冶金專業(yè)外文翻譯----鉻-鉬-_v鋼的回火脆性

1、<p>　　鉻-鉬- V鋼的回火脆性</p><p>　　可逆回火脆性（RTB）－ 在500-650℃溫度范圍內(nèi)回火或緩慢冷卻的鋼的脆性，它被認為是造成在前奧氏體晶粒邊界形成雜質(zhì)（P，銻，錫，砷） [1-4]的原因。然而，有資料顯示[5]，不僅有這些進程，而且有其他進程，像在500-600℃淬火鋼有助于回火脆性在發(fā)生。</p><p>　　在這項工作的關(guān)注是淬火鉻鉬合金，以防止在

2、釩和磷的含量RTB的鋼回火脆化。</p><p>　　熱件1（0.35％的V，0.015％P）重四十一噸是偽造的，以一個550毫米的大小, 15毫米厚的板材被切斷他們。另外，在一個100公斤的感應(yīng)爐加熱熔化。</p><p>　　釩被添加到在一個重達16公斤的鑄塊中，它被鍛造、軋制成10毫米厚的金屬板。在所有的10-15毫米厚加熱板放在油中980 ℃（1小時）淬火。從板淬火和在100-76

3、0℃回火10小時后制備樣品。</p><p>　　經(jīng)過淬火和低溫回火，所有加熱件呈現(xiàn)一個典型的馬氏體結(jié)構(gòu)，但在高（超過600攝氏度）的溫度回火后得到回火索氏體結(jié)構(gòu)。在實驗室加熱的奧氏體晶粒尺寸較?。?-9級）比商業(yè)熱（4級），它們的結(jié)構(gòu)部件具有更大的分散性和均勻性的。</p><p>　　脆性是從Tso 和 Ttemper變化中來確定， Ttemper是淬火鋼的回火溫度，Tso是韌-脆轉(zhuǎn)變

4、溫度，它具有最完整的脆化特征。</p><p>　　Tso的確定是通過5×5×27.5毫米的切口1毫米深（根半徑0.25毫米）沖擊試驗樣品。Tso被認為是在斷裂50％纖維下的測試溫度。拉伸強度通過5個直徑為3毫米的樣品確定的為20°。</p><p>　　圖1顯示了熱件1回火溫度與力學(xué)性能的變化。這種熱鑄塊的力學(xué)性能在100-600 ℃回火后幾乎不變。當回火溫

5、度從600-760度提高時，強度特性急劇下降，而韌性增加。</p><p>　　更改影響回火鋼的穩(wěn)定的情況下釩濃度，因為不含釩的鋼的強度會下降至約500 °。</p><p>　　熱件1開始時回火溫度大約300℃，Tso的增加達到500-600度的最高值。如同淬火條件對峰值相比增加了100 ℃?；鼗饻囟冗M一步提高，Tso 會下降，這與削弱開始相吻合。</p><

6、;p>　　對于熱件2-7的Tso和 Ttemper變化的總體特點是相同的，盡管在較低溫達到峰巔。這類鋼在是Tso?90℃條件冷卻，低于熱1大概 - 70 ± 10 ℃。在730 ℃后的Tso鍛煉價值與熱2-7（-110至-130 °）實際上是相同。</p><p>　　它的峰高和立場取決于釩鋼的內(nèi)容。當釩濃度變化從0到0.55％的高峰上升％60 °，同時轉(zhuǎn)移?100 °

7、;。比較熱件1和3，6，在史前的冶金學(xué)上有區(qū)別的，但相同釩含量（?0.3％）是相同的，人們可以看到，在Tso峰地區(qū)的增長幾乎是與淬火條件相同。</p><p>　　在0.005和0.022％的磷的熱處理后的測試中，磷對Tso沒有什么影響。</p><p>　　在回火期間的鋼的力學(xué)性能改變顯然取決于精細結(jié)構(gòu)的變化。圖3顯示了在熱淬火Tso顯微條件后，在600℃回火，對應(yīng)于對強度特性極值邊緣與

8、Tso高峰在760 °的回火之后，強度和Tso達到最低的。</p><p>　　淬火后的結(jié)構(gòu)由帶有位錯的板條馬氏體組成。在板條彼此間略有錯誤導(dǎo)向，平均寬度?0.3μ和長度?5μ，并歸為?5 ×5μ。板條充滿均勻分布的密度位錯?011cm-2。硬質(zhì)合金階段沒有觀察到淬火鋼。</p><p>　　在600℃回火10小時的后混亂陣列的整體性質(zhì)和晶體的碎片將被保留。木板條的平均

9、大小（寬度和長度）它的平均規(guī)模保持不變。唯一的變化是明顯的沉淀分散的碳化物階段。粒子的大小是150至200 A，平均密度為~1015cm-2。這些沉淀在混亂通過板條馬氏體空間分布。帶長條形直徑為?250A和?2?103A的較大沉淀物位于深的晶體的邊界附近（板條馬氏體）。這種類型的技術(shù)被作為MTCa電石鑒定。</p><p>　　在760 ℃回火后急劇變化的鋼的結(jié)構(gòu)：位錯是不規(guī)則的。大部分分布在整提失調(diào)，改建后更適

10、合了，與位錯領(lǐng)域相吻合。單元格的平均規(guī)模為?0.5μ。該階段碳化物形態(tài)的也有變化。微細分散的大小圓形的沉淀?250 A的位錯在位于網(wǎng)絡(luò)的連接處。隨著他們有更大的圓形沉淀?2?103 A位于邊界的十字路口。</p><p>　　析出的碳化物回火后的存在也證實了熱物理化學(xué)相分析熱件1。當回火溫度由原來?300-600度的滲碳體碳化物的數(shù)量逐漸減少， M7C3增加。強有力的碳化物含量形成的殘留增加元素：鉻，鉬，和釩.，

11、這個過程在500-600度是特別活躍。這很可能是VC和 MO2C膜類型不理化分析檢測，在高度分散的碳化物階段也有更多的熱力學(xué)穩(wěn)定相[6]。</p><p>　　因此，氫脆回火過程中鉻鉬釩鋼，作為對Tso與Ttemper出現(xiàn)明顯的高峰表現(xiàn)，從我們的實驗結(jié)果。它的峰值并沒有關(guān)聯(lián)的磷在鋼存在，但與釩的濃度各不相同，在碳化物的形成，與繼承發(fā)展脫位陣列的溫度范圍內(nèi)位置。Tso達到高峰時的溫度在Fe3C的變換更穩(wěn)，與M7C3

12、類似的現(xiàn)象已在鉻鋼觀察[5]。</p><p>　　脆化可能是由于體積和邊界的影響。在鉻鉬釩鋼的的情況二次硬化，體現(xiàn)在在 Ttemper = 300 - 500 ℃性能的增強可能與這些相關(guān)的影響有關(guān)。據(jù)了解，[7]，隨著晶粒尺寸不變（或板條馬氏體晶粒常量大?。㏕so通過εTso=σy+ C與低屈服強度σy（σ0.2相關(guān)的關(guān)系）同期相聯(lián)系，其中C和ε為常數(shù)。熱件1Tso呈線性關(guān)系中的屈服強度為σo.2大的變化范圍，

13、并只有在400-600度，特殊的碳化物開始制定改變，有偏離線性的關(guān)系。當然，這些過程首先影響邊界地區(qū)，那里的條件是為硬質(zhì)合金階段準備，由于晶體結(jié)構(gòu)的幾何缺陷有利合金相的形成?？梢栽O(shè)想，Tso峰值是由于在邊界條件的變化。這個假設(shè)是證實了電子顯微鏡分析結(jié)果：在矩陣（位錯密度，碎片大小條件）引起材料的[8，9]實力的增強，保持不變直至Ttemper = 600 °。這也證實了結(jié)構(gòu)進行檢查和對回火溫度與屈服強度的變化。因此，如果發(fā)生任

14、何更改的邊界，增加脆化，那么，Tso同Ttemper變化將有虛線的形式。</p><p>　　所有解釋RTB[1，2，10]的理論都是建立在對邊界的影響起主導(dǎo)作用基礎(chǔ)上的。然而，前奧氏體晶粒邊界損害并非我們的典型調(diào)查（熱件1紋裂在淬火、回火后顯現(xiàn)出來），沒有磷的影響，這顯然是對鉬的存在鋼鐵中的解釋。</p><p>　　由此可見，由實驗數(shù)據(jù)C–Mo–V鋼回火過程中的脆性主要受碳化物形成的影

15、響，更加準確地重建碳化物和的這些過程影響因素。</p><p>　　這原種影響的理可以概括介紹如下，隨回火溫度的增加滲碳體開始凝聚，在250-350℃開始沉淀。更有效地遏制電石，在鋼中容納更多的碳化物成形元素比鐵中350℃左右開始的特殊碳化物的原子核鋼的成形，尤其是M7C3 [8]。沉淀物均勻分布在整個體積的位錯位置，低角度的板條馬氏體的界限，高角度的板條殖民地邊界，他們加強了矩陣和削弱（脆化）邊界。</p

16、><p>　　在二次硬化溫度時，很可能最大程度的削弱碳化物與基體或與沉淀、連貫的最大密度的邊界，i.e.。如元素釩，促進碳化物細化，從而增加二次硬化[8]，增加脆化。</p><p>　　在沉淀和聚結(jié)硬質(zhì)合金階段，出現(xiàn)了混亂的位錯的同時，邊界變得更加完善和矩陣被削弱。這兩種效應(yīng)導(dǎo)致韌性、脆性轉(zhuǎn)變溫度下降，這是從實際觀察的結(jié)構(gòu)。</p><p><b>　　結(jié)論

17、</b></p><p>　　1。 15Kh3MFA類型的鋼都容易脆化，在給定的回火溫度下達到高峰。在高峰期（Tso）的上限，在正值溫度下材料開始削弱。 2。該峰的高度和回火溫度也相應(yīng)增加時，幾乎線性釩濃度從0提高到0.55％，但他們是在0.005-0.022％的磷濃度限制下表現(xiàn)出來的。 3?；鼗痄摰拇嘈匀Q于在碳化物相變（從滲碳特殊碳化物），保留的位錯優(yōu)先發(fā)生在碎片的邊界。<

18、;/p><p><b>　　文獻引用</b></p><p>　　1.L. M. Utevskii, Temper Brittleness of Steels [in Russian], Metallurgizdat, Moscow (1961),p. 138.</p><p>　　2. P. B. Mikhailov-Mikheev, Therm

19、al Embrittlement of Steels [in Russian], Mashgiz, Moscow--Leningrad (1956), p. 56.</p><p>　　3. J. Hollomon, Trans. ASM, 36, 473 (1946).</p><p>　　4. E. Houdremont, Special Steels [Russian transla

20、tion], Vol. I, Metallurgiya, Moscow (1966),p. 455.</p><p>　　5. V. A. Korablev, Yu. I. Ustinovshchikov, and I. G. Khatskelevich, "Embrittlement of chromium steels with formation of special carbides,"

21、; Metalloved. Term. Obrab. Met., No. I,16 (1975).</p><p>　　6. A. P. Gulyaev, I. K. Kupalova, and V. A. Landa, "Method and results of phase analysis</p><p>　　of hlgh-speed steels," Zavo

22、d. Lab., No. 3, 298 (1965).</p><p>　　7. J. Heslop and N. Perch, Phil. Mag., ~, No. 34, 1128 (1958).</p><p>　　8. V. V. Rybin et al., "The mechanism of hardening of sorbite-hardening steel an

23、d the possibility of determining it theoretically and experimentally," in: Metal Science [inRussian], No. 17, Sudostroenie, Leningrad (1973), p. 105.</p><p>　　9．L. K. Gordienko, Substrucutral Hardening

24、of Metals and Alloys [in Russian], Nauka, Moscow(1973), p. 64.</p><p>　　10．E. E. Glikman et al., "Nature of reversible temper brittleness," Fiz. Met. Metalloved.,36, 365 (1973).</p><p>

25、;　　11. Oliver WC, Pharr GM (2004) J Mater Res 19:3</p><p>　　12. Kim JY, Lee BW, Read DT, Kwon D (2005) Scr Mater 52:353</p><p>　　13. Kim JY, Lee JS, Lee KW, Kim KH, Kwon D (2006) Key Eng Mater 3

26、26–328:487</p><p>　　14. Kim JY, Lee JJ, Lee YH, Jang JI, Kwon D (2006) J Mater Res 21:2975</p><p>　　15. Kim JY, Kang SK, Lee JJ, Jang JI, Lee YH, Kwon D (2007) Acta Mater 55:3555</p><

27、p>　　16. Dowling NE (1993) Mechanical behavior of materials. Prentice Hall, Englewood Cliffs</p><p>　　17. Kim JY, Lee KW, Lee JS, Kwon D (2006) Surf Coat Technol 201:4278</p><p>　　18. DIN 1717

28、5-79 (1979) Seamless steel tubes for elevated temperatures</p><p>　　19. Ahn JH, Kwon D (2001) J Mater Res 16:3170</p><p>　　20. Dieter GE (1988) Mechanical metallurgy. McGraw-Hill, Singapore</

29、p><p><b>　　外文原文</b></p><p>　　TEMPER BRITTLENESS OF Cr--Mo-- V STEEL</p><p>　　Reversible temper brittleness (RTB) – embrittlement of steels during tempering or slow cooling

30、in the temperature range of 500-650℃ is considered to result from the formation of impurity segregates (P, Sb, Sn, As) in prior austenite grain boundaries [1-4]. However, there are data indicating [5] that not only these

31、 processes but other processes favoring embrittlement occur in quenched steel during tempering at 500-600 ° .</p><p>　　This work concerns embrittlement during tempering of quenched chromium steels alloy

32、ed with molybdenum to prevent RTB in relation to the vanadium and phosphorus concentrations.</p><p>　　Ingots of heat 1 (0.35% V, 0.015% P) weighing 41 tons were forged to a size of 550 mm. Plates 15 mm thick

33、 were cut from them. The other heats were melted in a 100-kg induction furnace with use of ZhS-0 iron in the charge.</p><p>　　Vanadium was added to the steel during pouring of an ingot weighing 16 kg, which

34、was forged and rolled to a plate 10 mm thick. All heats in the form of plates 10-15 mm thick were oil quenched from 980 ° (1 h). Samples were prepared from the plates after quenching and after tempering at 100-760 &

35、#176; for 10 h.</p><p>　　After quenching and low-temperature tempering, all heats had a typical martensitic structure, but after high-temperature tempering (above 600 ° ) a sorbite structure. The austen

36、ite grain size of the laboratory heats was smaller (grade 8-9) than in the commercial heat (grade 4), with greater dispersity and homogeneity of the structural components.</p><p>　　Embrittlement was determin

37、ed from the variation of Tso with Ttemper , where Ttemper is the tempering temperature of the quenched steel and Tso is the ductile-- brittle transition temperature, which most completely characterizes embrittlement.<

38、/p><p>　　Tso was determined on impact test samples 5 × 5 × 27.5 mm with a notch 1 mm deep (root radius 0.25 mm). Tso was taken as the testing temperature at which the fracture was 50% fibrous. The ten

39、sile strength was determined at 20 ° on five samples 3 mm in diameter.</p><p>　　Figure 1 shows the variation of the mechanical properties of heat 1 with the tempering temperature. The mechanical propert

40、ies of this heat are almost constant after tempering at 100-600 ° . When the tempering temperature is raised from 600-760 ° the strength characteristics decrease sharply, while the ductile characteristics incre

41、ase.</p><p>　　Changing the vanadium concentration affects the stability of the steel during tempering. For the steel without vanadium the strength begins to decrease around 500 °.</p><p>　　

42、For heat 1, beginning with tempering at temperatures around 300 °, Tso increases and reaches a maximum value at 500-600 °. As compared with the quenched condition the increase of Tso at the peak is 100 °.

43、With further increase of the tempering temperature Tso decreases, which coincides with the beginning of weakening.</p><p>　　For heats 2-7 the overall character of the variation of Tso with Ttemper is the sam

44、e, although the peaks occur at lower temperatures. For the steels in the quenched condition Tso is ~90 ° lower than for heat 1, and amounts to -70 ± 10 °. After tempering at 730 ° the value of Tso is

45、practically the same for heats 2-7 (-110 to -130°).</p><p>　　The height of the peak and its position depends on the vanadium content of the steel. When the vanadium concentration is changed from 0 to 0.

46、55% the peak rises %60 ° and at the same time shifts ~100 °. Comparing heats 1 and 3, 6, differing in their metallurgical prehistory but similar in vanadium content (~0.3%), one can see that the increase of Tso

47、 in the region of the peak is almost the same as for the quenched condition.</p><p>　　In amounts of 0.005 and 0.022%, phosphorus has no effect on Tso after the heat treatments tested.</p><p>　　T

48、he changes in the mechanical properties of the steel during tempering evidently depend on changes in fine structure. Figure 3 shows the microstructure of heat in the quenched condition, after tempering at 600 ° , co

49、rresponding to the edge of the plateau of the strength characteristics and the peak of Tso, and after tempering at 760 ° , where the strength and Tso are lowest.</p><p>　　After quenching, the structure

50、consists of lath martensite with well-developed dislocation arrays. The laths are slightly misoriented with respect to each other, with an average width of ~0.3 μ and length ~5μ, and are grouped in colonies ~5 × 5μ

51、. The laths are filled with evenly distributed dislocations with a density ~1011cm-2. No carbide phase was observed in the quenched steel.</p><p>　　After tempering at 600 ° for 10 h the overall characte

52、r of the dislocation arrays and the fragmentation of the crystals are retained. The average size of the laths (width and length)and the average size of the colonies remain unchanged. The only noticeable change is the pre

53、cipitation of finely dispersed carbide phases. The size of the particles is 150-200 A and the average density ~1015cm-2. They are precipitated on dislocations and evently distributed through the bulk of the martensite la

54、ths. L</p><p>　　The structure of the steel changes sharply after tempering at 760 °: the dislocations are polygonized. The dislocations distributed throughout the bulk are rebuilt into energetically mor

55、e suitable configurations–cell walls–not coinciding with long-range fields. The average size of the cells is ~0.5 µ. The morphology of the carbide phases also changes. Finely dispersed rounded precipitates with a si

56、ze of ~250 A are located in the junctions of the dislocation network. Along with them there are lar</p><p>　　The presence of carbide precipitates after tempering was also confirmed by physicochemical phase a

57、nalysis of heat 1. When the tempering temperature is raised from ~300-600 ° the quantity of Fe3C carbide gradually decreases and the quantity of M7C3 increases. The concentration of strong carbide-forming elements

58、in the residues increases: Cr, Mo, and V. This process is particularly well developed at 500-600 ° . It is highly probable that among the highly dispersed carbide phases there are also mo</p><p>　　It fo

59、llows from our experimental results that embrittlement occurs during tempering of Cr–Mo–V steels, manifest as a distinct peak on the curve of Tso vs Ttemper. The peak is not associated with the presence of phosphorus in

60、the steel but varies with the concentration of vanadium and is located in the temperature range of carbide formation, which develops with inheritance of the dislocation arrays. Tso reaches a peak at those temperatures wh

61、ere Fe3C transforms to more stable M7C3. A similar phen</p><p>　　Embrittlement may be due to both volume and boundary effects. In the case of Cr–Mo–V steels the secondary hardening manifest in the increase o

62、f the strength at Ttemper = 300–500 ° may be associated with the first of these effects. It is known [7] that with a constant grain size (or constant size of colonies of martensite laths) Tso is associated with a lo

63、w yield strength σy (σ 0.2 in first approximation) by the relationship εTso = σy + C, where C and ε are constants. For heat 1 Tso varies linear</p><p>　　All theories explaining RTB [1, 2, 10] are based on th

64、e recognition of the dominant role of boundary effects. However, damage in the boundaries of prior austenite grains was not typical in our investigation (the fracture of heat 1 after quenching and after tempering was qua

65、sibrittle), and no effect of phosphorus was observed, which is evidently explained by the presence of molybdenum in the steel.</p><p>　　It follows from the experimental data that the embrittlement of C–Mo–V

66、steels during tempering is affected mainly by carbide formation, more precisely the rebuilding of carbides and factors affecting these processes.</p><p>　　The mechanism of this effect can be presented in gen

67、eral terms as follows. With increasing tempering temperatures cementite begins to coalesce, precipitating at 250-350 °. In steels containing more effective carbide-forming elements than iron the formation of nuclei

68、of special carbides begins around 350 °, especially M>C3 [8]. Precipitating evenly throughout the volume on disloca tions, low-angle boundaries of martensite laths, and high-angle boundaries of colonies of laths,

69、 they strengthen the </p><p>　　It is probable that the boundaries are weakened most with the maximum density of carbides coherent with the matrix or with their precipitation, i.e., at secondary hardening tem

70、peratures. Such elements as vanadium, promoting refining of carbides and thus increasing secondary hardening [8], increase the embrittlement.</p><p>　　With precipitation and coalescence of carbide phase, occ

71、urring simultaneously with polygonization of dislocations, the boundaries become more perfect and the matrix is weakened. Both these effects lead to a drop of the ductile-brittle transition temperature, which is in fact

72、observed.</p><p>　　CONCLUSIONS</p><p>　　1. Steels of the 15Kh3MFA type are susceptible to embrittlement, which reaches a peak at a given tempering temperature. The upper limit of the peak (Tso)

73、coincides with the temperature at which the material begins to weaken.</p><p>　　2. The height of the peak and the tempering temperature corresponding to it increase almost linearly when the vanadium concentr

74、ation is raised from 0 to 0.55%, but they are independent of the phosphorus concentration within limits of 0.005-0.022%.</p><p>　　3. Temper brittleness of the steel investigated depends on the change in the

75、carbide phase (from cementite to special carbides) that occurs with retention of the dislocation arrays preferentially in the boundaries of fragments.</p><p>　　LITERATURE CITED</p><p>　　1.L. M.

76、Utevskii, Temper Brittleness of Steels [in Russian], Metallurgizdat, Moscow (1961),p. 138.</p><p>　　2. P. B. Mikhailov-Mikheev, Thermal Embrittlement of Steels [in Russian], Mashgiz, Moscow--Leningrad (1956)

77、, p. 56.</p><p>　　3. J. Hollomon, Trans. ASM, 36, 473 (1946).</p><p>　　4. E. Houdremont, Special Steels [Russian translation], Vol. I, Metallurgiya, Moscow (1966),p. 455.</p><p>　　5

78、. V. A. Korablev, Yu. I. Ustinovshchikov, and I. G. Khatskelevich, "Embrittlement of chromium steels with formation of special carbides," Metalloved. Term. Obrab. Met., No. I,16 (1975).</p><p>　　6.

79、 A. P. Gulyaev, I. K. Kupalova, and V. A. Landa, "Method and results of phase analysis</p><p>　　of hlgh-speed steels," Zavod. Lab., No. 3, 298 (1965).</p><p>　　7. J. Heslop and N. Perc

80、h, Phil. Mag., ~, No. 34, 1128 (1958).</p><p>　　8. V. V. Rybin et al., "The mechanism of hardening of sorbite-hardening steel and the possibility of determining it theoretically and experimentally,"

81、; in: Metal Science [inRussian], No. 17, Sudostroenie, Leningrad (1973), p. 105.</p><p>　　9．L. K. Gordienko, Substrucutral Hardening of Metals and Alloys [in Russian], Nauka, Moscow(1973), p. 64.</p>

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83、2. Kim JY, Lee BW, Read DT, Kwon D (2005) Scr Mater 52:353</p><p>　　13. Kim JY, Lee JS, Lee KW, Kim KH, Kwon D (2006) Key Eng Mater 326–328:487</p><p>　　14. Kim JY, Lee JJ, Lee YH, Jang JI, Kwon

84、 D (2006) J Mater Res 21:2975</p><p>　　15. Kim JY, Kang SK, Lee JJ, Jang JI, Lee YH, Kwon D (2007) Acta Mater 55:3555</p><p>　　16. Dowling NE (1993) Mechanical behavior of materials. Prentice Ha

85、ll, Englewood Cliffs</p><p>　　17. Kim JY, Lee KW, Lee JS, Kwon D (2006) Surf Coat Technol 201:4278</p><p>　　18. DIN 17175-79 (1979) Seamless steel tubes for elevated temperatures</p><