Microstructural evolution in 13Cr8Ni–2.5Mo–2Al martensitic precipitation-hardened stainless steel

Materials Science and Engineering A394(2005)

285–295

Microstructural evolution in13Cr–8Ni–2.5Mo–2Al martensitic

precipitation-hardened stainless steel

D.H.Ping a,?,M.Ohnuma a,Y.Hirakawa b,Y.Kadoya b,K.Hono a

a National Institute for Materials Sciences,Sengen1-2-1,Tsukuba305-0047,Japan

b Takasago R&D Center,Mitsubishi Heavy Industries Ltd.,Takasago676-8686,Japan

Received15October2004;received in revised form15November2004;accepted2December2004

Abstract

The microstructure of13Cr–8Ni–2.5Mo–2Al martensitic precipitation-hardened(PH)stainless steel has been investigated using transmis-sion electron microscopy,three-dimensional atom probe and small-angle X-ray scattering.A high number density(～1023–25m?3)of ultra-?ne (1–6nm)?-NiAl precipitates are formed during aging at450–620?C,which are spherical in shape and dispersed uniformly with perfect coherency with the matrix.As the annealing temperature increases,the size and concentration of the precipitates increase concurrently while the number density decreases.The Mo and Cr segregation to the precipitate–matrix interface has been detected and is suggested to suppress precipitate coarsening.In the sample aged for500h at450?C,the matrix decomposes into Cr-rich(? )and Cr-poor(?)regions.The decrease in the strength at higher temperature(above550?C)is attributed to the formation of larger carbides and reverted austenite.

?2004Elsevier B.V.All rights reserved.

Keywords:Martensitic PH stainless steel;Microstructure;Precipitate;Microanalysis;Three-dimensional atom probe(3DAP)

1.Introduction

Martensitic precipitation-hardened(PH)stainless steels are widely used in many engineering applications due to their high strength,high fracture toughness,good weldabil-ity and ease of machinability.They are the low-carbon steels containing certain amounts of Cr and Ni together with other substitutional elements,such as Mo,Co,Ti,Ni and Al.The high strength,combined with good toughness,is achieved by the precipitation of?ne,uniformly dispersed intermetal-lic precipitates in a martensitic?-Fe(bcc,a=0.2866nm) matrix[1–4].These?ne precipitates are formed during ag-ing at intermediate temperature(400–600?C)after the steels are solution-heat treated in the fully austenitic region,and then quenched to a completely martensitic structure.The type of the?ne precipitates depends on alloying elements. In a commercial17–4PH stainless steel,which contains Cu, nanosized bcc-Cu or fcc-Cu precipitate during aging[5].In ?Corresponding author.Tel.:+81298592717;fax:+81298592701.

E-mail address:PING.De-hai@nims.go.jp(D.H.Ping).a Fe–20Ni–23Co–0.07Al–0.17Ti PH steel,the precipitation hardening is mainly due to?ne Ni3Ti intermetallic phase[6].

A number of studies have reported that the martensitic PH stainless steels containing nickel and aluminum can be hard-ened by?ne?-NiAl precipitates(B2,a=0.2887nm)after aging at temperatures above400?C[7–11].

Transmission electron microscopy(TEM)studies re-vealed that the?-NiAl precipitates are spherical in shape and are dispersed uniformly with perfect coherency with the matrix?-Fe phase.However,these investigations were mostly carried out on the overaged samples,in which the in-termetallic precipitates were coarsened to the size that are large enough to be visible by the TEM technique.This is because the precipitates that form prior to the peak hardness condition are in a few nanometer dimensions.Thus,it is dif?-cult to obtain quantitative information on the size,distribution and chemical composition of the precipitates by TEM obser-vations.Recent three-dimensional atom probe(3DAP)study on a PH13–8stainless steel containing nickel and aluminum [12]has revealed that the precipitate composition is much lower than the stoichiometric composition of?-NiAl phase

0921-5093/$–see front matter?2004Elsevier B.V.All rights reserved. doi:10.1016/j.msea.2004.12.002

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after aging for4h at510?C.The size of the precipitates is only a few nanometers(2–8nm),and the number density is in the order of1024m?3.The observed hardening was attributed to the high-density?ne precipitates.

Compared to the wrought15–5PH and17–4PH stain-less steels,the present13Cr–8Ni–2.5Mo–2Al martensitic PH stainless steel is known to have high precipitation harden-ability.This steel is now being considered for applications as steam power plant turbines.For further improvement of the mechanical properties of these types of steels,a fundamen-tal understanding of the detailed microstructural features with various aging conditions,in particular,quantitative character-ization of the size distribution,chemical composition and vol-ume fraction of the?ne precipitates,is necessary.The present work is focused on the quantitative microstructural charac-terization of the13Cr–8Ni–2.5Mo–2Al martensitic PH stain-less steel during aging using TEM combined with3DAP and small-angle X-ray scattering.

2.Experimental procedure

The chemical composition of the steels investigated in the present study is given in Table1.Vacuum induction melting and electro slug re-melting techniques were applied to pre-pare the steel.Bars50mm in diameter were manufactured by forging.The samples were solution-heat treated at925?C for 1h followed by water quenching,then aged for4h at differ-ent temperatures ranging from450to620?C.In addition,the steel was isothermally aged at450?C for as long as500h to observe the kinetics of precipitate coarsening.Hardness was measured by a Vickers hardness tester using a1.0kgf load. Tensile tests were carried out at room temperature.

Square rods of approximately0.2mm×0.2mm×10mm were cut out from the bar sample.These rods were then electropolished to sharp needle-shape specimens for?eld ion microscopy(FIM)observation and three-dimensional atom probe(3DAP)analysis.For atom probe analyses,a locally built energy-compensated three-dimensional atom probe(ECTAP),equipped with the CAMECA optical tomo-graphic atom probe detection system[13],was employed. Atom probe analyses were performed at tip temperatures of Table1

Chemical composition of the13Cr–8Ni–2.5Mo–2Al martensitic precipitation-hardened(PH)stainless steel

Element Wt.%At.% Ni8.477.98 Cr12.3413.12 Mo2.151.24 C0.040.18 Si0.070.14 Mn0.040.04 P0.0030.006 S0.0040.008 Al1.102.25 N0.0040.016about90K under ultrahigh vacuum(<1×10?8Pa)with a pulse fraction(a ratio of pulse voltage to the static voltage) of0.2and a pulse repetition rate of1500Hz.A Philips CM200 TEM operated at200kV was also used in the present study. Thin foils for TEM observations were prepared by means of a twin-jet polishing technique using an electrolyte contain-ing15ml of perchloric acid,45ml of n-butanol and90ml of methanol.Small-angle X-ray scattering(SAXS)investi-gation was also carried out for evaluating time evolution of average size,number density and volume fraction of the?ne precipitates by using pin-hole collimated X-ray beam with Mo target(PSAXS-3S Rigaku)and two-dimensional detec-tor(Nanostar,Bruker),scattering signal collected in transmis-sion mode under vacuum.Samples are mechanically ground down to20mm for SAXS measurement.

3.Results

3.1.Mechanical property

The room temperature yield strength of the13Cr–8Ni–2.5Mo–2Al martensitic PH stainless steel aged for4h is plotted as a function of aging temperature in https://www.wendangku.net/doc/ba6670740.html,-pared to12Cr and17–4PH steels,the present steel shows a higher strength when aging temperature is below510?C. Hardness increase was observed above450?C and the maxi-mum yield strength occurs on aging at510?C.Above550?C, the usual overaging behavior occurs,which is associated with the rapid decreases of strength.For understanding the pre-cipitation hardening behavior,the isothermal aging was also carried out on the steel at450?C.Fig.2shows the hardness as a function of the annealing time at450?C.The hardness shows continuous increase up to500h,but the increase is not signi?cant from4to500h.

3.2.TEM observations

TEM observations were carried out on the samples aged at450,510,550and620?C for4h.A typical lath

martensite Fig.1.Variation of yield strength behavior in13Cr–8Ni–2.5Mo–2Al PH steel with annealing temperature.The annealing period is4h.

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287

Fig.2.Variation of hardness of13Cr–8Ni–2.5Mo–2Al PH steel with aging time at450?C.

structure was observed in all the samples.Fig.3a shows a TEM bright?eld micrograph of the sample aged at510?C for4h,which corresponds to the peak strength condition.

{010}superlattice spots are observed in the[001]zone axis pattern of the?-Fe matrix(Fig.3b),which can be indexed as

?-NiAl intermetallic compound with the B2structure.The precipitates are dif?cult to be distinguished clearly in the high-magni?cation TEM bright?eld micrograph due to their ?ne scale,and also the lack of strain contrast.The correspond-ing dark?eld image(Fig.3c)taken from the(010)spot of the ?-NiAl phase in the[001]diffraction pattern(Fig.3b)clearly shows a high number density of?ne?-NiAl precipitates in the matrix.The precipitates are fully coherent with the matrix as shown in the high-resolution TEM image(Fig.3d).The cir-cled regions show ordered lattice planes,which correspond to the?-NiAl precipitates.The precipitate size is about2–4nm and their inter-particle distances are about several nanometers (less than10nm).Such kind of precipitates were observed in all the samples aged at the temperature range from450to 620?C.Fig.4a–d show high-magni?cation TEM bright?eld micrographs of the samples aged at450,510,550and620?C for4h,respectively.Those?ne precipitates are randomly distributed in the martensite matrix.Their nucleation is apparently homogeneous,not at grain boundaries and/or dislocations,although some visible dislocations are in associ-ation with the?-NiAl precipitates.This is because the density of the precipitates is much higher than that of dislocations. One interesting point is that no signi?cant growth of the?-NiAl precipitates was noticed in the temperature range from 450to620?C.The precipitates have a spherical morphology and their size is around1–6nm in the above temperature range.

In the sample annealed at450?C for4h,only?-NiAl precipitates were observed.On the other hand,precipita-tion of carbides was detected above510?C.Fig.5a is

a Fig.3.TEM results of the13Cr–8Ni–2.5Mo–2Al PH steel aged at510?C for4h.(a)Low-magni?cation TEM bright?eld image showing martensitic lath structure;(b)electron diffraction pattern from[001]zone axes of the ?-Fe matrix.The weak super-lattice diffraction spots are from the?-NiAl precipitates with B2structure;(c)dark-?eld TEM image reveals a high density of?ne precipitates;and(d)high-resolution TEM image shows the precipitates are coherent with the matrix.

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Fig. 4.High-magni?cation TEM bright ?eld micrographs of the 13Cr–8Ni–2.5Mo–2Al PH steel aged for 4h at (a)450,(b)510,(c)550,and (d)620?C.The size difference of the ?-NiAl precipitates among the four samples is not signi?cant.

TEM bright ?eld image,the corresponding electron diffrac-tion pattern showing ?ne (less than 10nm)carbide particles (indicated by arrows).The diffraction spots can be indexed

as the [111]zone axis of ?-Fe and [11ˉ2

0]of hexagonal M 2C phase (hexagonal,a =0.289nm,c =0.456nm).The cor-responding high-resolution TEM image in Fig.5b reveals that the M 2C carbides have a well-known orientation rela-tionship with the matrix phase,{0001}M 2C //{110}?[3].At the aging temperature of 550?C,most of the carbides are spherical (Cr,Mo)23C 6,as shown in Fig.5c.The size of the (Cr,Mo)23C 6is about 20–50nm,and they are not uniformly dispersed in the matrix.At higher temperatures,in addition to the formation of these carbide precipitates (both carbide precipitates were found to be present in a small amount),austenite reversion has also been noticed,primarily at block boundaries [14].Fig.6a is a bright ?eld TEM image,which shows reverted austenite islands.Some austenite island grains are observed within martensitic laths,this may be due to a section of the austenite,which formed at a boundary and grew into the martensitic matrix grain.No austenite phase was observed in other heat treatment conditions.Fig.6b is an SAED pattern from the [110]zone axis of the austen-ite (?-Fe,fcc,a =0.366nm).Another set of [110]diffrac-tion spots,which has a larger lattice parameter in this ?g-ure,is from the (Cr,Mo)23C 6phase (fcc,a =1.066nm).This has also been con?rmed by energy-dispersive spectrometry (EDS)measurements.The corresponding dark ?eld image of the carbides is shown in Fig.6c.The reversion process is accompanied by a sharp decrease in the hardness of the

alloy.

Fig.5.(a)TEM bright ?eld image with the corresponding electron diffrac-tion pattern and (b)high-resolution TEM image show ?ne (less than 10nm)(Cr,Mo)2C carbides precipitated out from the matrix after the steel was aged at 510?C for 4h.(c)TEM bright ?eld image which shows the existence of (Cr,Mo)23C 6carbides in the samples aged at 550?C for 4h.

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Fig.6.(a)Bright?eld TEM image showing reverted austenite grains in the samples aged at620?C for4h.(b)Electron diffraction pattern from[110] zone axis of the austenite and[110]from(Cr,Mo)23C6phase.(c)Dark-?eld image taken from(Cr,Mo)23C6

phase.Fig.7.Al mappings obtained from the samples aged at450?C for(a) 0.5,(b)4,(c)50,and(d)500h.Each dot in the maps represents one Al atom.One Al-enriched region corresponds to one?-NiAl precipitate.(e) Concentration–depth pro?les of Al in the precipitates at different conditions.

3.3.3DAP analysis

Since the size of the?-NiAl precipitates is very?ne,it is

dif?cult to estimate the density,size difference and chemical

compositions(which is related to the degree of order of the ?-NiAl phase)by the TEM technique.For understanding the precipitation hardening behavior,the isothermally aged

samples were analyzed by3DAP.Fig.7is the3DAP Al maps

obtained from the samples aged for(a)0.5,(b)4,(c)50and(d)

500h at450?C.Each dot in the maps represents one Al atom.

The Al-enriched regions correspond to?-NiAl precipitates.

The size of the?-NiAl precipitates increases(up to4nm in

(d)),and the number density decreases from～1025m?3in

(b)to～1024m?3in(d)when aging time has been increased

from0.5to500h.As the size becomes larger,the Al

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Fig.8.3DAP analysis which revealed that phase decomposition occurred in the samples aged at450?C for500h.(a)Map of Al and Cr atoms.The larger dot corresponds to one Al atom,the smaller one is one Cr atom;(b) concentration–depth pro?les of the solute elements.

concentration in the precipitates appears to become higher as shown in Fig.7e.Fig.7e shows the concentration–depth pro?les of Al in the precipitates at different aging times. The Al concentration is about12,18,28and35at.%in the samples aged for0.5,4,50and500h,respectively.

After aging for500h at450?C,3DAP analysis also re-vealed that a different type of phase decomposition occurred in the remaining matrix phase.Fig.8a shows the mapping of Al and Cr atoms in a volume of about1nm×12nm×20nm. The larger dot corresponds to Al atom,the smaller one Cr atom.In the matrix phase containing?-NiAl precipitates,a high number density of Cr-enriched particles are observed. The concentration pro?le of Fe,Cr,Ni,Al,Mo,C and N in the Cr-enriched and Cr-depleted regions is shown in Fig.8b. The depth pro?les were obtained by analyzing a small subset volume(～1.5nm×1.5nm×20nm)in Fig.8a.The apparent low concentration of Ni and Al in some?-NiAl precipitates

is Fig.9.(a–d)Al maps in an analyzed region from the samples annealed at 450,510,550and620?C for4h,respectively.

because the selected subset volume for analyzing local con-centrations partially cuts the precipitates.It can be seen that the ratio of Fe and Cr are almost1:1in the Cr-rich regions. TEM observations did not show any other diffraction pattern apart from those of the matrix and?-NiAl,suggesting that the martensitic matrix decomposed into Cr-rich regions(? phase,but the concentration is signi?cantly lower than that of the equilibrium? phase)and Cr-poor regions(?phase, bcc-Fe).This matrix phase decomposition was observed only after long-term aging(500h),which is believed to give an ad-ditional hardening effect[15].No signi?cant segregation of Mo,C and N was detected at this aging stage.

Fig.9a–d are3DAP Al maps obtained from the samples annealed at450,510,550and620?C for4h,respectively. Each dot corresponds to the position of one Al atom,so the aggregates of Al atoms in the map correspond to?-NiAl precipitates.The number density of the precipitates was es-timated to be in the order of(a)～1025,(b)～1024,(c)～1024 and(d)～1023m?3.It decreases as the aging temperature in-creases,and the size of the precipitates and the concentration of Al and Ni within the precipitates increase with the aging temperature.The average precipitate size is about1nm in (a),3nm in(b),4nm in(c),and6nm in(d).The precipitates grow as the aging temperature increases,but the size is still ?ne,about4nm even at550?C,and the spacing among the precipitates is less than10nm.It is not easy to estimate the

D.H.Ping et al./Materials Science and Engineering A394(2005)285–295291 average size from the sample aged at620?C due to the low

number density of the precipitates.

3.4.SAXS results

It is not easy to estimate the average size of precipitates

accurately from the3DAP data due to the limitation of the

analysis volume and the arbitrariness of the scale.To measure

the average size of the NiAl precipitates accurately,SAXS

was measured from these samples.Fig.10a shows the scat-

tering intensity,I(q),from the samples annealed at450?C in

log scale versus square of momentum transfer q(=4πsin q/λ).

From the gradient of the slope of the linear part of the plot

(Guinier plot),the average precipitate sizes are plotted as a

function of aging time in Fig.10b.The time evolution of

the average size of precipitates were?tted using a relation-

ship of d=kt n,where k is a constant related to the activation

energy of the growth process and n the growth exponent,

which is related to the mechanism of the growth.As shown

in Fig.10b,the curve of exponent0.2shows good match with

the experimental data,while the ideal value is0.5for the case

of diffusion-controlled growth.This result indicates that the

growth of the NiAl particles is much lower than expected

from the normal diffusion-controlled growth.

In addition to the average precipitate size,the value of

I(q)at q=0,which can be extrapolated from the above?tting

in Fig.10a,gives the information on the time evolution in

the number and the total volume of the?ne precipitates.The

value I(0)can be described as follows,

I(0)=N(ρ?ρ0)2V2p V s

where N is the number of the precipitates,ρandρ0the elec-tron densities in the precipitates and the matrix,respectively. V p is the volume of the precipitates,which can be calculated from the obtained particle size.V s is the volume irradiated by X-ray,i.e.,if the sample thickness is the same,it is constant for all the samples.Assumingρandρ0are constant,the time evolution of the number of precipitates can be evaluated from I(0)and the average precipitate size.The results are shown in Fig.10c.As clearly seen from Fig.10c as well as from Fig.7, the number density of the precipitates,N/V s,decreases with the aging time.The total volume of the particle,NV,or the volume fraction normalized by the initial value is shown in Fig.10d.Although the error is large,it is almost constant or decreasing slightly under the assumption thatρ?ρ0is con-stant.Ifρ?ρ0becomes large as suggested from the atom probe data,decrease in N must be more pronounced than that plotted in Fig.10c.This suggests that the total volume or the volume fraction of the precipitate decreases during aging. This is because the concentration of Al and Ni in the precip-itates gradually increases by aging.The volume fraction of the precipitates will not be a constant parameter,which is dif-ferent with the suggestion of the classical nucleation theory. In the classical nucleation theory,it has been assumed

that Fig.10.(a)Small angle X-ray scattering pro?les from the samples aged at450?C.The gradient of the slope gives average precipitate size.(b)The average precipitate size in diameter(d:nm)against aging time(t:hours)for the samples aged at450?C.The number density ratio(c)and total volume ratio(d)of the?-NiAl precipitates as a function of aging time at450?C. The curve shown in(b)is the best?t of power for of the time evolution of d (d–t0.2).

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the chemical composition of the precipitate is constant,thus the volume fraction became constant after the nucleation and growth.

In contrast to the sample isothermally aged at 450?C,where only ?-NiAl precipitates,the evaluation of the pre-cipitate size in the samples aged at higher temperatures

was

Fig.11.3DAP analytic results,which show the concentration variety of Cr and Mo elements around one ?-NiAl precipitate in the samples aged at 510?C for 4h.(a)Ni,Al,Cr,Mo and C elemental maps in an analysis vol-ume (2nm ×2nm ×12nm).The quantitative concentration–depth pro?les are shown in (b),which reveal the segregation of Cr and Mo atoms at the precipitate/matrix interfaces.

dif?cult by SAXS because of the overlapping effect of the scattering from the ?-NiAl precipitates and the carbides.Es-pecially the high-q tail of the scattering from the carbides with a size of less than 30nm distorts the Guinier region of the ?-NiAl precipitates seriously in the samples aged at 510and 550?C.On the other hand,the size of the (Cr,Mo)23C 6carbides in the sample aged at 610?C is larger than 30nm,the scattering from which does not interfere with the Guinier region of the ?-NiAl precipitates.Hence,the size of ?-NiAl was evaluated to be 6.4±0.2nm.

3.5.Solute segregation at an interphase interface

The partitioning behavior of the other alloying elements around the ?-NiAl precipitates was carefully analyzed.Fig.11shows the results of 3DAP analysis revealing the con-centration variation of Cr and Mo around a ?-NiAl precipitate in the sample aged at 510?C for 4h.A small volume subset (2nm ×2nm ×12nm)was selected crossing one of the ?-NiAl particles as shown in Fig.9b.Fig.11a shows the Ni,Al,Cr,Mo and C elemental maps in the analyzed volume.The region between the two dashed lines corresponds to the ?-NiAl precipitate.It can be clearly seen that Cr and Mo atoms are partitioned into the ?-Fe matrix,being rejected from the ?-NiAl precipitate.The segregation of Cr and Mo atoms at the interfaces between the ?-Fe matrix and the precipitates is also observed from the elemental maps.In order to analyze the partitioning behavior of the Cr and Mo atoms quantita-tively,the integral concentration pro?les obtained from the interface are shown in Fig.11b,where the number of detected solute atoms is plotted as a function of the total number of detected atoms.In this ?gure,the slope of the plot corre-sponds to the local solute concentration,and their values in each phase are indicated in the ?gure.The integral depth pro-?les reveal that the Cr and Mo atoms are slightly enriched at the interfacial region,which is indicated by double arrows in Fig.11b,and the enrichment is in the matrix side.Since the diffusivity of Mo is slower than Ni and Al,the Mo segrega-tion at the precipitate interface would impede the growth of the particle [9,16,17].

4.Discussion

When the sample was aged at 450?C,precipitation hardening occurs as shown in Fig.2.Even after aging for 500h at this temperature,the size of the ?-NiAl precipitates was only about 4nm.The hardness keeps increasing up to 500h (Fig.2),while the number density of the NiAl precipitate decreases with aging time (Fig.7).Fig.9also shows that the number density of the precipitate is not the highest at the temperature where the peak hardness is observed.The microstructural observation suggests that the hardening is not due to the Orowan mechanism.If the hardening is governed by the Orowan mechanism,the peak hardness should appear where the precipitate density is

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the highest (inter-particle distance is the shortest).This is because the mechanism would be operative only when the precipitate size exceeds a critical value (r c ),which is about several nanometers [8,18,19].The hardness curve and the 3DAP results suggest the hardness increases as the NiAl particles grow,indicating that the hardness increase can be explained by the cutting mechanism.In the early stages of aging,the precipitates are around 1nm in diameter,thus the strengthening is associated with the stress required for dislo-cations to cut through the coherent precipitates.As discussed by Seetharaman et al.[8]for the case where dislocations cut ?ne precipitates,the contribution to the strengthening is due to three factors:coherency strains,shear modulus difference between the precipitates and matrix,and the ordering in the precipitates.Gerold and Haberkorn [18]have proposed a simple model for the strength increase due to ?ne,coherent and ordered precipitates,which can be expressed as: τ∝f 1/2r 1/2,

where τis the change in ?ow stress;f the volume fraction of precipitates;and r the radius of precipitate.

The ?ne ?-NiAl precipitates have been con?rmed to have a spherical morphology;thus,the volume fraction (f )is simply related to the precipitate size:

f =n 4

3 πr 3,

where n is the number density of the precipitates.Hardness is empirically assumed to be proportional to the precipitation strengthening portion of the ?ow stress [20].Thus,the rela-tionship between the hardness and the volume fraction can be written as follows: τ∝H v ∝f 1/2r 1/2.

Replacing f by f =n (4/3)πr 3,then H v ∝n 1/2r 2.

Since both n and r show different values under different an-nealing conditions,the hardness change will not simply fol-low r or n .As shown in Fig.7,the number density decreases with annealing time while r increases.At the beginning of the aging at 450?C,the magnitude of the number density decrease may be smaller than that of the precipitate growth,thus the hardness increases with aging time.Eventually,n and r will reach a certain value,which gives peak hardness.Based on the values of n and r measured by 3DAP and SAXS,the value of n 1/2r 2in the samples with various aging condi-tions is plotted in Fig.12.Since H v ∝n 1/2r 2,the hardness in the samples has also the same tendency as shown in Fig.12.Thus,the hardness variation agrees with the measured hard-ness as shown in Fig.2.Fig.12a is the n 1/2r 2dependence on aging time at 450?C.It keeps increasing up to 500h.The dependence of the n 1/2r 2on the aging temperature is shown in Fig.12b.The n 1/2r 2increases in the temperature range of 450–620?C.However,above 550?C,there is a tendency that the n 1/2r 2increase gradually saturates.Further increase in

the

Fig.12.Plot of the n 1/2r 2value in the samples with various aging conditions.(a)Different aging time at 450?C and (b)different aging temperature for 4h.

particle size and inter-particle distance at longer aging times will lead to a decrease in strength and hardness,since the yield strength is then governed by the Orowan mechanism.One more parameter,which has not been considered by the existing precipitation hardening theories,is that the precip-itates do not hold the stoichiometric composition when the precipitate is very ?ne.Since the Al and Ni concentration in the precipitates are lower in the early aging stage,these precipitates would not work as strong pinning sites for dislo-cations,allowing the particles to be cut by dislocations.This may be the reason why the hardness is not high,although the number density of the precipitates is higher in the early stage.Even after the 500h aging at 450?C,the sample did not overage.

The observed ?-NiAl precipitates are resistant against coarsening during aging.These results are in agreement with previous studies [7,8],but no good explanation for the sup-pression of the coarsening has been given.In 9Ni–12Cr–2Cu maraging steels,Stiller et al.[21]observed a Mo-rich phase both in the matrix and at lath boundaries,often in direct con-tact to the Ni-rich precipitates,and they became larger after longer time aging.Mo segregation has been detected by TEM in Fe–Ni–Al–Mo ferritic alloys with a relatively large ?-NiAl precipitates [22].Although the segregation of Mo in?uences the precipitate morphology by reducing the surface strain [22],no Mo segregation was detected in the similar alloy to that of the present investigation [12,23].However,the present 3DAP investigations have shown that the segregation of Mo

294 D.H.Ping et al./Materials Science and Engineering A 394(2005)285–295

atoms occurred at the precipitate/matrix interfaces when the NiAl precipitates are very ?ne.Since Mo atoms have a rela-tively larger radius,the diffusion of the Mo atoms is expected to be slower than that of other alloying elements.At 550?C,in bcc-Fe,the diffusivities of Al,Cr,Fe and Ni were estimated about Al ～9.3×10?18m 2s ?1,Cr ～5.0×10?22m 2s ?1,Fe ～2.3×10?20m 2s ?1and Ni ～3.6×10?20m 2s ?1,respec-tively [9].Mo diffusivity (～5.8×10?21m 2s ?1)was esti-mated to be lower than that (～1.6×10?20m 2s ?1)of Cr at 510?C in bcc-Fe [17].Since the diffusion of Mo is expected to be slower than that of the other alloying elements,the Mo enrichment may play an important role for the kinetics of the growth of the ?-NiAl precipitates.

The Cr atoms were also found to be rejected from the ?-NiAl precipitates and partitioned in the matrix phase with a slight enrichment at the precipitate/matrix interfaces.The Cr segregation explains the phase decomposition that occurs in the matrix phase after a long-term aging (500h)at 450?C.Since Cr is rejected from the ?-NiAl precipitates,the Cr con-centration in the matrix phase becomes higher during aging.Fig.13shows concentration–depth pro?les of Al and Cr in the samples aged at 450?C for 50h.The Al-enriched regions correspond to the NiAl precipitates.The Cr concentration in the matrix is estimated to be higher than 15at.%,which is higher than the solubility limit,but still outside the spinodal region in an Fe–Cr binary alloy.The Fe–Cr alloy with a Cr composition between the solubility limit and the spinodal re-gion will decompose via a nucleation and growth mechanism [24].This is the reason why the matrix phase decomposition into Cr-rich and Cr-poor phases occurs only after a long-term aging at 450?C.No signi?cant partitioning of the other al-loying elements was observed in the Cr-rich region.3DAP results suggest that the morphology of the Cr-rich regions is spherical with a size less than ～5nm,not inter-connected.Accompanying with the decomposition in the matrix phase,additional hardening effect would overlap with that of the ?

-

Fig.13.Concentration–depth pro?les of Al and Cr in the samples aged at 450?C for 50h for estimating the Cr content in the matrix after the formation of the ?-NiAl precipitates.

NiAl precipitates in the late stage of aging at 450?C.Such decomposition was also reported in a PH 13–8stainless steel aged at 400?C for 17,000h [12]and a PH 17–4alloy aged at 400?C for 100h [25].

A high number density of the ?-NiAl precipitates are formed in all the samples.The ?ne ?-NiAl precipitates have a spherical morphology and are coherent with the ma-trix.As the annealing temperature increased from 450to 620?C,the size was increased from 1to 6nm,and the number density was decreased from ～1025to ～1023m ?3.No other precipitate was detected in the samples aged at 450?C for 4h.However,small amount of ?ne (less than 10nm)(Cr,Mo)2C carbides were observed in the samples aged at 510and 550?C,(Cr,Mo)23C 6carbides with the par-ticle size in the range of 20–50nm were also detected.Not only carbides,but also reverted austenite phase formed at the annealing condition of 620?C for 4h.At the temper-ature lower than 550?C,the precipitation hardening is the dominant reaction;at higher temperatures,the large carbides and the reverted austenite signi?cantly deteriorates the yield strength.

5.Conclusions

The age hardening of the 13Cr–8Ni–2.5Mo–2Al PH steel is due to the precipitation of a high density (～1023–25m ?3),ultra-?ne (1–5nm)and fully coherent ?-NiAl precipitation with the B2structure during annealing at a temperature range of 450–620?C for 4h.Both the precipitate size and con-centration increase concurrently while the aging temperature increases,but the number density is decreased from ～1025to ～1023m ?3.The resistance against the coarsening of the ?-NiAl precipitates is due to the segregation of Mo and Cr atoms at the precipitate/matrix interfaces.

At 450?C,further annealing up to 500h resulted in the phase decomposition,which caused an additional age hard-ening effect.The martensitic matrix decomposed into Cr-rich (? -phase)and Cr-poor (?-phase)regions,the Cr-enriched re-gions contain about 50at.%Cr and 50at.%Fe.

Small amount of ?ne carbides ((Cr,Mo)2C and (Cr,Mo)23C 6)are formed in the martensitic matrix at higher annealing temperatures.At 620?C,in addition to the ?-NiAl precipitates and carbides,reverted austenite grains were ob-served.The decrease in strength at higher temperature (above 550?C)was attributed to the decrease in the number density of the ?-NiAl precipitates and the reversion of austenite,and also larger (Cr,Mo)23C 6carbides.Acknowledgment

This work was partly supported by Special Coordi-nation Funds for Promoting Science and Technology on “Nanohetero Metallic Materials”from the Science and Tech-nology Agency.

D.H.Ping et al./Materials Science and Engineering A394(2005)285–295295

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青岛市重点用能企业名单

南车四方机车车辆股份有限公司 青岛喜盈门集团公司 青岛广源发集团有限公司 青岛美高集团有限公司 济南山水集团有限公司青岛水泥分公司青岛正进集团有限公司 青岛大农服装有限公司 山东黄岛发电厂 青岛金晶股份有限公司 青岛恒源热电有限公司 青岛浮法玻璃有限公司 青岛压花玻璃有限公司 青岛市圣戈班韩洛玻玻璃有限公司 青岛高合有限公司 青岛浦项不锈钢有限公司 青岛北海船舶重工有限责任公司 青岛经济技术开发区热电燃气总公司 青岛赛轮子午线轮胎信息化生产示范基地 1 即墨市热电厂 青岛即发集团控股有限公司 青岛新源热电有限公司 青岛三湖制鞋有限公司 青岛正大有限公司 青岛高丽钢线有限公司 青岛北汇玻璃有限公司 即墨市双春水泥有限公司 青岛红领服饰股份有限公司 青岛恒光热电有限公司 青岛恒源化工有限公司 青岛天元化工股份有限公司 青岛海王纸业股份有限公司 青岛琅琊台酒业（集团）股份有限公司青岛胶南明月海藻工业有限责任公司 胶南易通热电有限责任公司 青岛泰发集团股份有限公司 青岛东亚轮胎有限公司

青岛康大外贸集团有限公司 胶南供电公司 胶南市水泥厂 2 胶南市海龙福利板纸有限公司 青岛振华工业集团有限公司 青岛德固萨化学有限公司 青岛龙发热电有限公司 青岛恒祥化肥有限公司 青岛世原鞋业有限公司 青岛华威建材有限公司 青岛广源发玻璃有限公司 青岛大明皮革有限公司 青岛昌新鞋业有限公司 青岛衣东纺织有限公司 青岛海尔金塑制品有限公司 山东金湖水泥有限公司青岛分公司 青岛福生食品有限公司 青岛信五皮革有限公司 青岛多福康食品有限公司 胶州天成玻璃工艺品厂 胶州市新纪元帘子布有限公司 青岛昌华集团股份有限公司 青岛热电集团金莱热电有限公司 青岛金浪热电有限公司 3 青岛泰光制鞋有限公司 青岛现代人热力发展有限公司 青岛金浪化工集团有限公司 青岛凤凰东翔印染有限公司 青岛九联集团股份有限公司 青岛海升果业有限责任公司 青岛交河技工塑料有限公司 青岛东方化工股份有限公司 海尔集团公司 青岛崂山玻璃有限公司 青岛啤酒第五有限公司

青岛恒源热电

注意：以下内容请进一步总结！ 青岛恒源热电有限公司 目标公司主要从事蒸汽、热水的生产及供应、蒸汽余热发电业务，同时提供供热管道及设施维修、安装业务。据介绍，目标公司开发了循环水供热工程项目，该项目是青岛市获批的第一个清洁发展机制(CDM)项目；前处该项目处于施工建设阶段，预计将于2009年上半年内正式投产。据介绍，目标公司主要负责临港工业区辖区内的蒸汽供应及热网管理，发电业务，对居民的用热服务。 公司成立于2001年，主要从事蒸汽、热水的生产及供应、蒸汽余热发电业务。 青岛恒源热电有限公司位于开发区B区供热范围，拥有12MW的抽凝式汽轮发电机组1台及12MW的背压机组1台，75t/h循环流化床锅炉3台和150t/h锅炉1台，最大供热能力是355t/h，担负着B区的生产、民用供热负荷，主要满足热电厂东部居民小区供热和山东科技大学供热。 青岛恒源热电有限公司位于青岛经济技术开发区临港工业区的中北部，海尔大道与渭河路交界处东北角，渭河路777号。厂区所在地东侧隔宽约100m绿化地为鑫龙物流公司，该公司东侧、距离本项目最近300m处为澳柯玛人才公寓；厂区南侧隔渭河路、绿化带100m处为东小庄村（原村庄平房已搬迁，现建有多座两层复式楼房），该村庄南侧、距离本项目约420m处为山孚日水食品有限公司；项目隔渭河路东南方向约200m处为澳柯玛工业园；西及西南方向隔海尔大道、渭河路均为浦项制铁有限公司；北侧与开发区消防大队以及正友砼业相邻。 企业所在地厂址东南距市中心约8km，东面距前湾港区约4.5km。 现有工程内容：青岛恒源热电有限公司主要服务于黄岛供热分区B 区（齐长城路以北、疏港高速以南、镰湾河以西、柳花泊和珠山以东片区（包括柳花泊），总占地面积约60平方公里）。企业现有锅炉规模为3×75t/h+1×130t/h 循环流化床蒸汽锅炉，总计约355t/h锅炉容量；发电机组规模为1×12MW C12-34.9/0.98（抽凝）+1×12MW B12-4.9/0.98（背压），总计发电装机容量24 MW。 近几年，恒源热电强化能源管理，合理调整运行方式，加强节能技术改造，企业能源管理工作上了一个新台阶，先后通过了“企业能源审计”、“热电联产机组认定”等审核认证工作，被评为“青岛市清洁生产企业”，2007年度“山东省节能先进企业”。 为进一步加强企业能源管理，完善优化企业节能减排工作，公司在本年度开始推行循环经济试点工作。目前，作为试点工作重点项目之一的企业冷渣机改造项目已基本完成，初步具备投运条件，预计本年度六月份正式投入运行。该项目是将循环流化床锅炉的人工排渣（温度一般在900℃），通过加装冷渣机把炉渣余热加热除盐水，将锅炉效率提高1-3%，同时解决人工放渣存在安全隐患、能源浪费以及不环保等问题，项目投资为85万元，年可节标煤700吨。

认识实习报告(青岛东亿热电厂)

热能与动力工程专业制热方向认识 实习报告 学院：机电工程学院 班级：热能一班 姓名：徐国庆 学号：201240502013

一．认识实习的目的和任务 1.认识实习的目的： （1）认识实习是四年制高等学校教学活动的实践环节之一； （2）认识实习是对学生进行火力发电厂主机（锅炉、汽轮机）、辅机（换热器、风机、水泵）及其制造厂的设备系统、生产工艺进行认识性训 练，对发电厂热力系统进行整体初步了解。 2.认识实习的任务： （1）对火力发电厂主机的认识实习 实习对象：锅炉本体、汽轮发电机本体。锅炉形式包括煤粉锅炉、循 环流化床锅炉、链条炉、余热锅炉等。汽轮机形式包括凝气式汽轮机、 背压式汽轮机、调节抽汽式汽轮机。 认识内容：设备外形特点、摆放位置、主要性能参数、安全生产常识。 （2）对火力发电厂辅助机械设备的认识实习 实习对象：制粉系统、除尘除灰系统、烟风系统、回热系统、润滑冷 却系统、水油净化系统等。 认识内容：设备外形特点、摆放位置、主要性能参数、安全生产常识。 （3）对火力发电厂设备系统的认识实习 实习对象：火力发电厂主机和辅机工程的系统。 认识内容：设备之间的空间关系、安全生产常识。 3.认识实习的意义 （1）强化学生对专业基础课程的理解 （2）国内火力发电厂的技术发展出现了新进展 CFB锅炉、燃气轮机、余热锅炉、超临界机组、烟气脱硫、布袋除尘、集中控制运行等新技术。 （3）认识实习有利于培养学生的职业精神 （4）认识实习有利于了解机组 （5）认识实习有利于了解机组建设过程 二．捷能汽轮机厂 （1）简介：汽轮机是火力发电厂三大主要设备之一。它是以蒸汽为工质，将热能转变为机械能的高速旋转式原动机。它为发电机的能量转换提供机 械能。 青岛捷能汽轮机集团股份有限公司始建于1950年，是我国汽轮机行业重 点骨干企业。拥有各种数控、数显等机械加工设备2200余台，以200MW 及以下“捷能牌”汽轮机为主导产品，拥有电站汽轮机和工业拖动汽轮 机两大系列产品，能够满足发电、石化、水泥、冶金、造纸、垃圾处理、燃气-蒸汽联合循环、城市集中供热等领域需求，年产能达500台/600万 千瓦以上。中小型汽轮机市场占有率居国内同行业首位，是目前国内中 小型汽轮机最大最强的设计制造供应商和电站成套工程总包商。 公司积极推进品牌战略，率先在汽轮机行业内取得了美国FMRC公司双重 ISO9001国际质量体系认证和ISO1400环境管理体系认证，率先在汽轮机 行业内第一个获得了“中国名牌产品”称号，先后获得了“全国AAA级 信用企业”、“中国优秀诚信企业”、“全国用户满意产品”、“山东

供热管网检修作业指导手册[青岛热电集团]

供热管网检修作业指导手册[青岛热电集团] 供热管网检修作业指导手册[青岛热电集团] 供热管网检修作业指导手册[青岛热电集团] 作者:佚名更新时间:2008-12-5 15:55:38 字体: 供热管网检修作业指导手册 1 总则 1.1 为使公司供热管网的维护、检修工作更为规范和科学合理，确保安全运行，制定作业指导手册。 1.2 本作业指导手册适用于公司供热管网的维护、检修及事故抢修。 本作业指导手册供热管网的工作压力限定为: 工作压力不大于1.6MPa(表压)，介质温度不大于300?的蒸汽供热管网。 1.3 管网的检修工作应符合原设计要求。 1.4 执行本作业指导手册时，尚应符合国家现行有关标准的规定。 2 术语 2.1 热网维修 热网的维护和检修。本作业指导手册中简称维修。 2.2 热网维护 供热运行期间，在不停热条件下对热网进行的维护工作。本作业指导手册中简称维护。 2.3 热网检修 在停热条件下对热网进行的检修工作。本作业指导手册中简称检修。 2.4 热网抢修

供热管道设备突发故障引起蒸汽大量泄漏，危及管网安全运行或对周边环境、人身安全造成威胁时进行的紧急检修工作。本作业指导手册中简称抢修。 2.5 供热管网 由热源向热用户输送和分配供热介质的管线系统。本作业指导手册中简称热网。 3 维护、检修机构设置、检修人员及设备 3.1 维护、检修机构设置及人员要求 3.1.1客户服务中心是公司高新区内供热管网运行、调度、维护、检修的责任机构，负责高新区内供热管网的维护、检修工作。 3.1.2 供热管冈的维护、检修人员必须经过培训和专业资格考 试合格后，方可独立进行维护、检修工作。供热管网维护、检修人员必须熟悉管辖范围内的管道分布情况、设备及附件位置。维护、检修人员必须掌握管辖范国内供热管线各种附件的作用、性能、构造以及安装操作和维护、检修方法。 3.1.3检修人员出门检修时应穿公司工作服，配戴上岗证，注意礼貌用语，维护公司形象。 3.2 维护、检修用主要设备与器材 3.2.1 供热管网的维护检修部门，应备有维护、检修及故障抢修时常用的设备与器材。 3.2.2检修设备、工具平时摆放在规定位置，检修设备和专用工具要有专人保管，所有设备、工具应保证完好，须保证检修时能够立即投入使用。检修物资也应分门别类码放整齐，方便查找，以保证检修、抢修时不会因为寻找物资配件而耽误时间。每次检修完后都应检查备品备件数量，发现不够时要及时与物质采购部联系进行必要地补充，确保检修时不会因无备品备件而影响检修时间与质量。

青岛西海岸公用事业集团易通热电有限公司新能源分公司_中标190922

招标投标企业报告 青岛西海岸公用事业集团易通热电有限公司新 能源分公司

本报告于 2019年9月22日 生成 您所看到的报告内容为截至该时间点该公司的数据快照 目录 1. 基本信息：工商信息 2. 招投标情况：中标/投标数量、中标/投标情况、中标/投标行业分布、参与投标 的甲方排名、合作甲方排名 3. 股东及出资信息 4. 风险信息：经营异常、股权出资、动产抵押、税务信息、行政处罚 5. 企业信息：工程人员、企业资质 * 敬启者：本报告内容是中国比地招标网接收您的委托，查询公开信息所得结果。中国比地招标网不对该查询结果的全面、准确、真实性负责。本报告应仅为您的决策提供参考。

一、基本信息 1. 工商信息 企业名称：青岛西海岸公用事业集团易通热电有限公司新能 源分公司 统一社会信用代码：91370211334195493K 工商注册号：370211120004502组织机构代码：334195493法定代表人：赵军田成立日期：2015-04-23 企业类型：有限责任公司分公司(非自然人投资或控股的法人 独资) 经营状态：注销 注册资本：/ 注册地址：山东省青岛市黄岛区相公山路723号 营业期限：2015-04-23 至 / 营业范围：为上级公司联系业务。（依法须经批准的项目，经相关部门批准后方可开展经营活动）联系电话：*********** 二、招投标分析 2.1 中标/投标数量 企业中标/投标数： 个 （数据统计时间：2017年至报告生成时间）

2.2 中标/投标情况（近一年） 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 2.3 中标/投标行业分布（近一年） 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 2.4 参与投标的甲方前五名（近一年） 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 2.5 合作甲方前五名（近一年） 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 三、股东及出资信息 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 四、风险信息 4.1 经营异常() 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 4.2 股权出资() 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 4.3 动产抵押() 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。 4.4 税务信息() 截止2019年9月22日，根据国内相关网站检索以及中国比地招标网数据库分析，未查询到相关信息。不排除因信息公开来源尚未公开、公开形式存在差异等情况导致的信息与客观事实不完全一致的情形。仅供客户参考。

青岛热电集团有限公司简介

青岛热电集团有限公司成立于1993年，属于国有独资大型热电联产企业，主要担负着青岛市企、事业单位和居民供热及部分发电任务，同时，供热市场辐射黄岛、平度、莱西、即墨、城阳等县市区域。集团公司先后成立了工程公司和具有甲级设计资质的设计院，逐步形成了热电联产、区域锅炉、热网输配等多种供热形式并存，集供热、发电、热力设计、工程施工、热力产品制造经营为一体的完整产业链。 目前，热电集团为全省地方最大供热企业。企业资产总额48亿元，年销售收入16.2亿元，所属企业16个，职工2200余人，年供蒸汽312万吨，年发电能力9.3万千瓦，已建成蒸汽管网145.43公里，热水管网1552.93公里，供（换）热站294座，供热面积3561万平方米，拥有单位用户292家，居民用户28.8万余户。 集团公司先后被评为全国AAA级信用企业、全国建设系统文明服务示范窗口单位、思想政治工作先进单位、企业文化建设先进单位、精神文明建设先进单位；山东省文明单位、节能先进企业、思想政治工作优秀企业；青岛市和工商年度免检企业、安全生产先进单位、廉洁勤政先进单位；山东省供热协会副理事长单位。 自成立以来，公司始终秉承“关爱社会、服务民生”的企业宗旨和“励精图治、锲而不舍”的企业精神，贯彻科学发展，创新经营管理，实现了企业快速发展。1996年在全国供热行业首家推出社会服务责任赔偿制度，1997年在山东省供热行业首家进行了股份制改造，1998年在山东省供热行业首家成功地进行了集团产权制度改革，1999年在全国同行业中首家通过了ISO9001国际质量认证，并先后通过了ISO14001环境管理体系和GB/T28001-2001职业健康安全管理体系认证，2001年公司成为全国供热行业中首家申请注册服务商标的企业，推出“暖到家”服务品牌，并被评为山东省著名商标和服务名牌。“青岛热电”正在逐步步入标准化、规范化、品牌化的发展轨道。 招聘专业及人数: 1、结构专业1人(研究生)； 2、建筑专业1人(研究生)； 3、技经专业1人(研究生)； 4、焊接技术与工程1人； 5、无损检测专业1人；

五大电力发电厂及下属详细

华能集团所属电厂： 华能丹东电厂华能大连电厂华能上安电厂华能德州电厂华能威海电厂华能济宁电厂华能日照电厂华能太仓电厂华能淮阴电厂华能南京电厂华能南通电厂华能上海石洞口第一电厂华能上海石洞口第二电厂华能长兴电厂华能福州电厂华能汕头燃煤电厂华能汕头燃机电厂华能玉环电厂华能沁北电厂华能榆社电厂华能辛店电厂华能重庆分公司华能井冈山电厂华能平凉电厂华能岳阳电厂华能营口电厂华能邯峰电厂 大唐集团所属： 长山热电厂湖南省石门电厂鸡西发电厂洛阳首阳山电厂洛阳热电厂三门峡华阳发电公司河北马头电力公司唐山发电总厂北京大唐张家口发电总厂兰州西固热电有限公司合肥二电厂田家庵发电厂北京大唐高井发电厂永昌电厂北京大唐陡河电厂南京下关发电厂安徽淮南洛河发电厂保定热电厂略阳发电厂微水发电厂峰峰发电厂含岳城电站天津大唐盘山发电公司内蒙大唐托克托发电公司保定余热电厂华源热电有限责任公司阳城国际发电有限公司辽源热电有限责任公司四平发电运营中心长春第二热电有限公司晖春发电有限责任公司鸡西热电有限责任公司佳木斯第二发电厂台河第一电厂江苏徐塘发电有限公司安徽省淮北发电厂安徽淮南洛能发电公司安阳华祥电力有限公司许昌龙岗发电有限公司华银电力株洲发电厂华银株洲发电公司金竹山电厂华银金竹山火力发电厂湘潭发电有限责任公司湖南省耒阳发电厂灞桥热电有限责任公司灞桥热电厂陕西渭河发电厂陕西延安发电厂陕西韩城发电厂永昌发电厂甘肃甘谷发电厂甘肃八0三发电厂甘肃连城发电厂甘肃兰西热电有限公司广西桂冠电力股份公司桂冠大化水力发电总厂广西岩滩水电厂陈村水力发电厂王快水电厂张家界水电开发公司贺龙水电厂鱼潭水电厂陕西石泉水力发电厂石泉发电有限责任公司甘肃碧口水电厂百龙滩电厂华电所属： 1中国华电工程(集团)有限公司2华电煤业集团有限公司3华电财务有限公司4华电招标有限公司5华信保险经纪有限公司6北京华信保险公估有限公司7河北热电有限责任公司8包头东华热电有限公司（在建）9内蒙古华电乌达热电有限公司（在建）10华电国际电力股份有限公司11华电国际电力股份有限公司邹县发电厂（扩建）12华电国际电力股份有限公司莱城发电厂13华电国际电力

(集团发布)青岛热电集团有限公司关于实施供热计量收费工作的意见

青热电〔2010〕121号 青岛热电集团有限公司 关于实施供热计量收费工作的意见 各单位、处室： 为全面贯彻《山东省物价局、山东省住房和城乡建设厅关于推进供热计量改革的指导意见》，根据青热办【2010】25号文件要求，自2010年开始，新供热建筑及完成供热计量改造的既有居住建筑，取消以面积计价收费，实行按用热量计价收费，为做好供热计量收费工作，经研究确定以下实施意见： 一、实施计划 （一）对已经改造完成的既有居住建筑实施供热计量收费，明细如下：第一热力海信慧谷、丰华园、弘信花园、都市名家小区；第二热力公司天宝苑小区；金河热力公司荣馨苑小区。

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