Carbideprecipitationin austenitic stainlesssteel carburized at lowtemperature
F.Ernst * Y. Cao G.M. Michal A.H. Heuer
Deperiment ef Materialb Sciesce oml Engimering. Case Westerw Rerene U/nisersity. Clereland. OfH 44706-7204 US.A
Received 16 August 2006; received in revised form 14 September 2006; acepted 14 September 2006 Available online 12 January 2007
Abstraet
Low-temperature gas-phase carburization can significantly improve the surface mechanical properties and corrosion resistance ofaustenitic stainless steel by generating a single-phase “case”* with concentrations of interstitially dissolved carbon exceeding the equilib- rium solubility limit by orders of magnitude. Upon prolonged treatment however carbides (mostly x M;C) can precipitate and degradethe propertie. High-resolution and spatially resolved analytical transmission eletron microscopy revealed the precise carbideaustenite orientation relationship a highly coherent interface and that precipitation only occurs when (i the carbon-induced lattice expansion ofthe austenite has reached a level that substantially reduces volume-misfit stres and (i) diffusional transport of nickel chromium andiron enhanced by structural defects can locally reduce the nickel concentration to the solubility limit of nickel in z-carbide.
316 the case is ≥25 μm thick and X-ray diffractometry0.12 [4 6]. Recent scanning Auger microprobe (SAM) mea- and X-ray photoelectron spectrometry indicated x|0| surements based on calibration standards even suggestexceeds the equilibrium solubility limit of carbon by a fac- x[0] = 0.14 (Avishai et al. unpublished data). Thistor of ≥700 at the processing temperature and a factor 10² at room temperature. Such a “colossal” supersatura-tion with carbon is possible because Z is high enough forinterstitially dissolved carbon atoms to retain considerable mobility but low enough to limit the mobility of the metalinto the alloy to form a single-phase carbon-rich “case". in metal carbides cannot readily occur [13]. Another atoms to a level at which the usual precipitation of carbonimportant requirement we have identified for obtaining acolossal supersaturation with carbon is the presence of suf- ficient atom fractions of elements with a high affinity forcarbon. In the 316L-type alloys the main element playing this role is Cr (chromium).
In recent work [16] we reported that the surfacemechanical properties and corrosion resistance of 316-typeaustenitic stainless steel (FeCrNi alloy) can be signifi-lok Company. This process after activating the alloy sur- gas-phase carburization process developed by the Swage-face by removing the passivating chromium oxide scalesphere at a processing temperature T = 748 K for a dura- [7] supplies carbon from a conventional carburizing atmo-tion t between 26 and 38 h. As a result carbon difusesWith increasing depth : into the case the concentrationX[] smoothly decreases from its maximum X[0] at the surface to the intrinsic carbon level X? of the non-carbu-rized alloy core. For low-temperature-carburized bulk
As long as carbon stays in solid solution a colossalsupersaturation of the austenite greatly improves the sur- face hardness resistance to fatigue crack nucleation wearresistance and corrosion resistance [1 2 4 6 8]. Precipitation
of carbon-rich second phases in contrast generallycarburization stage of the Swagelok process is designed degrades the properties. Therefore the duration tp of theto maximize the uptake of carbon while entirely retaining it in solid solution. However on prolonged carburizationor multiple application of the Swagelok process whichmay be of considerable interest for increasing the case depth metal carbides can eventually precipitate. They formin a zone below the surface in which the average carbon fraction exceeds the non-equilibrium solubility limit X(≥0.12 for bulk 316L). Carbon uptake in excess of this limit correspondingly increases the volume fraction of thecarbide phase and the depth of the carbide-containing zonebelow the surface.
Previous work [5 6] revealed two different carbide mod-the metal atoms ie. Fe or alloying elements substituting ifications: MC (or ≥ phase) and MC ("M" stands forFe in its lattice sites in appropriate proportions). The(MC). Laterally the particles are not distributed uni- majority of the carbide particles by far are phaseformly within the carbide zone. Rather they occur in groups or "“colonies"” with virtually carbide-free austeniteregions between them. Fig. la presents a typical conven-tional transmission electron microscopy (TEM) bright-field image of a colony [6]. The particles appear with largeaspect ratios typically featuring a short dimension of ≥20 nm and a long dimension up to several micrometers.They contain a high density of planar faults orientedroughly orthogonal to the long axis.
The x particles have the crystal structure of Fe;Cters [5 6] Table 2 piles the lattice parameters of we (*"Hagg carbide') [911] with very similar lattice parame-have determined by X-ray diffractometry (XRD) fromsoftware packages Celref [12 13] and RIETAN [14 15]. low-temperature-carburized 316L powder [6] using theagree very well with those found for Fe;C in two other The lattice parameters obtained from our measurementspublications [16 17] also listed in Table 2. The space groupis C2/c (15). Table 3 lists the atom positions as determined by Retief [16]. The unit cell contains 20 Fe and 8 carbonatoms consistent with the stoichiometry M;C. Fig. 2 pro- jects the structure of FesC in [010] the direction of thesymmetry axis.
Fig. 1. Conventional TEM of z (M;C needles that precipitated in a (austenite) matrix with a colossal supersaturation of carbon (X [0] = 0.12. recorded with a 200 kV Philips CM20 transmission ekeetroncorectly oriented with respet to (a) The patt constitutes a supepo- microscope). (a) Bright-field image. (b) Selected-area difraction pattern jo ued sxeoz[11] pu 7 jo und stxeoz[010] jo uos The white spots superimposed onto the pattem in the lower rightquadrant were obtained by simulating the [010] diffraction pattern of Fe C2.
tation relationship (OR) with the matrix [5]. Based on The x particles grow in a unique crystallographic orien-Fig. Ib originating from a region including several carbide particles and about 20 similar TEM selected-area diffrac-tion (SAD) patterns we have described this OR as [5]
planar faults apparent in the image and identify the fault ing [100]-oriented rows originate from the high density ofplane as (200). As all the carbide particles constituting the colony in Fig. 1a have the same OR with the matrix ed e s red se m se sxe u pueto each other a unique correspondence exists between the (idealized) particle shape and the orientation of the crystallattice therein. According to Fig. la the typical particle extension in [001] is much smaller than in [100] . TEM
(1)
(2)
The direction (2) corresponds to the viewing direction ofan angle of 90° as do [211] and the normal of (111) - Fig. 1. Note that [010] and the normal of (001) make
with respect to Fig. Ia. The streaks between the spots form- The diffraction pattern of Fig. Ib is correctly oriented
Fig. 2. The crystal structure f Fe; C the Hagg” carbide shown in [010]projection. The structure belongs to the monclinic crystal system and has thespace group is C2/c (15). The shaded region corrsponds to one unit cell.
images in [001] [5 6] show that the extension in [010] isthan in [100]. Accordingly. the particle shape resembles parable to that in [oo1l i.e. it is also much smallerthat of needles or laths but not plates or discs.
As carbides generally degrade the surface mechanicalproperties and corrosion resistance it is important to iden-mental understanding of the micromechanism by which tify the parameters that control X? and to gain a funda-the carbide particles nucleate and grow in order to develop strategies for avoiding them. In the work reportedhere we have studied the atomistic structure of the tenite matrix) and the spatial redistribution of atom species interface (the interface between M C particles and the aus-associated with carbide formation under the constraints of low-temperature carburization by advanced methods ofTEM.
2. Experimental methods and procedures
The material we investigated was a low-temperature-car-burized foil of 316L-type austenitic stainless steel (ADT15)with a thickness of 38 μm and annealed at 1338 K for 1.5 h. Table 1 lists the results of a wet-chemical analysis. Carbideprocess corresponding to a total carburization time formation occurred after two applications of the Swagelokfp = 44 h at 7 =748 K.
carburization with a Scintag X-1 X-ray diffractometer uti- X-ray diffractograms were obtained before and afterlizing Cu Kα radiation with the wavelength = 0.154056 nm in the BraggBrentano (020) mode [18]From the diffractograms we determined the peak centerssponding to / = 1 2 3 4 respectively). From these data 0 of the reflections {1111 {200) {220) {311} (corre-
Composition of the AT15 foil (at%) Table 1
C Mn Si P S Cr Ni Mo AI Cu0.16 1.61 0.98 0.05 100 18.76 10.65 1.10 0.210.35N 0.24 0 0.07 0.07 0.190.05 Co B Nb 0.02 Ca 0.140.07 65.29 Se Fe
we obtained best estimates a and β for the austenite latticeparameter and the NelsonRiley coeficient respectively this approach a and β are given by the values of a andthe measured 0 and their theoretical values β that minimize the sum of the square residuals between
where p := of the planes reflecting under 0 The relation (3) follows di- ² and hare the Miler indicesrectly from the Bragg equation [18] and the NelsonRileycorrection [20 18]. For evaluating [p] (an implicit func- tion) and for fitting the parameters a and β we developeda C implementation of the downhill simplex algorithm [21]. To determine standard error limits for the best esti-mate a we generated 10* new data sets {0′) drawing ran-dom values from normal distributions centered on the theoretical [p] obtained from (3) with the best estimatesa and β. The standard deviation of these distributions was chosen to [19]
where P = 4 is the number of peaks and f = 2 the numberof fitting parameters (or degrees of freedom). Fitting (3) to each new data set {0) we obtained histograms of the 104best values af and β’ for the fitting parameters a and β respectively.Sorting these data yielded corresponding cumulative distributions from which we eventually deter-the 68% level of confidence. The symmetric *standard er- mined (asymmetric) confidence intervals for a and β atrors" indicated for the data in the following sectionscorrespond to the largest possible deviation within the asymmetric 68% confidence interval. A corresponding pro-cedure was applied to difusion data from the literature to obtain pre-factors and activation energies.
The particles and the interface were studied byhigh-resolution TEM (HRTEM) scanning TEM (STEM) with a high-angle annular dark-field (HAADF) detector
(3)
(4)
electron-spectroscopic imaging (ESI) and X-ray energy dispersive spectroscopy (XEDS）. The instrument weemployed was a Tecnai F30 S-TWIN (FEI) operating with e dp oe ss tion resolution limit of 0.14 nm.The instrument isequipped with a HAADF detector (Model 3000 Fischione) d 2001" Gatan retrofitted with a 2k x 2k CCD camera) Li detector. Employing these attachments we performed and an X-ray energy-dispersive spectrometer with an Sielemental mapping using the “three-window method" [22] and XEDS line scans. The latter were conducted in STEMmode based on Z-contrast imaging with the aid of theHAADF detector. In these experiments the full width at half intensity maximum of the electron probe wassize # 6 in nanoprobe mode. Specimen drift was pen- ≥0.5 nm corresponding to gun lens setting # 7 and spotsated by the instrument's built-in capability of monitoringthe position of a characteristic image feature.
Fig. 3. NelsonRiley plot of XRD (X-ray diffractometry) data obtainedfrom the 316L-type austenitic stainless steel foil before and after low- temperature carburization. (After [6])
(8)
The increased scatter about the least-square-fit line ob-served for the points representing the carburized material in Fig. 3 does not reflect statistical error but is a systematicand reproducible effect [6] which we attribute to peakshifts introduced by an increased density of stacking faults [3133]. The peak shifts can be modeled [34 35] and thea third fitting parameter in the non-linear regression anal- stacking fault density can be quantified by including it asysis (Ernst unpublished data). However the value for athe result (6) obtained by treating the peak shifts as a sto- obtained in that way is not significantly different fromchastic effect.
Cross-sectional specimens for TEM were preparedfrom the low-temperature-carburized steel foils following the method developed by Strecker et al. [23] using spe-cially prepared specimen holders.Afer mechanicallyreducing the thickness to 80 μm and further reducing the thickness in the center to ≥20 μm by means of a dim-ple grinder (Gatan Inc.) the final thinning to electronprecision ion polishing system (PIPS Gatan Inc.) fromboth sides until perforation occurred. Prior to loading the specimen into the microscope we cleaned the surfacefrom adsorbants and contamination by a means of an Ar* plasma cleaner (Fischione Model 1020). For simulat-for generating ball models of the crystal structures of and we employed the JAVA version of the EMS soft-[9]d r
3.2. Concentional and high-resolution TEM
Fig. 4 presents another conventional TEM bright-fieldthe hole of the TEM specimen. The carbide particles one image of a colony. This image includes the thin edge atof them marked " appear with uniform gray levels indi- cating that they have the same crystallographic orientation.While TEM images recorded in certain crystallographicorthogonal to the long axis of the particles (e.g. Fig. la) directions reveal a high density of planar faults on planesthese defects do not produce much contrast under the imaging conditions of Fig. 4. The regions of carbon-super-saturated matrix (marked *") also have the same crystal-lographic orientation (average gray level) implying that the matrix regions belong to a single grain and the 7ZDifferent from the carbide particles the matrix regions orientation relationship is constant across the entire image.exhibit distinct gray level variations on the subnanometerlength scale indicating a high density of dislocations. While part of these dislocations will have resulted fromfrom cold work (foil rolling) prior to carburization a sig- nificant fraction may have formed by yielding under biaxialpressive stress introduced by the (depth-dependent) lat-tice expansion acmodating the interstitially dissolved carbon [2].
3. Results
3.1. X-ray diffractometry
difractograms of the 316L-type austenitic stainless steel Fig. 3 shows a NelsonRiley plot obtained from X-rayfoil recorded before and after carburization [6]. The best estimates for the lattice parameter a° of the non-carbu-rized austenite and the lattice parameter a*[0] at the sur-face of the low-temperature-carburized material are
(5)
(6)
This corresponds to an expansion factor of 1.024 ± 0.005.Based on the empirical relationship [2730]
(7)
and α = 0.104 nm [30] the value (6) for a[0] corresponds [9] 01
Fig. 4. Conventional TEM bright-field image of x (M C) particles embedded in / (austenite) with a colossal supersaturation of carbonedge and the hole of the TEM specimen. (X[0] 0.084). The upper part of the image features the amorphous
Fig. 5a is an HRTEM image of the interface between a particle (top) and the 7 matrix (bottom). The interfaceresides in the lower third of the image and lies approxi- mately horizontal and parallel to the viewing direction（"edge-on). It can be identified with the transition between the contrast patterns in the upper and lower halfof the image. Within the region included in Fig. 5a theinterface does not exhibit considerable roughness.
Fig. 5b shows the interface at higher magnification. Thecontrast pattern observed inin the lower half featurese te s e pss jo ss om of 71°. Accordingly. the fringes correspond to [111}it exhibits only one set of lattice fringes the viewing direc- and the viewing direction corresponds to [110]. In 7 astion does not coincide with a low-indexed zone axis. Differ- ent from what might have been expected based on (1) and(2) the lattice fringes parallel to the one set of (111}fringes do not originate from (002) planes. This bees obvious by considering the misfit between the correspond-ing plane spacings. According to the lattice parameters ofa = (0.368 ± 0.002） nm of the expanded austenite dis-(0.212 ± 0.001) nm. Taking into account that these planes make an angle of Φ ≥ 20° with the plane of the interface the misfit
Fig. 5. Atomistic structure of the x interface. (a) High-resolution TEMmatrix (bottom). The viewing direction corresponds to [110] in 7 and image of an interface between a needle (top) and the supersaturated 7specimen close to the hole (upper right). The lattice fringes that [124| in x. This image was recorded in an ultra-thin region of the TEM q o pa 1eue ue oms aoq oq o qoe edde coherence of the field-emission electron gun employed for imaging. (b)Enlarged region of the interface in (a).
*Metollorgical ansd Materials Deparnsmt; COPPE; Fedleral Uitersity ef Rio de Janeiro Brazi Mechanical Department - Federal Uwitersity af Ceari Brazil“Plysical Institate - Federal Unfeersity of Rio de Janeiro Brazi?
Abstract
solution treatment water quenching and magnetic aging at two different temperatures. Coercive force and maximum Magnetic properties of the Fe20Mo5NixC (wt% x varying from 0.016 to 0.14) aloys were studied after hot rolling.energetic product as high as 450 Oe and 1.8 MGOe respectively were obtained. C 2001 Elsevier Science B.V. Allrightsreserved.
Keywonds: Magnetic materials; Magnets; Cobalt free magnets
1. Introduction
aging treatmenl on the magnetic properties of The aim of this work is to determine the influence ofFe20Mo-5Ni alloys with carbon additions ranging from 0.016% to 0.14% (w1%) and aging time varyingfrom 0 to 1080min.
alloys show high mechanicalplasticity when cold Since the 1980s it has been known that FeMoNiVicalloy L deformed as well as magnetic properties similar to
with coercive force H = 210 Oe and maximum energetic Tiefel et al. [1] obtained a Fe20Mo5Ni (wt%) alloyproduct (BH) 1.1 MGOe after hot rolling solutiontreatment at 1200°C and magnetic aging at 610°C for 4 h. With aging temperature of 650°C the resuls were similar[1].
2. Materials and methods
melting furnace using high-purity materials. The chem- The ingots were prepared in a vacuum induction-ical position of the alloys is given in Table 1. The ingots were soaked at 1250°C hot rolled (60% reduc-tion) water quenched and magnetically aged at two different temperatures 610°C and 650°C. The magneticproperties were determined in a vibrating sample mag- netometer. A maximum field of 10000 Oe was appliedduring measuremenL.
Recently Teodosio et al. [2] studied the influence ofon the magnetic properties of the Fe20Mo5Ni alloy. small contents of carbon from 0.016% to 0.14% (wt%) The alloys containing carbon achieved H = 262267 Oe and (BH)mx = 0.830.97 MGOe after hot rolling solu-the necessty of any aging heat treatments or cold defor- tion treatment at 1250°C and water quenching withoutmation.
3. Results and discussion
with the aging time for the temperatures of 610°C and Figs. 1 and 2 present the variation of coercive force650°C respectively. It can be seen that there is a very similar behavior for all the alloys with different carboncontent. The coercive force starts from about 260 Oe
Fig. 1. Coercive foroe H against aging time after aging at610°C.
pides soseao n ug pue “(uean Sue ou qm)until 4 h of aging time for 610°C and 1 h for 650°C and stabilizes in a plateau around 450 Oe for both temper-atures. The alloys that present the highest coercive forceare in the region of 0.046% and 0.075% (wt%) of carbon content. There is a great similarity between the behaviorof H for 610°C and 650°C of aging temperature. The only difference between the two plots is that at 650°C theincrease of Hf is even faster than the increase observed a1 610°C because of the kinetic effect of the temperature.
time for the alloys with different carbon content for an Fig 3 shows B saturation induction against aginga great symmetry with the equivalent H plot (Fig. 1). For aging temperature of 610°C. From this figure one can seeall alloys B stats around 17 500 G with no aging treal-ment) drops very fast until 4h of aging time and stabil- izes at a plateau around 14 000 G. The B behavior wasalso investigated at 650°C; the values obtained were similar to those obtained at 610°C. The only differencebetween the two temperatures is that the drop of B a1 650°C is faster than the drop observed at 610°C becauseof the kinetic effect of the temperature.
According to Becker et al. [3] the microstructure ofthe material has no influence on B . On the other hand
Fig. 2. Coercive fore H against aging time after aging at650°C.
Fig.3. Maximm eergetic produet (BH) against agingtime after aging at 610°C.
this property is strongly influenoed by the chemical - position of the alloy. The precipitation of carbides cha-nges the x-matrix position. But the microstructure ofthe material has a strong influence on the coercive force.
Tiefel et al. [1] and Magat et al. [4] proposed thatFe Mo precipitates are responsible for the high H values. More recently Teodosio et al. [5] showed thepresence of M C carbide precipitates in the alloys con- taining carbon where M can be Fe or Mo. These carbideshancing H values. The precipitation of carbides during are ferromagnetic and hinder domain wall motion en-magnetic aging decreases the quantity of Mo dissolved increases. the z-matrix and consequently the saturation B de-
alloys after magnetic aging at 610°C. This value of max- Fig 4 shows the maximum energetic product of theimum energetic product is higher than that obtained for the alloy without carbon (1.1 MGOe) [1]. The time ofheat treatment (4 h) for the highest value of the maximum energetic product is the same that gives the highest valueof H. The temperature of 650°C was also investigated
1. After 4h of aging al 610°C and 1h at 650°C the coercive forces of the alloys reached the maximumvalues from 400 to 450 Oe and maximum energeticproduct reached the maximum values from 1.3 to 1.8 MGOe These values of H and (BH)x are aboutcarbon and aged at the same temperatures. twice as high as those of the same alloy without2.The highest magnetic properties were obtained for the alloys with carbon contents between 0.046 and3. The perfect symmetry between H and B together 0.075% (wt%).with the presence of precipitates observed in previous works shows a strong indication that the mechanismof magnetic hardening is developed by the formationof precipitates in the x-matrix during low-temperature aging.
Fig. 4. Saturation B agains aging time after aging at610°C.
and the maximum energetic product was achieved after1 h of aging time. The values obtained for (BHf) were similar to those obtained at 610°C.
References
[2] B. Adolfo J.R. Teodosio Joel M. Neto Julio Scr. Mater. 42 [1] S. Jin T.H. Tiefel J. Appl. Phyx. 52 (3) (1981) .[3] JJ. Becker F.E. Luborsky D.L. Martin IEEE Trans. (2000) 711.[4] L M. Magat G.M. Makarova L.P. Lapina E. Belozerov Magn. MAG-4 (1968) 8498.[5] S.S.M. Tavares J.R. Teodosio M.M. Pires Scr. Metall. Fiz. Met. Metalloved. T. 5 (6) (1981) .Mater. 33 (2) (1995) 251257.
4. Conclusions
Fe20Mo-5NixC(w1%) magnetic allys x varyingfrom 0.016 to 0.14 were hot rolled solution treated a1 1250°C and magnetic aged at 610°C and 650°C. Thefollowing conclusions can be made:
Design and experimental study of self-acting gas bearings for cryogenic turboexpanders
Liu Guo-jun²，Wu Qi-meng.Zhang Feng-hua² Xu You-liang²
Hou Yu'，Xiong Lian-you' Wang Jin'. Chen Chun-zheng'. Tang Run-quan*
1. ne f Crmic EnginringXen Jig Uiy an 709 :2.S Osn P Co. ad S2151)
Abstract: The ne serdnamie fol jouml berings with elastic support and spial guoved thst bearin hse beenapphed to a turtboexpander associated with 150me /h oxygen plant They can support the rotlor revolving al 12. 16 ×10'rpm whirh overspeeds by 14% . By using these gas bearings. smooth operatioe is cbtained. Is additioo the pefomance of spiral groved thrust bearings has been studied. The esperimental results bave proved that the applicationof this kind of self-acting gas bearings to the cryogenic turboexpanders is fesible.
Keywords: Gas bearing: Experimental study: self-acting gas bearing
（1]模子，全动压气体轴承低醒透平影账机的研究[博士学位论 文西安：西安交通大学、1999[2.照联发，新型箔片动压气体轴承的研究及其在低温造平攀胺机 上的应用[博士学位论文】，西安：西安交通大学，198“3]刘立强，陈纯正等高通低温机械螺驼槽气体止推轴承的发 展、操冷技卡，1996（11;8-13[4] W X.Ka. C.Z. Chen. The chaerirnnr audy of ga-lbneeed elf-Cryugreics 30 Sept. 1990 ating spml rved thnat bering equippe on epense surlme.S]Y Hou L. Y. Xoeg ad C. Z. Chn et al. The eqperiernal stuly of usng pleefoil jomal bnnnp for lagh spend crogrmcturboesperder.Proemdinof SemhIaematenlCrognic Engneeing Coeferesct Boumeuth U. K.、113 117 1998. 7
Btitste of Polmer Techology n Mterials Engeering (IPTME) Lohrogh Umierity Loghorogh Leicesterhre LE11 3TU UK
plant stees using a modified modelling technique. MTDATA is used to calceulate some of the thermodymamic parameters for the This paper reports a systemati study on the effets of different alloying lments on the precipitation of MCg particles in poweralloying system. The elements studied include tungsten aluminium silicon manganese cobalt copper nickel and molybdenumprecipitates at low concentrations. Coarsening rate decreases progressively with increasing concentration of Si Co Al and Cu. Ni with varying degrees of concentration. Results show that all the elements are beneficial in reducing the coarsening rate of MCcan be very harmful by increasing coarsening rate when its concentration is greater than about 0.4 at.%. Both Mn and W are quite effective in reducing coarsening rate when the concentrations are low. Extremely sharp increases in coarsening rate are found whenthe concentrations of these two elements exced about 0.5 at.%. Mo is significant in reducing the coarsening rate even at very low concentrations and the effect is more pronounced with increasing concentration. All these effects result from the change of solubilityof chromium in the matrix caused by the alloying elements.
material. Therefore numerous experimental studieshave been carried out on the effects of different alloying elements on the microstructure and especially onproperties of steels [2.3]. The difficulty of this kind ofexperimental work on such a plex system is obvious. In recent years puter modelling and/or simulationof microstructure evolution and mechanical properties has bee more and more important. Compared withexperimental work puter modelling does not havemodelling can give insights of structure and property the difficulties of being lengthy or costly. Computerevolution and therefore is able to predict long timebehaviour of the material [4 5]. Examples of such modelling include the model developed by Bhadeshia’sgroup [6.7] and the model used in the DICTRA software [8.9]. Carolan and Faulkner [10] and more recently Jiang and Faulkner [11] proposed a model for inter-granular precipitation that is based on the concept of collector plate. It is unlike the above models that areconcentrated on intra-granular precipitation. This model has been reasonably successful when applied toaluminium alloys and austenitic steels [12 13]. However
In recent years there is increasing demand forincreasing power plant efficiency to reduce both theThis promotes the development of new steels used insupercritical power plant [1]. It is well known that microstructural evolution of the steels determines thelife and long time behaviour of the material. One of theprecipitate population. The coarsening rate of these main features of the microstructure of steels is its.mechanical properties of the steels. The stability of these second phase particles determines the degradation of theprecipitates depends on mainly the chemical posi-tion and heat treatment of the material.
It is of vital importance in the development of newalloys to understand the effects of different alloying elements on the microstructure and properties of the
there are no systematic studies of the effects of differentsystem and the lack of thermodynamic data available. alloying elements in steels due to the plexity of theMTDATA [14] developed by the National PhysicalLaboratory (NPL) provides a useful tool to calculate thermodynamic paramcters.
main alloying elements on the precipitation and coar- This paper reports a systematic study on the effect ofpower plant steels using puter modelling based on sening of inter- and intra-granular MzsC precipitates in[10 11] and MTDATA.
2. Modelling details
2.1. Inter-gramular precipitation model
Details of the model for inter-granular precipitationcan be found elsewhere [10 1l]. The model assumes that in Fig. 1. Basically the modelling includes four stages of the inter-granular precipitates are cap-shaped as shownacross the matrix is determined ie. segregation of solute calculation. (1) The solute concentration distributionatoms to grain boundaries during quenching (2) Calcu-lation of nucleation site density at the agcing tempera- ture at grain boundaries according to the soluteconcentration calculated during stage one. (3) The nuclei formed in stage two are allowed to grow at the ageingtemperature due to solute diffusion from grain centre tocentre to grain boundaries coarsening occurs where the boundary. (4) When all solute has diffused fromlarge precipitates grow while smaller ones dissolve although only the average size of particles is consideredin the actual modelling.
2.1.1. Non-equiibriom and equilibrium segregotion
causes solute re-distribution in the material. It is thought Segregation of solute atoms to grain boundariesthat at ageing temperatures higher than 0.55 the meltingmust be considered while equilibrium segregation (ES) point of the alloy non-equilibrium segregation (NES)also is the main consideration at lower temperatures. The solute distribution across the matrix is a pli-
Fig. 1. The grain boundary precipitates formed by two sections of thesame sphere of radius r. GB stands for grain boundary and 9P is the contact angke.
cated function of the distance from the grain boundary.profile are as in Ref. [1l]. It is well known that the The expressions for calculating solute concentrationdriving force for nucleation is related to the super- saturation of solute atoms. Therefore this segregationprofile provides the basis for the subsequent nucleationdensity calculation.
2.1.2. Nucleation site density
grain boundary is given by The number of nuclei formed per unit area of the
phase G* is the Gibbs free energy of critical nucleus where x is the solute concentration in the precipitateformation N is the number of atom sites at the grainboundary k is the Boltzmann constant and 7u is the nucleation temperature. G* depends on the free energyinterfacial energy s at the interface between the change per unit volume of nucleus AGv and thenucleus and the matrix in accordance with the followingequation:
where y is the angle between the surface of nucleus andthe grain boundary (see Fig. I). This defines the shape of the nucleus. AGv is the driving force for the phasetransformation and can be estimated from
volume of nucleus C the solute concentration at grain where R is the universal gas constant V. the molarboundary and Ci is the equilibrium solute concen- tration at nucleation temperature. Nucleation is impor-tant in the following precipitation calculation. Thegrowth rate of precipitates depends on the average distance between the nuclei which is given by /I/N
2.1.3. Growth of precipitates
In the constant square collector plate model thesolute atom arrival rate at the collector plate determinesis the square with side length equal to the distance the growth rate of a precipitate. The collector plate areabetween adjacent precipitates.Therefore the initial collector plate area would be
N is defined in Eq. (1). Assuming that the ageing time is divided into a series of small time intervals &r the sizeof the precipitate at time rr is L1 and that at time
(1)
(2)
(c)
(4)
it is L then
2.2.1. Coarsering orset tine
determined using Eq. (6). This is based on the assump- In the previous model coarsening onset time istion that coarsening occurs when all solute atoms haveMC in ferritic steels this assumption is no longer been dragged from grain centre to the boundaries. Forvalid. Typical carbon content in 9 and 12 wt.% Cr ferritic steels is about 0.15 wt.%. Therefore only a smallfraction of chromium is precipitated in MsCg due to thefraction of M23Cs is determined by the carbon content very low carbon content. Thus the maximum volumefor the solute atoms supply from intra-granular pre- rather than chromium content. Also there is petitioncipitates.
(5)
4m1 is the collector plate area at 1&r D is thevolume diffusion coefficient at ageing temperature Cis the average solute concentration in the region from the grain boundary to a distance /2D&r into the grain C is the solute concentration in the precipitate phase temperature and p. and p. are molar density of thematrix and precipitate phase respectively. f() = 2π/sin²(2/3)cos (1/3)cos² ) defines the rela-tionship between the precipitate volume and the cube ofthe size parameter L. Thus given the initial collector plate area and precipitate size (the critical nucleus size atnucleation) precipitate size L at any time 1 afternucleation can be evaluated using Eq. (5). Nucleation times are very small of the order of milliseconds.
Generally when either all non-rate controlling atomsor all rate controlling atoms have been precipitated thevolume fraction of the precipitates would stop increas- ing with time. Thus the maximum possible volumefrom one of the following cquations: fraction of the precipitates Vmax can be estimated
(8)
(6)
2.1.4. Coarsening
would occur when all solute atoms have diffused from In the original model it was thought that coarseninggrain centre to the boundary. Therefore the coarseningonset time is determined by
In Eqs. (8) and (9) Ng and Nc are the number of ratethe formula MC of the precipitate C and C are controlling and non-rate controlling atoms as defined inthe contents of non-rate controlling and rate controllingfraetion of precipitates depends on which of the two elements respectively. The actual maximum volumevalues determined by Eqs. (8) and (9) is smaller. In the case of MCs in ferritic steels V is more than onesening onset time is determined by carbon contents.
(9)
where d is the grain size. During coarsening the size ofprecipitates as a function of time is given by [15]
(7)
When intra-granular precipitation occurs the intra-granular precipitates will also take solute atoms. There- fore Eq. (6) does not apply in such cases. However ifwe take the volume of both inter- and intra-granularvolume fraction either Eq. (8) or Eq. (9) still can be precipitates into account in calculating the precipitateapplied. The maximum volume fraction of MCg in the steels studied here calculated using Eq. (8) is in goodagreement with the equilibrium value calculated usingMTDATA. This to some extent justifies Eqs. (8) and (9). Therefore Eq. (8) or Eq. (9) is used instead of Eq. (6) tocalculate the coarsening onset time.
d is the grain boundary width D is the grain boundarydiffusion coefficient Lg is the precipitate size at the onset of coarsening Ic is the coarsening onset time asdefined in Eq. (6) 7 is the ageing temperature 4 is aconstant depending on the fraction of grain boundary constant determined by surface energies and B is also a
2.2. Modifications made to the model
In order to use Eq. (8) or Eq. (9) to determine thecoarsening onset time the total volume fraction of bothinter- and intra-granular precipitates is calculated after each time interval r from the precipitate size and thevolume fraction is pared with the maximum values number of the precipitates. Then the calculated totalas shown in Eqs. (8) and (9). When the calculated
The above model has been reasonably successful inpredicting grain boundary segregation and precipitation in aluminium alloys and austenitic steels [12 13]. How-ever when applied to ferritic steels there are some limitations in the model. Therefore modifications havebeen made to the model and are introduced below.
The nucleation of intra-granular precipitates is treatedhomogeneous nucleation the free energy can be written in a similar way to inter-granular precipitates. Assumeas
precipitate volume fraction reaches the maximum coarsening takes place.
2.2.2. Effects of tempering on ageing
In the previous model only one ageing temperaturecan be dealt with. Therefore tempering and subsequent ageing are treated separately. The modelling of temper-ing and ageing is exactly the same; i.e. precipitates nucleate and grow from the critical nucleus size at theageing. This means that tempering has no effect on the appropriate temperature both in the tempering andcerned. This is certainly not the case in reality. Thewould continue to grow. Therefore the model hasbeen modified to treat quench tempering and subse- quent ageing in a sequence that is during ageing thetempering. This treatment not only overes the precipitates will transform from the state at the end ofshortings in the previous model as mentioned above but also can be generalised to more plicated heat treatment conditions.
critical nucleus radius and corresponding free energy are where r is the radius of the spherical nucleus. Thethus as follows:
(11)
(z1)
Eq.(3） with solute concentration in the matrix C The free energy change Gv can be calculated fromWe believe that as this is a non-equilibrium situation; instead of solute concentration in grain boundary C.the chemical potentials within the grains leading up tothe grain boundary are never levelled out. Therefore the driving forces for inter- and intra-granular precipitationdiscussed above. The average intra-granular precipitate are different and can be estimated through Eq. (3) assize then is determined using Eq. (10) with initial nucleussize as r * rather than 0:
2.2.3. Interaction between intra- cnd inter-granular precipitation
In the previous model intra-granular precipitation isconsidered separately from inter-granular precipitation. The model used for calculating the growth of intra-granular precipitates is the simple Zener model in which the average particle size of intra-granular precipitates x is described by the following equation [13]:
(13)
The nucleation site density here is determined also usingEq. (l) with N being the number of atom sites per unitvolume of the matrix and using the intra-granular precipitate driving force. Thus it is also possible tocalculate the average particle spacing for intra-granular precipitates.
(10)
concentration in the matrix Ca the equibrium where D is the diffusion coefficient C is the solutethe precipitate phase and the time. A few problems are solute concentration C the solute concentration inassociated with such a model. Firstly no nucleation isconsidered. Therefore no information about the spacing and volume fraction of the precipitates is available.Secondly as intra-granular precipitation is considered separately the interaction between intra- and inter-granular precipitation is not considered. For very longageing time this might not be a big problem because a vast majority of the precipitates will be in the grainboundaries. However in the short term as intra- andthey would definitely affct each other’s growth rate. In inter-granular precipitates pete for solute atoms addition if intra-granular precipitates exist they will take up carbon and solute atoms and therefore shortenthe coarsening onset time. Therefore intra-granular precipitation has been included in the previous modelto give an integrated precipitation model.
When coarsening occurs the size of intra-granularprecipitates changes with time according to the follow- ing equation [16 17]:
(14)
where r is the radius of the precipitate at the beginning of coarsening.
2.3. Applications of the model
on the alloy 5 supplied by CORUS. The chemical Modelling has been carried out on a model steel basedposition of the steel sample is shown in Table 1.The alloy is initially treated as a FeCrC ternary system as the basis for parison. Then the alloy isfollowing main alloying elements: Al Si W Mn Mo treated as FeCrCX system with X being one of theCo Cu and Ni. MTDATA is used to calculate the
Table 1Chemical positions of Corus alloy 5 (Fe balance)
Cr 0.15 C 0.55 W Mo Mn !S P S Co 9N V 0.061 N Cu B A1wt.% at.% 11.4 12.2 0.69 0.17 1.52 0.88 z0 0.40 0.81 0.82 0.27 0.53 0.014 0.025 0.009 0.016 2.24 2.11 0.07 00 0.24 0.26 0 0.48 0.42 0.005 0.026 0.01 0.021
equilibrium solute concentration as a function of temperature and of concentration of different alloyingelements. In the calculation using MTDATA all fourelements in the FeCrCX systems were considered and all possible phases identified by MTDATA wereincluded in the calculation (refer to Figs. 57 for the phases included). The result then is used as an input tocach system at different concentrations of alloying the modelling and model calculations are carried out onelements and at two ageing temperatures 600 and650C for up to 100 000 h. The solution heat treatment q . sdto ambient temperature. The alloy is then tempered at 760 °C for 1 h. These heat treatment conditions are inaccordance with the heat treatment of the CORUS alloy5. The main parameters used in the modelling are shown in Table 2.
steels do not have significant effects on either the [23]. The results show that these major phases in ferriticvolume fraction or the coarsening rate of M3Cs particles. We believe that this is because of the highchromium content of the steels studied in our modellingand simulation. The amount of NbC present in the steel is small due to the low content of Nb. Therefore theand coarsening rate of MCg particles very much. This presence of NbC would not affect the volume fractionto some extent justifies the application of the approachdiscussed above to the steels studied. However attempts are being made to revisit this problem using our morerecently developed Monte Carlo approach where weinto consideration. can take multi-ponent and several second phases
In such a treatment the interactions between differentalloying elements are neglected. The reason is that the plexity of the phase diagrams and thus the difficultyin getting reliable results on the equilibrium concentra- tion increases sharply with increasing the number ofalloying elements. However even with this limitation this study can give insights and guidelines of the effectsof different alloying clements on the precipitation andcoarsening of MsCg in ferritic steels.
3. Results and discussion
3.1. Precipitation cuerres
base alloy (i.e. FeCrC system) aged at 600 °C for up Fig. 2 shows the modelled precipitation curve of theto 100o00 h. The data points are experimental measure-MC precipitates in Corus alloy 5 creep tested (stress: ments of the average size of inter- and intra-granular154 MN m 2) at the same temperature for 6378 h. Forcquivalent circle radius (ECR) i.e. the radius of a circle inter-granular precipitates the size is described by thewith the same area. Although the sample is creep tested and strain will accelerate coarsening we do not expectthere is much difference between the sizes of the particlesin the aged and creep tested samples as the time is not very long. Therefore the prediction in Fig. 2 is in
We also acknowledge that in the stels studied here there are many second phases in addition to M2sC6such as laves phase and vanadiumnitride particles. In our more recent simulations of precipitation in a similarwe considered simultaneous precipitation of MsC6 ferritic steel P92 using a new Monte Carlo approach laves phase and vanadiumnitride particles and theMTDATA by taking all the clements into consideration thermodynamic parameters are calculated using
Table 2Main parameters used in the modelling of precipitation kinetis
Parameter Unit Value ReferenoesPo molar density precipitate Px. molar density matrix molm mol m" 5496 110497 [18] [19]D grain boundary diffusion pre-exponential term Q grain boundary diffusion activation enengy m²s- kJ mol-1 502 283.1 [20]° [20]Dv volume diffusion pre-exponential term Qv volume difusion activation energy kJ mol1 m² s1 8.5 × 104 240 [21] [21] contact angle interfacial energy degree 0.668 57 [10] [22]
* Calculated from triple produet taking = 5 × 1o * m and = 1.
ay-Zau-umNNV(1.ChinaDepartment of Power Engineering North China Electric Power University Beijing 102206 China;2.100080School of Energy Science and Engineering Harbin Institute of Technology Harbin 150001 china;3.
Abstract: The nonlinear question of bearing journal swirl has been researched in this article by the way ofnonlinear luvrication mechanism analysis. In the basic of oil film pressure analysis and the swirl frequencyn n isn uosp nssd o un jo o e pr m equation of oil film force and the bearing coupling motion we give the journal swirl principal in differentand load unstability history area. As the same time the waveform conversion between the convergence oilwedges positive pressure and divergence oil wedge negative pressure has been finished in the three regions. Inthe different off center of the bearing journal we go deep into discuss the alteration of swirl frequency alongwith bearing joumal autorotation rate extrusion speed and load. In addition give the swirl character ofnonlinear dynamic oil film force. So in this article a new way of rotating machine bearing journal swirlfrequency analysis and diagnosis and active control has been put forward.
Key words: sliding bearing: swirl frequency; nonlinear character
Recentprogress:kinetics of thebainite transformation in steels
Manabu Takahashi *
Stel Rsrh Larrres Ni St/ Corton St 20-1 Fats Ch 293-8511. Jap
Received 15 August 2004; received in revised form 19 August 2004; accepted 23 August 2004
Abstraet
s fully accepted experimentally. The nucleation-controlled model based on the displacive and diffusionless nature of the bainite trans adopt the solute drag efeet and a quantitative approach is proposed although the difusion of substitutional alloying elements is notformation bined with the kinetics of cementite formation in austenite and in ferrite as the simultaneous reactions. 2004 Elsevier Ltd. All rights reserved.
ment conditions for a given production line. Even when steel contains substitutional alloying elements such as Siand Al which are known to retard the formation of cementite the untransformed austenite could de-pose to cementite and ferrite during extended holdingat bainite transformation temperatures. It is therefore essential to model the bainite reaction and the cementiteformation individually to permit the prediction of a ‘processing window' where it is possible to obtain themixed microstructure of ferrite bainite and retainedmation are still not well established as Bhadeshia austenite. However the kinetics of the bainite transfor-nite transformation [5]. In this review the author will try pointed out in his article on unresolved issues in the bai-to assess the recent work that is considered to contributetransformation in steels. the quantitative expression of the kinetics of the bainite
of the intensely discussed phenomena in steels. The The bainite transformation is as well recognized onemajor interest of the work is focused not only on themechanism of the bainite transformation itself but also on a quantitative expression of the kinetics of the bainitetransformation. As it is well known the bainite transfor- mation has been widely exploited in industry to improvethe properties of steels. One of the successful applica-tions of the bainite transformation is so called low-al- loyed TRIP steels [14]. An enrichment of carbon intomass% is maintained in the absence of cementite precip- untransformed austenite up to more than about 1itation during the bainite transformation. The untrans-formed austenite then remains stable even at ambient temperature. Since the rate of the bainite reaction de-pends strongly on the chemical position of the austenite at the beginning of the reaction the fulldescription of the kinetics of phase transformations islect the right alloy chemistry and appropriate heat treat- extremely useful from an engineering point of view to se-
2. Maximum degree of the reaction
It is always very important to know the maximumpossible degree of the transformation as a function ofthe reaction temperature. When bainitic ferrite forma- tion is acpanied by the cementite precipitation the
of bainitic ferrite can be casily calculated from the car- transformation is known to plete. And the amountbon content of the steel. It is not however the casewhere cementite precipitation is retarded effectively. When the formation of cementite is separated from thatof bainitic ferrite by the addition of Si or Al a two-stepreaction can be observed. The first step corresponds to the formation of bainitic ferrite and the onset of the sec-ond step is recognized as due to the start of cementite precipitation from carbon-supersaturated austenite.
The maximum degree of the first step is extremelyimportant since it determines the amount of retained austenite in low-alloyed TRIP steels. Two deferentmechanisms which govern the maximum degree of the reaction have been discussed. One is the effect of partic-ular alloying element on the movement of the transfor- mation interface and the other is the diffusionlesstransformation nature of bainite. In principle the recon-structive formation of ferrte procceds until the equilib- rium or para-equilibrium condition is maintained whenthe formation of cementite is absent. In order to explain the existence of the transformation stasis in the recon-structive reaction an additional mechanism such as sol-ute drag effect is always needed to introduce. Wu and his co-workers [*6] tried to explain the transformation stasisquantitatively by introducing the drag effect of Mo on the interface movement. Although they adopted 20 or30kJmol′ of the binding energy of Mo at the centerof the boundary the result showed only a qualitative agreement of the tendency of the degrees of the reactionas a function of the reaction temperature. Further improvement of the model may be nceded to deal quan-titatively with the effect of the various kinds of substitu-reaction. Furthermore there is a discrepancy among tional alloying element on the maximum degree of thethe measurements of substitutional alloying element segregations to the transformation interfaces. Althoughan extended segregation of Mo at the vicinity of aboundary is reported as a result of a STEM-EDX analysis [7] vast numbers of finest scale analyses usingAP-FIM preclude the existence of any segregation of substitutional alloying element [812].
On the other hand the calculation of the maximumpossible degree of the diffusionless ferrite formation is straightforward. Observed carbon concentrations in re-tained austenite are reported to agre well with the cal- culated T curve in different steels [13 14]. The negativeslope of the To curve predicts the carbon concentrationin residual austenite to be greater at lower bainite trans- formation temperatures as is observed experimentally.The effect of alloying elements on the carbon concentra- tion in retained austenite can also be explained as theireffect on the To curves [15]. Most of the alloying ele-ments decrease the carbon concentration at the T curve except the case of Si Al and Co. A1 [16 17] and Co [17]increase the carbon concentration at the T curve and
Si hardly changes it. This tendency quantitativelycorresponds to the effect of those alloying element on the maximum possible carbon concentration in re-tained austenite namely the maximum degree of the reaction.
3. Overall reaction kineties
nucleation and growth mechanism the overall reaction Since the bainite transformation proceeds by therate depends on the driving force and the mobilityAlthough the transformation pletes when the resulting in a C-curve nature of the TTT diagram.cementite formation is allowed to occur it is essential to treat the formations of bainitic ferrite and cementitecope with a large variety of alloy chemistries that are individually and allow them to pete cach other toadopted in industries. There are two major approachesto express the overall kinetics of the bainite transforma- tion one of which is the diffusion controlled growthmodel of the bainitic laths and the other is the nuclea-tion of bainitic ferrite plates controlled model.
3.1. Diffusion controlled growth model
The growth rate measurements of bainite bainiticferrite and sub-unit have been reported for various - binations of alloy chemistry. There seem to be big dis-Trivedi's model is reported to be applicable to predict crepancies among these measured growth rates.the growth rate of the bainite laths [18] although theculated one. It is however worth noting that bainite measured growth rate is somewhat smaller than the cal-consists of fine sub-unit of ferrite whose thickness is measured to be around 0.1 or 0.2 μm at 400°C and de-creases with the transformation temperature and anincreasing the alloy carbon content [19]. This result is consistent with the analysis made by Singh and Bhades-hia [20]. This fine structure of bainite was also repro- duced in a thre-dimensional way [*6 *21] supportingthe autocatalytic nucleation of ferrite. From these inves-tigations the growth rate discussed by Quidort and Bre- chet [18] could be that of sheaves of bainite rather thansub-units. Observed larger growth rates in the low Si steel can be understood in a different manner from theirdiscussion. Because of the existence of cementite precip-itation carbon supersaturation in untransformed auste- nite is reduced and then the total amount of bainitewas analyzed to be larger growth rates. In the case of increases. This results in larger sheaves of bainite thatthe diffusion controlled growth model of bainite laths ost9] it is always necessary to adopt the solute drag effectelement such as Mo is reported vast numbers of finest
scale analyses using AP-FIM support no segregation ofinterfaces [812] substitutional alloying element at the transformation
and the carbon concentration in the untransformedis inevitable within a realistic period of time. In this case austenite. In general however the carbide formationfrom the untransformed austenite and the sub-units of the model should include the formation of carbide bothbainite which petes with the formation of the sub- units. The overall process is illustrated in Fig. 1 [31].Bainite nucleates with the partitioning of carbon butgrows without any partitioning. The nucleation theory [32] of bainite was confirmed to be applicable to highcarbon steels [33] that is important when the carbonformed austenite at ambient temperature is required. enrichment up to an amount to stabilize the untrans-So the sub-unit of bainite is fully carbon supersaturated at the growth stage. Carbon atoms are however evacu-ated to the surrounding austenite after the formation be- cause the reaction temperature is high enough forcarbon to diffuse along a long distance. This carbonredistribution can cause a carbon enrichment of the untransformed austenite. When the carbide forms insidepleted the lower bainite is obtained whereas car- the sub-unit of bainite before the carbon evacuation isbide precipitates from the carbon-supersaturated auste-Azuma et al. recently modeled whole of this process nite resulting in the upper bainite microstructure.the formation of bainitic ferrite was adopted. And successfully [**34]. Rees and Bhadeshia's model forfor the cementite precipitation both in austenite and in ferrite. Avrami's extended volume concept was modifiedThis model can be used to predict the overall transfor- [35 36] to deal with the simultaneous transformations.mation kinetics with and without the cementite precipi-tation. As a result the model can reproduce the two-step reaction at 450°C and appearance of the lower bainite at
formers such as Mo and W are included alloy carbides In particular alloy systems in which strong carbideformation reported to occur at the transformation inter- faces which suggest a long range diffusion of substitu-tional alloying element during or after the formationquence of the reactions will be needed. of bainite [25 26]. Detailed investigations of the se-
3.2. Nucleation controlled model
sionless nature of the formation of sub-units of bainitic Another model is based on the displacive and diffu-ered to be large enough to simplify the overall transfor- ferrite. Since the growth rate of the sub-unit is consid-mation rate being controlled by the successivenucleation of bainitic laths or plates. Quantitative expression of this model is first reported by Bhadeshia[28 29] This approach is applicable to steels where no [27] and latterly modified by Rees and Bhadeshiacarbide precipitation is expected to occur until the car-bon concentration in untransformed austenite reaches the To curve above which the diffusionless transforma-tion is thermodynamically impossible. In this model it is essential to determine the size of sub-unit of bainite[19 20]. A theoretical approach was conducted recentlysub-units is determined as a point where the chemical by Van Dooren et al. [**30] in which the thickness offree energy released by the formation of a sub-unit is equal to the mechanical energy dissipated in the auste-nite due to the growth of a sub-unit. When the forma-tion of carbide is inhibited these models can be applied to predict the degree of the transformation
Fig. 1. Schematic illustration of the overall process of tbe bainite transformation in steels.
Fig. 2. Calculated change in volume fractions of bainitic ferrite cementite (q) and austenite in a Fe0.6% C1.5% Si1.5%Mn alloy transformed at 300 and 450°C. Plots are experimentally observedvolume fraction munication). of bainitic ferrite (M Azuma.personal
350°C in Fe0.6% C1.5% Si1.5%Mn alloy as shown inFig. 2 (M. Azuma personal munication). Effect of Si on the retardation of the cementite precipitation canbe understood as the result of a diminished driving force for the formation of para-equilibrium cementite. In thecase of the bainite transformation in low-alloyed TRIPsteels the sub-unit of bainite nucleates side by side due to a geometrical limitation of small austenite grains[*37 *38]. It is therefore important to take into account the nucleation of sub-units both at the austenite grainboundaries and at the sub-units. Matsuda and Bhades-hia [*39] reported a model that deals with these two dif- ferent nucleation sites in carbide free bainiticapplied to the microstructural development in low-al- microstructures. This approach could be extended andloyed TRIP steels. Further investigations are needed.
4. Effect of pre-deformation and external stress
There is a continued interest in the effect of externalstresses and pre-deformations on the kinetics of the bai- nite transformation [40] since the grain refinement of ashot rolled multiphase steels including low-alloyed TRIPsteels is one of the successful approaches to improve mechanical properties. It was confirmed experimentallythat the plastic deformation of austenite reduces the
5. Conclusions
Acknowledgement
References
total amount of the bainite transformation due tomechanical stabilization of austenite [4144] although external stresses accelerate the bainite transformation[40 41]. The effect of pre-deformation on the kinetics of the bainite transformation is rather plicated.retarded [40]. The retardation of the overall reaction The transformation rate is first accelerated and finallywas again confirmed experimentally [4244]. Quantita-external stresses is needed to develop. tive treatment of the effect of pre-deformations and
Modeling of the kinetics of the bainite transforma-petition of bainitic ferrite and cementite can be pre- tion has e to the stage where the formation anddicted quantitatively. The displacive and difusionlessnature of the bainite transformation is now widely ac- cepted since the theory can predict quantitatively theproducts obtained in actual production lines. However amount and the chemical positions of the reactionfurther developments are required for the model to dealIt is also worth noting that the kinetics of the lower bai- with the effect of pre-deformations and external stresses.nite transformation just above the Ms temperature is an area where more detailed investigations are nceded. Aninteraction between the formations of bainite and mar-of the overall reaction rate at the temperatures [45]. tensite could play an important role in determination
The author is grateful to Prof. H.K.D.H. Bhadeshiasions on the manuscript. and Dr. N. Fujita for their useful ments and discus-
The papers of particular interest have been high-lighted as:** of very special interest. of special interest;
[1] Matsumura O Sakuma Y Takechi H Enhancement ofelongation by retained austenite in intercritical annealed 0.4C1.5Si0.8Mn steel. Trans ISIJ 1987;27:5709.[2] Matsumura O Sakuma Y Takechi H. TRIP and its kinetic aspects in austempered 0.4C-1.5Si-0.8Mn steel. Scripta Metall[3] Sakuma Y Matsumura O Akisue O. Influence of C content 1987;21:13016.and annealing temperature 0 on microstructure and mechanical propertiesof 400 C transformed steel containing retained[4] Kawano O Takahashi M Wakita J Esaka K. Abe H. austenite. ISIJ Int 1991;31:134853.Improvement of ductility by retained austenite in bot rolled
[23] Quidort D Brechet Y. The role of carbon on the kinetics of o ss o q[24] Enomoto M. Partition of carbon and alloying elements during the growth of ferrous bainite. Scripta Mater 2002;47:1459.[25] Shiflet GJ Hackenberg RE. Partitioning and the growth of bainite. Seripta Mater 2002;47:1637.[26] Hackenberg RE Shiflet GJ. A microanalysis study of the bainite reaction at the bay in FeCMo. Acta Mater 2003;51:[27] Bhadeshia HKDH. Overall transformation kinetic. Journal 213147.de physique colloque C4 Supplement au No.12 1982;43:437 41.[28] Rees GI Bhadeshia HKDH. Bainite transformation kinetics Part I modified model. Mater Sci Technol 1992;8:98593.[29] Rees GI Bhadeshia HKDH. Bainite transformation kinetics Part 1 modifed model. Mater Sci Technol 1992;8:9946.[**30] Van Dooren D. De Cooman B.C. Thibaux P. Theoreticalof Int. Conf. on Advanced High Strength Steels for Automotive Applications AIST 2004;24757.[31] Takahashi M Bhadeshia HKDH. Model for transition from upper to lower bainite. Mater Sci Technol 1990;6:592603.[32] Bhadeshia HKDH A rationalization of shear transformation in steels. Acta Metall 1981;29:111730.[33] Garcia-Mateo C Bhadeshia HKDH. Nucleation theory for high-carbon bainite. Mater Sci Eng A 2004;378:28992.xddn u1 un se N mn nzy [] and lower bainite transformation in steels. Mater Sci Forum[35] Robson JD Bhadeshia HKDH. Kinetics of precipitation in 2003;426432:140512.[36] Jones SJ. Bhadeshia HKDH. Kinetis of the simultaneous power plant steels. Calphad 1996;20:44760.deposition of austenite Into several transformation prod- ucts. Acta Mater 1997:.[°37] Girault E Jacques P Ratchev P Van Humbeeck J Verlinden B Aermoudt E. Studly of the temperature dependence of thebainitic transformation rate in a multiphase TRIP-assisted steel. Mater Sci Eng A 1999;273275:4714.[*38] Godet S Harlet Ph Delannay F Jacques PJ. Critical assess- ment of the bainite transformation of finely grained and[*39] Matsuda H Bhadeshia HKDH. Kineties of the bainite trans- deformed austenite. Mater Sei Forum 2003;426432:14338.[40] For example see Bhadeshia HKDH. Bainite in Steels. Lon- formation. Proc R Soc London A 2004;460:170722.[41] Yao K-F Liu C Shi W Liu Z. The infuence of stress and don: Institute of Materials; 1992. 20124.strain on kinetis of bainite transformation. J Shanghai Jiaotong Univ 2000;E-5-1:2259.[42] Larn RH Yang J-R. The effect of pressive deformation of austenite on the bainitic ferrte transformation in FeMnSC[43] Godet S Harlet Ph Delannay F Jacques PJ. Critical asess- steels. Mater Sci Eng A 2000;278:27891.ment of the bainite transformation of finely grained and deformed austenite. Mater Sci Forum 2003;426432:14338.[44] Godet S Jacques PJ Harlet Ph Delannay F. Simulation of the hot rolling of TRIP-assisted multiphase steels. Mater Sci Forum[45] Chang LC. Microstructures and reaction kinetics of bainite 2003;426432:12738.transformation in Si-rich steels. Mater Sci Eng A 2004;368: 17582.
high strength steels. In: Proc. of Int. Conf. On Phys. Metall. on Thermomech. Proc. of Stels and Other Metals ISU[5] Bhadeshia HKDH. The bainite transformation: unresolved 1988;2:6928.[*6] Wu KM Kagayama M. Enomoto M. Kineties of ferite issues. Mater Sci Eng A 1999;273275:5866.transformation in an Fe0.28mass% C3mass%Mo alloy. Mater Sei Eng A 2003;343:14350.[7] Enomoto M Maruyama N Wu KM Tarui T. Alloying element accumulation at ferritelaustenite boundaties belowMo alloy. Mater Sei Eng A 2003;343:1517. the timetemperature-transformation diagram bay in an FeC[8] Peet M. Babu SS Miller MK. Bhadeshia HKDH. Three- u uop qe o sseu aqod e [9] Bhadeshia HKDH Waugh AR. Bainite: An atom-probe study temperature bainite. Scripta Mater 2004;50:127781.of the inplete reaction phenomenonActa Metall 1982:30:77584.[10] Stark I Smith GDW Bhadeshia HKDH. The distribution of substitutional alloying element during the bainite transforma-[11] Bhadeshia HKDH Waugh AR. An atom-probe study of tion. Metall Trans A 1990;21A:83744.bainite. In: Proc. of Int. Conf. on Solid-Solid Phase Transfor- mations ASM 1981:9938.[12] Josefsson B Andren HO. Microstructure and thermodynamic behavior of a chromium molybdenum submerged arc weldmetal in the as welded state Advances in the science and technology of welding. ASM Int 1989:2437.[13] Jacques P Harlet P Delannay F. Critical assessment of the phase transformations occurring during the heat treatment ofAided High Strength Ferrous Alloys GRIPS' Sparkling World TRIP-assisted multiphase steels. In: Proc. Int. Conf. on TRIP-[14] Matsuda H. Kitano F. Hasegawa K. Urabe T Hosoya Y. of Steel 2002;1:12934.sheet. In: Proc. Int. Conf. on TRIP-Aided High Strength Ferrous Metallurgy of continuously annealed high strength TRIP steel[15] Takahashi M Bhadeshia HKDH. A model for the microstruc- Alloys GRIPS' Sparkling World of Steel 2002;1:11319.ture of some advanced bainitic steels. Mater Trans JIM 1991;32:68996.[16] Traint S Pichler A Hauzenberger K Stiaszny P Wermer E. Influence of siicon aluminum phosphorus and copper on hephase transformation of low alloyed TRIP-steels. In: Proc. Int. Conf. on TRIP-Aided High Strength Ferrous Alloys GRIPS[17] Garci-Mateo C Caballoro FG Bhadeshia HKDH. Acceler. Sparkling World of Steel 2002;1:1218.[18] Quidort D Brechet Y. The role of carbon on the kinetics of ation of low-temperature bainite ISIJ Int 2003;43:18215.[19] Chang LC. Microstractures and reaction kinetics of bainite bainite transformation in steels. Scripta Mater 2002;47:1516.transformation in Si-rich steels. Mater Sci Eng A 2004;368: 17582.[20] Singh SB Bhadeshia HKDH Estimation of bainite plate- thickness in low-alloyed steels. Mater Sci Eng A 1998;245:729.[*21] Wu KM Enomoto M. Three-dimensional morphology of degenerate ferrite in an FeCMo alloy. Scripta Mater[22] Aaronson HI Reymolds WT Purdy GR. Coupled-solute drag 2002;46:56974.effeets on Ferrite formation in FeCX systems. Metall Matet Trans A 2004;35A:.
[1）关野清之助.联明，1973；59 762（2）浙江大学，9Ni钢两相区热延迎和然相变（技木报告），1980，4（3） Miller D L Motal1 Trens 1e72 3 905C5）JlomHcsm M I HccneaoBaxHe no nucoonpoHmM cDJasaM H mHresAHu （4）古林英一，本正人.联额，197763；1035KpcTaaaM Hsn AH CCCP 1963 108[6）白华等.纲铁.1986（4）：60
Influence of Thermomechanical Treatment on the Structure andProperties of Cryogenic Steel 9Ni
cryogenjc steel 9Ni was investigated for its room-temperature strength and Thermomachanically treated at 660 800 900，and 1000‘C repectively thethe notch toughness at -ie6℃， with its grain size fracture morphology andmicrostructure observed.The results indicated thal higher cryogenjc toughnessor (aβ) biphase heat treatment The cryogenic loughness will increase by and strength can be obtained by thermomechanical treatment ihan by ordinaryabout 12% with the maximum room-temperature strength up to 107.8 MPa.Such an improvement is altributed mainly lo the refinement of marlensite structure much more dislocation substructures retained in tempering structureand the uniform distribution of inverse austenite.
[Abstract]This paper establishes the galubricated Reynolds equation and bump foil elastic deformation equation of dynamic radial bump foil bearing. With the boundary condition these elastohydrodynamic lubrication (EHL) equations are numerically solved by Finite Diference Method (FEM). Then this paper ana-lyzes the effect of structure and other parameters on static load performance of dynamic radial bump foilbearing which provides some reference for bump foil bearing structural improvement with increasing its static load performance..
[Abstraet 】Particle Suwrm Optimization (PSO) algorithm is an effective evolutionary algorithm. In thispaper the antibody density was introdaced into PSO to improre searching characteristics by keeping the di-versity of the population. Disk Brake is on importan part of the vehicle lo ersure the security. The improed PSO is then applied to optimal design of disk brake in order to minimize the brake time under the conditionsofgeometry constrain intensity constrain and temperafure constrain. The result shous that brake effciencyand the security of the vehicle increase.
The nonlinear research of sliding bearing journalswirl princi pal
(1.ChinaDepartment of Power Engineering North China Electric Power University Beijing 102206 China;2. 100080School of Energy Science and Engineering Harbin Instiute of Technology .Harbin 15001 .china;3. Being 3d Powr Engineering Ltd Chinese Academy of Sciences Beijing)
-NrN
Abstract: The nonlinear question of bearing joumal swirl has been researched in this article by the way ofbearing journal we go deep into discuss the alteration of swirl frequency along with bearing journal autorotation rate
nonlinear luvrication mechanism analysis. In the basic of oil film presure analysis and the swirl frequency equationof oil film force and the bearing coupling motion we give the jourmal swirl principal in different swirl velocity theload attenuating region of bearing oil film pressure distribution unstability transition region and load unstabilityhistory area. As the same time the waveform conversion between the convergence oil wedges positive pressure andextrusion speed and load. In addition give the swirl character of nonlinear dynamic oil film force. So in thisarticle a new way of rotating machine bearing joumal swirl frequency analysis and diagnosis and active control hasbeen put forward.
Key vor dS: sliding bearing; swirl frequency; nonlinear charact
状态分析可知，在惯性收敛模形成的油膜力的方向与外载荷的方向有一定的角度，相应分解成与外载荷作用在同一条直线方向相反的力和与外载荷相垂 de直力的两个分量，前者是径向弹性挤压效应 ip ，它力）则是引起轴颈涡动效应（ 直接影响轴颈径向振动大小和方向，后者（失稳分 do ）的根源，它直接影响油膜及转子运行的稳定性.同样两者对油膜压 dt力影响都是产生使压缩油模侧压力提高，分离油稷侧压力降低的效应.
由于轴颈受油膜失稳分力及失稳载荷力的作用产生的涡动，使轴颈受到的离心力增加，这个离心力又反过来进一步加大轴颈在轴承内的偏心，使 de p大于零，油膜承受的动载荷加大，收敛油樱间隙减小，发散油模间隙增大，进一步推动了轴颈的涡动，促使轴承发生液固耦合涡动现象发生.
图1轴承结构及轴颈受力运动参数示意图Di agr am1: the beari ng str uct ure and t he beari ngj our nal stress neti on paraneter schenati c di agr am
Instinto de Fisio. Unisersidade Federal do Rio de Jaseiro CP 68528. CEP 21945-970. Brazil Departamo de Engenhrig Mecinic Uninersidade Federal Fluminene Brazil* Departmento de Engenharia Metalirgica COPPE(EE Uninersidade Federaf do Rio de Janeiro CP 58510 CEP 21945.970 Rio de Joneiro Brazil
Received 6 September 1996; received in revised form 27 August 1997
Abstract
The magnetic properties in an Fe20Mo5Ni0.12C alloy hot-rolled solution treated at 1230°C and aged at 610Cfordifferent times were measured.These properties were pared with those of the same alloy previously cold-rlled solution treated at the same temperature and aged under the same conditions The results for the same alloy withoutreachinga H value of 401 Oe after 240 min. of aging.The precipitation hardening was acpanied by thermomagnetic carbon obtained by other authors have also been considered. The alloy presented magnetic hardening during aginganalysis (TMA) in the range 25800°C and X-ray diffraction. C 1998 Elsevier Science B.V. All rights reserved.
PACS: 81.40.Rs
Keywordx: Magnetic materials; Magnets; Magnetic hardening
1.Introduction
610°C [5]. However. the addition of 0.12% C de-creases the ductility in cold rolling. On the otherhand this alloy is ductile when hot-rolled.
served in FeMo [1] Fe-Mo-Ni [24] and Fe Magnetic hardening during aging has been ob-20Mo5Ni0.12C [5] alloys. These alloys are pos- sible substitutes for Vicalloy I and other Co-Femagnets. The previous results on these materialsindicate that the 0.12 wt% carbon addition to the Fe20Mo5Ni improves its magnetic propertiesand B = 8000 G are obtained after 240 min at [5]. The best results H = 448 Oe B = 15 400 G
Results presented by Wan et al.[1] on theFe-Mo alloyssuggest that the increase in coercivity can be due to a Mo-rich solidsolution (n’ phase)and 入 (FeMo) precipitation. In the thermomag-observed the first in the range of 600700°C ow- netic analysis two transition temperatures (T ) werehigher temperatures (>700°C) attributed to the ing to the supersaturated ferrite and the second atphase formed during the TMA. The first transitiontemperature decreases with the Mo addition indicating a decrease in the exchange interaction inthe α-FeMo solid solution [1].
In this work the magnetic hardening of a hot-rolled Fe20Mo-5Ni0.12C was investigated bymeans of magnetization measurements (hysteresis loops and thermomagnetic analysis TMA) andX-ray diffraction.
mated Siemens type-F diffractometer with Co( 1.79 A) and Cu (. = 1.54 A) radiations. A graphitemonochromator was used.
3. Results and discussion
2.Materials and methods
samples hot-rolled and aged at 610°C. The coercive Fig. 1 shows the curve of magnetic hardening forforce of the hot-rolled material is not so high as inthan the values reached by the alloy without car- the previously cold rolled one [5] but is still higherbon [2].
An ingot of position Fe20Mo5Ni0.12C(wt%) was obtained by induction melting under vacuum.
rolled to a plate 11 mm thick and quencbed in oil. The 43 mm thick ingot was soaked at 1220°C Then it was soaked at 1230°C hot-rolled to 4 mmthick and water quenched.
It is interesting to note that in the initial condi-tion both the cold-rolled and the hot-rolled
For the magnetic measurements cylinders of3 mm diameter were machined and then sliced in dises of 0.20.5 mm thick. The disc-shaped sampieswere treated at 1230°C for 30 min in vacuum andthen quenched in water. After this they were aged at 610°C for times varying from 15min to 20hwith samples sealed in quartz. (Table I). All these treatments were undertaken
Magnetization measurements (hysteresis loopsand TMA) were made in a vibrating sample mag- netometer (VSM) EGG-PAR model 4500. Theoven model 151H in low-vacuum atmosphere TMA data were obtained with a high-temperature(0.1 mmHg)and an applied magnetic field of 100 Oe. The heating rate was 5°C/min and thecooling rate was 2°C/min.
for 1 h at 670 and 750°C and also in the sample X-ray diffraction was carried out in samples agedaged 4 h at 610°C. It was performed in an auto-
Fig 1. Curves of magnetic hardening at 610°C: (a) hot-rolledFe20Mo-5Ni0.12C (b) cold-rolledFe-20Mo5Ni-0.12C [5] and (c) Fe20Mo-5Ni [2]
Table 1 Identificatioe of samples submitted to magnetization measurements
Sample Condition Measurements1 Hot-rolled and treated at 1230C(not agn) Hysteresis loop TMA hysteresis loop2 3 Hot-rolled treated at 1230°C and aged at 610°C for 1 b Hot-rolled treated at 1230°C and aged at 610°C for 15 min Hysteresis loop TMA bysteresis loop Hysteresis loop TMA hysteresis loop4 Hot-rolled treated at 1230°C and aged at 610°C for 4 b As sampie I affer magmetic measurements Hysteresis loop TMA bystersis loop TMA6 Hot-rolled treated al 1230°C and aged at 610°C for 20h Hot-rolled treated at 1230°C and aged at 610°C for 10 h Hysteresis lop Hysteresis loop
Fe20Mo5Ni0.12C alloys have the same values of H (see Fig. 1). In part this is due to the heattreatment at 1230°C undertaken after cold and hot rolling On the otherhand the cold-rolling processin this case the maximum value of H is reached seems to increase the magnetic hardening rate andafter 40 min in parison with 240 min for thehot-rolled alloy (Fig 1).
The optimum properties of Fe-20Mo-5Nj0.12C hot-rolled and previously treated at 1230°Cwere H =401 Oe B =6387 G B =14813 G and (BH)mx = 0.67 MG Oe obtained after 240 min at610°C (Table 2 and Fig 1).
Fig. 2 shows the TMA curves of samples 14 (seeTable 1)at increasing temperatures.Samples 1 and 2 present the same behaviour: a magnetic phasc-FeMo supersaturated) with transition temper-
Table 2Magnetic properties obtained affer aging at 610°℃
Fig. 2TMA curves at increasing temperature fr the hot-rolled Fe20Mo-5Ni0.12C alloy: (1) not aged (2) aged 15 min at610 C (3) aged 1 h at 610°C and (4) aged 4 h at 610°C.
ature (7.) 625°C (taken at the inflection point). With the precipitation of a Mo-rich phase theCCC-matrix loses Mo and a second transition istransition belongs to an a-ferrite with lower Mocontent
Sample 3 (previously aged at 610°C by 1 h)shows an intermediate situation: the magnetization increasing may indicate additional precipitationobserved. during the TMA. The transition at 714°C is then
Finally sample 4 (aged at 610°C by 4 h) presentsonly the transition at 714°C. After 4h aging at 610°C the material presents a stable structure ofoccurs during the TMA. α plus precipitates and no additional precipitation
Fig.3 shows the TMA curves of samples 14at decreasing temperatures.Samples 1 and 4 have the same behavior presenting only the transitionat 714°C. Samples 2 and 3 present a similar be- havior with another phase formed at 387C.The net magnetization of all samples are higherat decreasing temperatures (see Fig. 4 for an cxample).
decreasing temperatures for sample 5 which is Fig. 4 shows the TMA curves at increasing andsample 1 annealed by a first TMA (see Table 1) 1tditional precipitation occurs during the TMA and presented a behavior very similar to sampie 3. Ad-
Fig3.TMA cuvesat decreaingtmperaturefo thehot-rolled Fe20Mo5Ni0.12C alloy: (1) not aged (2) aged 15 min at610′C (3) aged 1 h a1 610°C and (4) aged 4 b at 610°C.
after (A) TMA dsta coliectioe of sample 1. Fig. 5. Comparison between hysteresis loops before (B) and
the first TMA) Fig. 4. TMA curve for sample 5 (sample 1 after annealing due to
ing temperature a transition at 387°C is detected. the transition at 714°C is observed. With decreas-
canindicate the precipitation of two phases during 1 h respectively. The results piled in Table 4served in a previous work [6]. It is interesting to aging (FeMo) and a (Fe Mo) C carbide ob-note the change of d value of the (1 0 0) peak offerritedeoaging.Itmcansthatinthe solubilized state the lattice parameter of ferrite is 0.29133 rmand after aging it is 0.28709 m. As expected the degree of supersaturation of ferrite decreases withthe precipitation.
Hysteresis loops were obtained before and afterB are shown in Table 3. This table shows that all the TMA data collection. The results of H B anddue to the TMA annealing The hysteresis loops ofsample 1 before and after TMA data collection are shown in Fig. 5. As shown in this figure the TMAcauses magnetic hardening in this sample withthe H increasing from 275 to 392Oe. Also an increase in the remanence B while a decrease into that obtained in FeMo[1]. In the samples the saturation B is observed.This result is similarpreviously aged by 1 and 4h the TMA causesa kind of overaging which makes the H values to decrease.
4.Conclusions
The Fe20Mo5Ni0.12C alloy hot-rolled andhardening during aging at 610°C. Optimum proper- solution-treated at 1230°C presented a magneticties were B = 6387 G B=14813 Gand (B） 0.67 MG Oe obtained after 240 min at 610°C.
samples aged at 610°C for 4 h 670 and 750°C for Figs. 68 present the X-ray diffractograms of
Table 3Coercive field (H.) remanent inductance (B ) and saturation inductance (B) before and after TMA data collection
Sample Before TMA After TMAH (Oe) B (G) B (G) H (Oe) ()’g B (G)2 275 302 3654 4811 18165 17673 392 328 4708 5169 15182 154593 4 375 401 6557 6387 14509 325 369 6919 6420 15393 15352
Fig. 6. X-ray difractogram of sample aged 4 h at 610°C (KaCu radiation)
Fig. 7. X-ray diffractogram of sample aged 1 h at 670°C (KaCo radiation).
bilized sample showed two transition temperatures Thermomagnetic analysis (TMA) of the solu-(T= 625°C and T=714°C) The precipitationofFeMo and/or (Fe Mo)C at about 660°C is responsible for the second transition temperature since it causes the reduction of the Mo content of
the α-ferrite matrix. The T: belongs to a super- saturated α and T.2 belongs to a low Mo αL.
The sample previously aged at 610°C by 1 hpresented additionalprecipitation during the TMA. On the other hand in the sample aged at610°C by 4 h no additional precipitation occurs. It
Structure Design of Motorized Spindle Bearing SUN Tao QIN Lu-fang
(Xuzhou Institute of Technology Xuzhou 221008 China
Abstract: The paper analyses the steps and methods of motorized spindle bearing design on the preconditions of Sujeaq au sanpou a Bueaq apudis pezuu jo uisp aunngs Suissnosp fq auad ujeuao Bugaaugegn jo uoelas pue peogaud uueaq au jo uogeuaap uoaas ueaq au apnu seu suauo peugisapBuueeq apuids u]
K ey words: motorized spindle; bearing preloadf lubrication
Study on Laser Global Location System for Mobile Robot
CUI Mao-yuan? DONG Zai-1i² TLAN Yan-tao’ HONG Wei
( 1. fs L Chnee Araey f Scs Sof f bfi Se &: Enining Elis Cirsiy Jfin Ggrln 30025 2. Ralatics Ialomry Oinee Aslsy f Sicxces Sheryng hrine f Aswtion Iinsing Slemgsg 1J005 )
al location problem of mbile rbt in stcturd eviroent are mainly discusd. A new iteatise search location algri using leat Ahstract his paper inrodcs a new laser location sstem of mbile robt. Based om the location sstem the appraches to glb-squtes procedure is peugosnd. The location ospesimbs caried out in onidirectioeal mobile robot platfom demonstate the effectivenes of algoeithm.
Keywords : Robot ; Pasitioeing system / Lead squures method ; Ieration
Zhao Pengfei W u W anyun Hao B aoguoNorth China U niversity of Technobgy)(U niversity of Science and Technobgy)
bad are discu ssed the m ethod to diagno se the inside circle fault outside circle fault rolling body fault and ABSTRACT The specific property and fault diagno sis method of the bearing w ith kow speed and highhow to judge the level of the fault are advanced To prove the correctness of the m ethod tow examp les areanalyzed
KEYWORDS Low speed High bad Bearing Fault diagnosis
[1 ] Zheng W. P. * Extractin of Transient Signal fromSystem M easurem ent and Control V ol 117 1995[2]孟庆丰等：滚动轴承和齿轮故障的时频域识别”.《重 型机械). 1998 No. 1[3]陈克兴、李川奇，《设备状态监测与故障诊断技术》
Nonperiodic D anam ic Response". Journal of D ynam ic
毕四富，李宁²，王亚伟，刘海萍，屠振密2（1.哈尔滨工业大学（威海）海洋学院，山东威海264209；2.哈尔滨工业大学应用化学系，黑龙江哈尔滨150001）BI Si-fu' LI Ning'² WANG Ya-wei'.²， LIU Hai-ping'，TU Zhen-mil2(1.School of the Ocean，HarbinInstitute of Technology at Weihai Weihai 264209 China;2.Department of Applied Chemistry，Harbin Institute of Technology，Harbin150o01 China)
Abstract:The bath positions of trivalent chrome black passivation for zinc-iron alloy coating were optimized by orthogonaltest method.The effects of operation conditions on the appearance and the corrosion resistance of the passivation film were alsoinvestigated. An optimal passivation technology of trivalent chrome black passivation was confirmed based on the aboveexperiments.This passivation bath contains chromium （II1) ions，nitrate ions，nickel（II） ions，phosphate and organic acidwhich is used as plexing agent.The passivation film obtained from the optimal technology has a bright and uniform blacksurface，good adhesion to the electroplatedzinc-ironalloy coating，andno whiterust generated on the passivatedzinc-iron alloycoating after exposure in neutral salt spray for 96 h.
[2]褚德威，全成军，韩玉华，等.锌-铁合金镀层黑色铬酸盐钝化工艺[J].材料保护，1995，28（7）：14-16.[3]王云燕，彭文杰，舒余德.Zn-Fe合金镀层黑色钝化工艺研究1钝化液组成成分的优化[J].电镀与涂饰，2003，22(5)：1-6.[4]杨才千，龙志林，王先友，等.Zn-Fe合金镀层的黑色钝化[J].表面技术，2001，30（2)：46-48.[5]Long ZL Zhou Y C，Xiao I.Characterization of blackchromate conversion coating on the electrodeposited zinc-ironalloy[J]. Appl Surf Sci 2003 218(1-4):123-136.[6]Bech-Nielsen G Larsen M K Jensen K A et al. Corrosionand dissolutionvalence ofachromatedZn-Fealloy studiedbyCMT and EC measurements [J]. Corros Sci 1994 36(11):1 907-1 919.[7]李秀英，三价铬钝化在生产中的应用[J].电镀与环保，2006，26(3):42-43.[8]Zaki Nabil.Producing Cr（VI)-free black finishing forautomotive applications:AESF SUR/FIN 2004 &.Interfinish2004 World Congress[C]. Chicago Illinois 2004; 494.
Abstract:An environment-friendly anti-discoloration agent（FBS） for sliver coating is developed.The optimum massconcentration of FBS is determined as 0.2 g/L by experiments of HS atmosphere exposition NaS soaking andelectrochemistry.The electrochermical experiments show that the agent has a good synergistic effect with other corrosioninhibitors，and the corrosioninhibition efficiencyisgreatlyimproved.