a "super-solidus liquid phase sintering", (SSLPS), mechanism,[9]. Densification rates at temperatures below approximately 1225oc were reliant on solid state diffusion and were consequently extremely slow. Once liquid phases began to appear at temperatures above the "solidus" the densification rate became extremely rapid and showed a marked dependence on sintering temperature. Each of the ceramic/solid lubricant particulate additions also slightly raised the temperature at which SSLPS sintering began, but for apparently different reasons. The effects of MnS in raising the sintering temperature was almost certainly related to the high oxygen concentration of the original MnS powder, e.g., 5.9 wt%, which caused the solidus temperature of the M3/2 high speed steel to be raised by the decarburisation of the steel during sintering. Use of a purer grade of MnS powder containing only 1.1 wt% of oxygen was successful in lowering the sintering temperature for this M3/2 HSS/MnS composite and confirmed that decarburisation was the major cause of high sintering temperatures in this type of composite. Micro-hardness tests on the martensitic matrix structure for the various composites together with metaUographic evidence that the steel matrix contained low carbon lath type martensite in place of the more acicular high carbon martensite formed in the other composites further supported the idea that decarburisation was responsible for raising the sintering temperature in composites containing the impure grade of MnS The effects of TiC and NbC additions in raising the sintering temperature were probably caused by the dissolution of titanium or niobium into the steel matrix during the period spent in heating up to the sintering temperature, which also raised the solidus temperature.
.I.D. Bolton, A.J. Gant / Joun~al t~fMaterials Processing Technology 56 (1996) 136-147
4.2 Sintering kinetics A valuable insight into the possible mechanism of sintering was also gained by applying the Kingery liquid phase sintering model,[10] to the isothermal data used to determine sintering kinetics. At sintering temperatures below those at which liquid phase was present, e.g. 1225°C, Kingery slopes of well below 1/5th were obtained and supported the proposalof that solid state diffusion was the major cause of densification. Similar shallow Kingery slopes were also obtained at very high sintering temperatures, where densification was so rapid that full densities were achieved even during the stage when samples were still being heated up to their sintering temperature. Thus complete densification had occurred before the effects of isothermal holding on density could be properly measured. Only within an intermediate sintering range, just above the solidus temperature, was the true slope formed by a liquid phase sintering reaction shown. Within this sintering temperature range Kingery slopes of around l/5th were obtained and suggested that the predominant mechanism for densification consisted of the solution re-precipitation between some solid and its surrounding liquid phase. Quantitative measurements of the effects of sintering time/temperature on the relative proportions of MC and M6C carbides together with the metallographic evidence of a gradual transition from M6C into MC carbides at high sintering temperatures prompted the suggestion that this solution re-precipitation reaction could be the dissolution of M6C carbide into the liquid followed by re-preclpitation as an MC carbide. Such a process was prevalent amongst the TiC containing composites due to the natural encouragement given to MC carbide formation by the solution of titanium into the steel, and activation energy determinations for the changes in volume fraction of the MC and M6C carbides relative to sintering time for the M3/2 + TiC composites supported this idea. This activation energy was remarkably close to that determined for densification, see Figure 2, i.e. 1080 kJ/mol. 4.3 Chemical stability of the ceramic/solid lubricant particles Both types of solid lubricant, namely MnS and CaF 2 were shown to be chemically stable when sintered into a high speed steel matrix with little evidence that they dissolved in the steel. In the presence of TiC, however, there were indications that the MnS was thermodynamically unstable and that reaction between MnS and TiC could produce a new Ti4C2S 2 phase, [7]. Although such a reaction was thermodynamically possible at all temperatures the reaction kinetics were such that the Ti4C2S 2 phase was only observed in any significant quantities in samples sintered above 1280oc. Little is yet known concerning the possible effects of this phase on the wear/friction properties of the high speed steel. Both of-the ceramic carbides of TiC and NbC were found to react with the high speed steel matrix but the extent of this reactivity differed between the two types of carbide and was also a function of their particle size. Significant reactions were also only evident at sintering temperatures above those at which liquid phases were present so that it was possible to speculate that any new phases formed were as a result of reaction between the TiC or NbC carbide and the liquid phase formed from the steel matrix. In the case of TiC, reaction with this liquid phase at the surface of the TiC particles was naturally more evident with the finer 5gm grade of TiC particle and caused the formation of an interfacial layer of cubic MC mono-carbide phase which surrounded the TiC particles. Two distinctive varieties, either an iron rich or a vanadium rich form of this MC carbide were produced. The tendencies for titanium, vanadium, or niobium to form an isomorphous series of cubic MC carbides, mutually soluble in one another, is well known, [ 11 ], and would of course explain why MC alloy carbides were formed at the TiC/steel interface. Less clear, however, was the fact that the MC carbides did not form as a complete range of compositions involving a gradual change in their titanium and vanadium composition. Instead only two general compositions for this MC carbide arose with fairly specific metal atom/carbon ratios, i.e., M2X3C 5 (lattice parameter a o ~ 0.427 rlm)or M3X2C5, (lattice parameter a o ~ 0.429 qm); where M denotes the combined total of metallic elements capable of forming cubic MC mono-carbides, e.g., Ti Nb, and V, X denotes the combined total of metallic elements which do not normally form cubic MC carbides, e.g., Fe, Cr,W,Mo, and C is the total carbon, see Table 3. No such interfacial layers were produced in the composites containing NbC and unlike the TiC composites little or none of the original carbide addition was retained after sintering. Niobium carbide particles were totally converted into an alloy MC carbide by the diffusion of alloying elements from the high speed steel into the NbC with no interfacial layer or diffusion gradient from the centre to the outside of each particle. This possibly reflected a greater reactivity for the NbC particle with its surrounding steel matrix, compared with the
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reactivity of TiC, but also demonstrated the role of finer particle size in speeding up any reaction. Progressively more replacement of the niobium by both vanadium and by molybdenum occurred in the carbide as sintering temperatures were increased, but a ratio ofM4XC 5 was maintained between the MC forming and the non MC carbide forming metallic elements. Differences in behaviour between TiC and NbC also existed for the manner in which solution and diffusion of the ceramic carbide addition into the high speed steel affected the microstructure steel matrix. Slight solution of titanium into the steel caused the disappearance of the usual M6C carbides from the steel structure by encouraging MC carbides to form in their place. This did not occur in the NbC materials either because niobium has such limited solubility in the austenite at the sintering temperatures employed,[12] or because its diffusion rates are so much slower than those of titanium in the austenite phase. 4.4 Particle distribution The agglomeration and clustering effects observed with the finer TiC or NbC particles were partially due to the agglomeration of fine carbide particles during the powder mixing stage but this was not the only cause. Dramatic increases in size for the MnS and CaF 2 particles could also be explained by the presence of a liquid phase at the higher sintering temperatures. The smaller ceramic and solid lubricant particles, some of which originated from fragmentation during die compacting stage, were picked up by the liquid phase as it flowed via any interconnecting pores through the compact at the earlier sintering stage. These particles became trapped in the larger pore interstices to become agglomerates which subsequently either coalesced into large particles, as in the case of the MnS and CaF 2 additions, or were formed into clusters of small particles surrounded by a second phase due to their chemical reaction with the liquid phase. Such was the case for the fine grained TiC [5 !am] carbide additions. 4.5 Hardness The addition of hard TiC or NbC ceramic carbides to the high speed steel produced only a small increase in hardness compared with that of the baseline high speed steel both in the as-sintered and in the as-sintered plus double tempered condition, see Table 8, and the double tempering treatment produced higher hardness values than those of the as- sintered materials because of the removal of retained austenite and the development of secondary hardening by carbide precipitation. Significant reductions in hardness were produced by the addition ofMnS/CaF 2 solid lubricant to the various composite mixtures. The small increases in hardness, associated with the addition of TiC/NbC, were presumably due to the small volume fractions of carbide added and somewhat surprisingly the NbC was the slightly more effective of the two carbides, despite the fact that the volume fraction of NbC carbide added to the steel was smaller for the NbC than for the TiC carbide, e.g., 5 wt% ~ 8.3 vol% TiC and 5.4 vol% NbC. Higher hardness values for the NbC composite were also at odds with the observed changes in microstructure, namely that much of the TiC additions were retained in the structure and that TiC encouraged the formation of the harder MC carbides in place of the relatively soft M6C carbides, [4], in the steel matrix: pure NbC particles were not retained and no preferential MC carbide fo.rmation occurred within the steel matrix of the NbC composite. Poor hardness values were achieved in composites containing the solid lubricant and this was partially due to the decarburisation of the steel caused by the introduction of oxygen as an impurity, at least in the case of MnS additions. The more important reason for poor hardness was that both types of solid lubricant tended to collapse under the hardness indentor giving similar effect to that of porosity in under-sintered materials. Thus hardness readings in the composites containing solid lubricants were to some extent not a true value.
5. CONCLUSIONS [1] Metal matrix composites based on high speed steel and containing particulate additions of hard TiC or NbC ceramic plus MnS/CaF 2 solid lubricant particles were successfully sintered to full density. TiC particles additions were essentially preserved as a dispersion in their original chemical form, but the NbC addition was changed into an alloy MC carbide by the diffusion of alloying elements from the high speed steel. This was possibly encouraged by the fine grain size of the NbC used.
J.D. Bolton, A.J. Gant / Journal of Materials Processing Technology 56 (1996) 136-147
[2] Both types of ceramic carbide were chemically bonded into the high speed steel matrix, either as a result of forming an interracial vanadium rich MC carbide layer at the TiC/steel interface, or by diffusion from the steel into the ceramic carbide, as in the case of NbC carbide. [3] Densification was achieved by a liquid phase sintering reaction which appeared to involve solution reprecipitation as the predominant sintering process. Solution of the matrix M6C carbides together with the added TiC/NbC carbides, into a liquid phase, followed by a re-precipitation to form MC carbides was a possible reaction for the solution re-precipitation reaction involved. [4] Although the solid lubricant particles did not react chemically with the steel matrix during sintering they underwent considerable coarsening in size and did not appear to be all that well wetted by the steel matrix. [5] Double tempering at 550°C atter the sintering treatment raised hardness levels but only slight increases in hardness were brought about by the addition of hard ceramic carbide, and considerable, but apparent, reductions in hardness were caused by the addition ofMnS/CaF 2 solid lubricant.
6. REFERENCES [1] Trudel.Y, Cliloglu.C, Tremblay.S, Mod. Dev. in P/M. MPIF Princeton, NY, 15, 1985, 775. [2] Yuko. T, Kazunori U, Lubrication Engineering, 31, 8, (1976), 402. [3] Price.W.J.S, Rebbeck.M.M, Wronski.A.S, Amen.A.S, Powder Metallurgy, 28 (1985), 1, [4] Hoyle.G, "High Speed Steels", Butterworth, London, (1988). [5] BaI1.C.J, Metal Science, 18, (1984), 577. [6] Bandi.W.R, Krapf.G, Analyst, 104, (1979), 812. [7] Lui.W.J, Jonas.J.J, Bouchard.D, Bale.C.W, I.S.I.J Int., 30, (1990), 985. [8] Kubachewski.O, Alcock.C.B, Metallurgical Thermochemistry, 5th edition, 1979, Pergamon, London. [9] Wright.C.S, Powder Metallurgy, 32, (1989), 114. [10] Kingery.W.D, J. Appl. Physics, 30, (1959), 301. [11] Goldschmidt.H.J, J. Iron & Steel Inst. 160, 1948, 345. [12] Narita.K, Trans.I S I J, 15, 1975, 145.
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