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Materials Science and Engineering, A 132 ( 1991 ) L5-L9 L5 Letter On the tempered martensite embrittle- from Fig. 1, both apparent and sharp crack ment in AIS1 4140 low alloy steel toughness (denominated KA and K[c respectively) are significantly degrade
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  Materials Science and Engineering, A 132 ( 1991 ) L5-L9 L5 etter On the tempered martensite embrittle- ment in AIS1 414 low alloy steel F A Darwish Department of Materials Science and Metallurj,% Catholic Universi O' ( PU('/RJ), C.t'. 380(-)8, 22452-Rio de Janeiro, RJ (Brazil) L C Pereira and C Gatts D~Tmrtment of Metallur~.., and Materials Engineering, t:~'deral Universi O' (( OPPE/Ut:RJ), Rio de Janeiro, RJ (Brazil) M L Gra~za Materials l)ivision, 7~'chnical Aero~v)ace ('enter (I>MR/CTA), .~ dO Jos~; dos ('ampos, M' (Brazil) (Received June 11. 1990: in revised form September 10, 1990) Abstract In the present investigation the Auger electron spec- troscopy CAES) technique was used to determine local carbon and phosphorus concentrations on the fracture surfaces of as-quenched and quenched-and-tempered (at 350 °C) AISI 4140 steel specimens austenitized at low and high temperatures. The AES results were rationalized to conclude that, although carbide growth as well as phosphorus segregation are expected to contribute to tempered martensite embrittlement, carbide precipitation on prior austenite grain bounda- ries during tempering is seen to be the microstructural change directly responsible for the occurrence of the referred embrittlement phenomenon. 1 Introduction Tempered martensite embrittlement TME) has long been observed in high strength low alloy steels see for example refs. 1-7). Consistent with an earlier finding by Wood [1], the phenomenon has recently been observed [8, 9] for quenched- and-tempered AISI 4140 steel austenitized at low and high temperatures. As can be noticed from Fig. 1, both apparent and sharp crack toughness denominated K A and K[c respectively) are significantly degraded as the tempering temperature is increased from 200 to 350°C. Fractographic studies [9] have indicated that this toughness degradation is associated with the predominance of an intergranular cracking failure mode along prior austenite grain boundaries of the quenched-and-tempered material Fig. 2). The embrittlement of prior austenite grain boundaries in quenched and low temperature tempered commercially pure high strength low alloy steels has been related by a number of investigators see for example refs. 4-7 and 10-12) to the combined effect of carbide cementite) precipitation as well as impurity- element notably phosphorus) segregation. While phosphorus segregation occurs during austeniti- zation [4, 7, 10], cementite forms by precipitation from the tempered martensite and by thermal decomposition of retained austenite at the P ã 870 °C AUST. --. KI¢ li }e 240 i 1200 °C AUST ....... K A 200 I~ ~_ ill 160 Illl w uJ J#~ x x / i i i O0 f SO 200 350 5 TEMPERING TEMPERATURE ( C) Fig. I. Variation with heat treatment conditions of KLc and K.~ of AISI 4140 steel fracture tested at room temperature [9] (K,, is reported for a notch root radius of 0.25 ram). The data corresponding to the 500°C tempering were not reported in ref. 9 and are shown here to demonstrate the rapid increase in toughness as the steel softens. (Tempering treatments were carried out for 1 h.) 0921-5093/91/ 3.50 © Elsevier Sequoia/Printed in The Netherlands  L6 Fig. 2. Scanning electron micrographs of fracture surfaces for room-temperature-fracture-tested quenched and 350 °C tempered AISI 4140 steel austenitized at (a) 870°C and (b) 1200 °C. boundaries during tempering [2, 6, 7, 10-13]. Subsequent deformation-induced transformation on loading of the remaining intergranular austen- ite [2] adds more embrittling carbides (cementite) to the grain boundaries. The present study concentrated on using Auger electron spectroscopy (AES) to determine local carbon and phosphorus concentrations on the fracture surfaces of in situ broken AISI 4140 steel specimens. The AES measurements, which were made for both as-quenched and quenched- and-tempered material, were then used to gain an insight into the microstructural changes that can lead to TME. 2. Experimental procedure The material used in this investigation was commercially pure AISI 4140 steel, having the following composition (in weight per cent): C 0.38; Mn 0.78; P 0.014; S 0.024; Si 0.29; Cr 0.90; Mo 0.17; Ni 0.25. Pre-notched cylindrical specimens (25 mm long and 2.7 mm in diameter) were machined from the as-received steel plate and were then subjected to different heat treatments to obtain four distinct microstructures, representing, for low (870°C) and high (1200°C) temperature austenitizing, the as-quenched martensite and the tempered martensitic structure produced on tem- pering at 350 °C. These microstructures, termed 870- Q, 1200- Q, 870- 350 and 1200---350 (where Q indicates the steel in the as-quenched condition), reproduce those of the specimens used to obtain the corresponding toughness data reported in Fig. 1. The heat-treated specimens were introduced into the Auger spectrometer where they were fractured at approximately liquid nitrogen tem- perature. Intergranular and transgranular regions on the resulting fracture surfaces were first iden- tified and AES measurements were then made for several spots within these regions using a beam size of 1/~m and maintaining a high vacuum (4×10 -1° Torr) throughout the whole process. A primary beam energy of about 2 keV was chosen and the current density was main- tained at approximately 20 /~A cm -2. Approxi- mate atomic percentages of phosphorus, carbon and iron were determined from measurements of the amplitude of the corresponding Auger peaks in the differential d[N E) E]/dE vs. kinetic energy E spectrum using listed [14] sensitivity factors. No attempt was made to determine the sulfur concentration as it is believed [5] that the presence of manganese in the steel could very effectively precipitate free sulfur as MnS. 3 Results Typical AES spectra are shown in Fig. 3 and the analysis of the AES data are presented in Table 1, where it is observed that, for a given microstructural condition, the carbon concen- tration on intergranular fracture facets is invari- ably higher than that detected for transgranular fracture regions. However, Table 1 indicates no detectable presence of phosphorus on the  5 ._z_  o 3 INTERGRANULARA /~ P V r e ~~ ~.qTRANSGRANULAR Fe o ,~o ~o ~o ~o ~o ~o ~o ~ ,o oo KINETIC NERGY V Fig. 3. Typical Auger spectra from intergranular and trans- granular regions on the fracture surface of AES specimens. These particular spectra were obtained for a sample analyzed in the 1200 ~ 350 microstructural condition. TABLE 1 Analysis of AES data Condition Fracture Amount on fracture surface region (at. ) Fe C P 870 ~ Q Transgranular 83 14 3 870~Q Intergranular 72 25 3 870 ~ 350 Transgranular 80 13 7 870 ~ 350 Intergranular 73 19 8 1200 ~ Q Transgranular 86 14 -- 1200 ~ Q Intergranular 75 25 -- 1200 + 350 Transgranular 82 13 5 1200 ~ 350 Intergranular 63 32 5 fracture surface of 1200~Q AES specimens. Furthermore, the AES data indicate that phos- phorus segregation to both intracrystalline and intergranular fracture regions is favored by low temperature austenitization as well as by the 350 °C tempering treatment. 4 Discussion The results shown in Table 1 indicate a high carbon concentration on the intergranular regions of the fracture surface of the quenched- and-tempered AES specimens. This in turn indi- cates the presence of high carbon constituents along the fracture path, i e along prior austenite grain boundaries in room-temperature-fracture- tested quenched-and-tempered specimens, as L7 intergranular cracking was found to be the pre- dominant failure mode (Fig. 2). High local carbon levels in as-quenched and quenched-and-tempered steels were found [15[to be associated with retained austenite films present in the steel after quenching. In addition to an observed enrichment in the carbon level in the austenite films by a factor of about 2 with respect to the normal level in the steel, a particularly high concentration of carbon (peak values in the range of 10-24 at. ) at the martensite-austenite inter- face was also detected. This very large concentra- tion of carbon at the transformation interface is considered to provide a simple and direct expla- nation for the retardation of any subsequent reac- tion [15]. Accordingly, Sarikaya et al [16] conclude that TME accompanied by quasi- cleavage fracture in alloy steels containing retained austenite after quenching is concurrent with the decomposition on tempering of retained austenite into carbides (M3C) at the lath bound- aries. However, Bhadeshia and Edmonds [3J attribute TME detected in an Fe-V-C steel to the coarsening of interlath cementite resulting from thermal decomposition of interlath retained austenite. Transmission electron microscopy observations [17] made it clear that the carbides nucleate at the austenite-martensite interface and can grow into the austenite and/or martensite. The driving force for this process is the high carbon concentration at regions immediately adjacent to the interface. An increased stability of retained austenite during tempering favors car- bide growth mainly into the martensite or along the austenite-martensite interface. Thus coarsen- ing of the carbide precipitates is not necessarily related to decomposition of the retained austenite [17]. If one considers the ambient-temperature frac- ture behavior of the AISI 4140 steel tested in the 1200-' Q condition some interesting conclusions can be made. In particular, the observation [9] that specimens in this microstructural condition fail by a mixture of quasi-cleavage and inter- granular cracking should indicate that some sort of embrittlement had actually taken place, prob- ably due to deformation-induced carbide precipi- tation in highly carbon enriched regions along the retained austenite-martensite interfaces. The high carbon concentration detected on the frac- ture surface combined with the virtual absence of phosphorus segregation (Table 1) seems to indi- cate that the precipitation of carbides is the  L8 primary cause of the predominance of brittle fracture modes since carbide coarsening is not likely to occur during loading of the as-quenched steel. Based on the discussion presented above, it may be concluded that TME in the steel consid- ered in this investigation is primarily related to carbide precipitation during tempering along prior austenite grain boundaries already weak- ened by phosphorus segregation (Table 1). However, depending on the inherent toughness of the matrix, carbide precipitation alone can trigger brittle failure modes as demonstrated by the frac- ture behavior of 1200--'Q specimens tested at room temperature. The high carbon concentration associated with retained austenite films present at prior austenite grain boundaries favors carbide precipitation and growth along those boundaries during tempering and this is reflected by the AES data (Table 1) indicating considerable carbon segregation on the intergranular cleavage facets of the specimens' fracture surfaces. Carbide growth is expected to take place during the 1 h at 350°C tempering treatment, and this together with phosphorus segregation to prior austenite grain boundaries leads to the predominance of intergranular crack- ing as the failure mechanism in the quenched- and-tempered specimens. Considerable carbide growth along prior austenite grain boundaries has recently been reported [18] for an AISI 4340 sample tempered at 350 °C for 1 h after quench- ing from 1200°C. An elevated austenitizing temperature is expected to result in an increased stability of the austenite in virtue of the excessive grain growth and extensive carbide dissolution in the austenitic matrix during high temperature austenitization. Increased stability of the austenite can result in increases in the amount [19] of etained austenite present at prior austenite grain boundaries as well as in the carbon content. Accordingly, the density and size of carbides formed on those boundaries are expected to increase with increasing austenitizing tempera- ture. This is reflected by the AES data (Table 1) where the carbon concentration on the intergran- ular cleavage facets is seen to be higher for the 1200 ~ 350 specimens than for those analyzed in the 870--'350 microstructural condition. These conclusions seem to be in agreement with the toughness data shown in Fig. 1 (see also Table 2) where it can be noticed that the toughness degradation accompanying the 350 °C tempering is more pronounced for the 1200 °C austenitizing treatment. In fact, the 1200 °C austenitization is so inducive to embrittlement in AISI 4140 steel that it can provoke intergranular cracking in the as-quenched material despite the absence of detectable phosphorus segregation to prior austenite grain boundaries. For conventional austenitization, on the other hand, a ductile failure mechanism was found to prevail in as- quenched specimens [9]. Deformation-induced carbide precipitation on loading was apparently not sufficient to provoke brittle failure modes in room-temperature-fracture-tested 870--' Q speci- mens despite phosphorus segregation concurrent with low temperature austenitizing (Table 1). High temperature austenitization, on the other hand, is thermodynamically and kinetically un- favorable to phosphorus segregation [20]. This is borne out by the AES data which indicate that practically no phosphorus is present on the frac- ture surfaces of the 1200--'Q specimens. The AES results also indicate a considerable enrich- ment in phosphorus at prior austenite grain boundaries, for both austenitizing temperatures, as a result of the 350 °C tempering, in agreement with a finding by Paju and Moiler [21 ]. In addition to phosphorus diffusion enhanced by lattice defects, phosphorus rejection by the growing carbides during tempering [22] can contribute to the segregation of this element to prior austenite grain boundaries. TABLE 2 romparison of toughness values for as-quenched and quenched-and-tempered AISI 4140 steel [9] Condition Ktc Reduction in Kjc K A Reduction in K A (MPa m I/2) due to TME ( ) (MPa m I/z due to TME ( ) 870-Q 57+1 -- 1/3+4 -- 870~350 52+ 1 8.8 102+3 9.7 1200-Q 71 +4 -- 85+2 -- 1200~350 56+2 21.1 68-3 20
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